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## **Meet the editor**

Professor Amin had completed his Master of Manufacturing Engineering and his Ph.D. in Technical Sciences from Volgograd Poly Technique Institute of the former USSR in 1979 and 1982. He started his teaching career at Bangladesh University of Engineering and Technology (BUET) in 1993. After 1997 he had moved to his current job at the International Islamic University Malaysia

(IIUM) where he has been working as a Head of Department for the past 7 years. He currently holds position as a senior Professor at the Manufacturing and Materials Engineering Department. His current research interests are in: heat assisted machining of difficult-to-cut metals and alloys including titanium alloys, hard machining, high speed machining of metals and ceramics and chatter control during metal cutting.

Contents

**Preface IX** 

**Part 1 Manufacturing Processes and** 

Chapter 1 **Numerical Modeling of the Additive** 

Zhiqiang Fan and Frank Liou

Chapter 2 **Formation of Alpha Case Mechanism** 

Chapter 3 **Genesis of Gas Containing** 

**Inherent Defects in Titanium Parts 1** 

**on Titanium Investment Cast Parts 29** 

**Defects in Cast Titanium Parts 42** 

**Part 2 Properties of Titanium Alloys Under High** 

Chapter 5 **Hot Plasticity of Alpha Beta Alloys 87**  Maciej Motyka, Krzysztof Kubiak, Jan Sieniawski and Waldemar Ziaja

Chapter 4 **Titanium Alloys at Extreme Pressure Conditions 67** 

Chapter 6 **Machinability of Titanium Alloys in Drilling 117** 

**Part 3 Surface Treatments of Titanium Alloys for** 

**Titanium Alloys – Nitriding 141**  Iryna Pohrelyuk and Viktor Fedirko

Chapter 7 **Chemico-Thermal Treatment of** 

Si-Young Sung, Beom-Suck Han and Young-Jig Kim

Vladimir Vykhodets, Tatiana Kurennykh and Nataliya Tarenkova

**Temperature and Ultra High Pressure Conditions 65** 

Nenad Velisavljevic, Simon MacLeod and Hyunchae Cynn

Safian Sharif, Erween Abd Rahim and Hiroyuki Sasahara

**Biomedical and Other Challenging Applications 139** 

**Manufacturing (AM) Processes of Titanium Alloy 3** 

### Contents

#### **Preface XI**


X Contents

#### Chapter 8 **Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 175**  Elzbieta Krasicka-Cydzik

#### Chapter 9 **Surface Modification Techniques for Biomedical Grade of Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 201**  S. Izman, Mohammed Rafiq Abdul-Kadir, Mahmood Anwar, E.M. Nazim, R. Rosliza, A. Shah and M.A. Hassan

### Preface

Though titanium and its alloys are relatively new engineering materials, they have found wide application in the aerospace, shipbuilding, automotive, sports, chemical and food processing industries due to their extreme lightness, high specific strength and good corrosion resistance at temperatures below 500oC. They are also considered suitable materials for biomedical application due to their biological passivity and biocompatibility. However, besides these positive properties titanium alloys have a number of adverse favorable properties which are related to their processing, machinability and long time use in open and corrosive environment. Chemical reactivity of titanium with other materials at elevated temperature is high, which necessitates the development of non conventional melting, refining and casting techniques, making this material very expensive. Numerous research is directed towards addressing these issues in order to ease their processing and further applications.

This nine diverse chapters of this book are distributed under three sections and address problems related to the processing and application of this precious metal and its alloys. The book chapters are contributed by researchers who devoted long periods of their research career working on titanium and its alloys looking for solutions to some of these specific problems.. From this perspective this book will serve as an excellent reference material for researchers whose works is in anyway related to titanium and its alloys - from processing to applications. The chapters are designed to address the issues that arise at the material development, processing and the application stages. For instance, the α case formation defects that arise at the investment casting stage and the optimization aspects of the additive manufacturing (AM) processes through numerical modelling and simulation has been addressed in two different chapters. Similarly, the metallurgical defects resulting from entrapped gases during casting processes are common to titanium parts. The morphology of formation of these defects in the production of Ti–6Al–4V alloy is presented in another chapter of the book with the objective of addressing effectively the problem at the casting stage. Titanium parts are sometimes designed to work as die components or as projectiles and as such are subjected to high pressures and temperatures. However, high pressure raises a number of scientific and engineering issues, mainly because under such pressures the relatively ductile phase may get transformed into a fairly brittle ω phase, which may significantly limit the use of titanium alloys in high pressure applications. One of the chapters of the book deals

#### X Preface

with these issues and indicates how the formation of ω phase may be avoided through adoption of proper processing techniques.

In the Chapter 'Hot Plasticity of alpha Beta alloys' the authors have shared their invaluable experimental results explaining different aspects of hot plasticity of twophase titanium alloys and have indicated techniques for developing appropriate microstructure yielding optimum plastic flow stresses under elevated temperatures. The phenomenon of super plasticity is also addressed in the same chapter. Furthermore, application of titanium alloys under corrosive environment and friction requires additional strengthening through effective surface treatment. One chapter of the book addresses different aspects of a common chemico-thermal method - nitriding to apply an effective coating on titanium parts. Specific issues related to the intricacies of the nitriding process suitable for application on titanium parts have been elaborated in the same chapter with excellent illustrations. Titanium alloys are common options in biomedical and dental applications and the ternary alloys Ti-6Al-7Nb are widely used for these purposes due to their unique mechanical and chemical properties, excellent corrosion resistance and biocompatibility. Development of nanotube anodic layers for medical applications on these materials are addressed in one of the chapters while the mechanism of electrochemical deposition of calcium phosphate on titanium substrate and the related process parameters and optimization techniques are presented in another. Titanium and its alloys are well known for their poor machinability properties. Useful experimental materials are presented in one chapter dealing with machining of titanium alloys, specifically drilling, a common machining operation.

We hope the research materials presented in the different chapters of this book will contribute to the ongoing research works on titanium and its alloys and help further improvement in the properties and application of titanium alloys.

#### **Acknowledgements**

The Editor would like congratulate the publishing team of INTECH for taking up this vital project and successfully steering it through its various reviewing, editing and publishing stages. Deep appreciation is extended to all the authors of the book chapters for their contribution in composing this valuable book. He would also like to acknowledge his deep appreciation to the Publishing Process Managers of the book for their sincere cooperation in rendering Editor's duties during the entire period of the editing and compilation process. Finally, he would like to express his gratefulness to the publisher for choosing him as the Editor of this book.

**Prof. Dr. A.K.M. Nurul Amin** 

Department of Manufacturing and Materials Engineering, Faculty of Engineering, International Islamic University of Malaysia, Malaysia

X Preface

adoption of proper processing techniques.

with these issues and indicates how the formation of ω phase may be avoided through

In the Chapter 'Hot Plasticity of alpha Beta alloys' the authors have shared their invaluable experimental results explaining different aspects of hot plasticity of twophase titanium alloys and have indicated techniques for developing appropriate microstructure yielding optimum plastic flow stresses under elevated temperatures. The phenomenon of super plasticity is also addressed in the same chapter. Furthermore, application of titanium alloys under corrosive environment and friction requires additional strengthening through effective surface treatment. One chapter of the book addresses different aspects of a common chemico-thermal method - nitriding to apply an effective coating on titanium parts. Specific issues related to the intricacies of the nitriding process suitable for application on titanium parts have been elaborated in the same chapter with excellent illustrations. Titanium alloys are common options in biomedical and dental applications and the ternary alloys Ti-6Al-7Nb are widely used for these purposes due to their unique mechanical and chemical properties, excellent corrosion resistance and biocompatibility. Development of nanotube anodic layers for medical applications on these materials are addressed in one of the chapters while the mechanism of electrochemical deposition of calcium phosphate on titanium substrate and the related process parameters and optimization techniques are presented in another. Titanium and its alloys are well known for their poor machinability properties. Useful experimental materials are presented in one chapter dealing with machining of

titanium alloys, specifically drilling, a common machining operation.

improvement in the properties and application of titanium alloys.

the publisher for choosing him as the Editor of this book.

**Acknowledgements** 

We hope the research materials presented in the different chapters of this book will contribute to the ongoing research works on titanium and its alloys and help further

The Editor would like congratulate the publishing team of INTECH for taking up this vital project and successfully steering it through its various reviewing, editing and publishing stages. Deep appreciation is extended to all the authors of the book chapters for their contribution in composing this valuable book. He would also like to acknowledge his deep appreciation to the Publishing Process Managers of the book for their sincere cooperation in rendering Editor's duties during the entire period of the editing and compilation process. Finally, he would like to express his gratefulness to

**Prof. Dr. A.K.M. Nurul Amin** 

Malaysia

Department of Manufacturing and Materials Engineering,

Faculty of Engineering, International Islamic University of Malaysia,

**Part 1** 

**Manufacturing Processes and** 

**Inherent Defects in Titanium Parts** 

## **Part 1**

**Manufacturing Processes and Inherent Defects in Titanium Parts** 

**1** 

*USA* 

**Numerical Modeling of** 

Zhiqiang Fan and Frank Liou

**the Additive Manufacturing (AM)** 

It is easy to understand why industry and, especially, aerospace engineers love titanium. Titanium parts weigh roughly half as much as steel parts, but its strength is far greater than the strength of many alloy steels giving it an extremely high strength-to-weight ratio. Most titanium alloys are poor thermal conductors, thus heat generated during cutting does not dissipate through the part and machine structure, but concentrates in the cutting area. The high temperature generated during the cutting process also causes a work hardening phenomenon that affects the surface integrity of titanium, and could lead to geometric inaccuracies in the part and severe reduction in its fatigue strength [Benes, 2007]. On the contrary, additive manufacturing (AM) is an effective way to process titanium alloys as AM is principally thermal based, the effectiveness of AM processes depends on the material's thermal properties and its absorption of laser energy rather than on its mechanical properties. Therefore, brittle and hard materials can be processed easily if their thermal properties (e.g., conductivity, heat of fusion, etc.) are favorable, such as titanium. Cost effectiveness is also an important consideration for using additive manufacturing for titanium processing. Parts or products cast and/or machined from titanium and its alloys are very expensive, due to the processing difficulties and complexities during machining and casting. AM processes however, have been found to be very cost effective because they can produce near-net shape parts from these high performance metals with little or no machining [Liou & Kinsella, 2009]. In the aerospace industry, titanium and its alloys are used for many large structural components. When traditional machining/cast routines are adopted, conversion costs for these heavy section components can be prohibitive due to long lead time and low-yield material utilization [Eylon & Froes, 1984]. AM processes have the potential to shorten lead time and increase material utilization in these applications. The following sections 1.1, 1.2 and 1.3 summarize the fundamental knowledge for the modeling

Additive manufacturing can be achieved by powder-based spray (e.g., thermal spray or cold spray), sintering (e.g., selective laser sintering), or fusion-based processes (or direct metal deposition) which use a laser beam, an electron beam, a plasma beam, or an electric arc as an

**1. Introduction** 

of additive manufacturing processes.

**1.1 Additive manufacturing** 

**Processes of Titanium Alloy** 

*Missouri University of Science and Technology* 

### **Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy**

Zhiqiang Fan and Frank Liou *Missouri University of Science and Technology USA* 

#### **1. Introduction**

It is easy to understand why industry and, especially, aerospace engineers love titanium. Titanium parts weigh roughly half as much as steel parts, but its strength is far greater than the strength of many alloy steels giving it an extremely high strength-to-weight ratio. Most titanium alloys are poor thermal conductors, thus heat generated during cutting does not dissipate through the part and machine structure, but concentrates in the cutting area. The high temperature generated during the cutting process also causes a work hardening phenomenon that affects the surface integrity of titanium, and could lead to geometric inaccuracies in the part and severe reduction in its fatigue strength [Benes, 2007]. On the contrary, additive manufacturing (AM) is an effective way to process titanium alloys as AM is principally thermal based, the effectiveness of AM processes depends on the material's thermal properties and its absorption of laser energy rather than on its mechanical properties. Therefore, brittle and hard materials can be processed easily if their thermal properties (e.g., conductivity, heat of fusion, etc.) are favorable, such as titanium. Cost effectiveness is also an important consideration for using additive manufacturing for titanium processing. Parts or products cast and/or machined from titanium and its alloys are very expensive, due to the processing difficulties and complexities during machining and casting. AM processes however, have been found to be very cost effective because they can produce near-net shape parts from these high performance metals with little or no machining [Liou & Kinsella, 2009]. In the aerospace industry, titanium and its alloys are used for many large structural components. When traditional machining/cast routines are adopted, conversion costs for these heavy section components can be prohibitive due to long lead time and low-yield material utilization [Eylon & Froes, 1984]. AM processes have the potential to shorten lead time and increase material utilization in these applications. The following sections 1.1, 1.2 and 1.3 summarize the fundamental knowledge for the modeling of additive manufacturing processes.

#### **1.1 Additive manufacturing**

Additive manufacturing can be achieved by powder-based spray (e.g., thermal spray or cold spray), sintering (e.g., selective laser sintering), or fusion-based processes (or direct metal deposition) which use a laser beam, an electron beam, a plasma beam, or an electric arc as an

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 5

Fan et al., 2007a; Gandin & Rappaz, 1994; Grujicic et al. 2001; Rappaz & Gandin, 1993; Zhu et al., 2004], front tracking [Juric & Tryggvason, 1996; Sullivan et al., 1987; Tryggvason et al., 2001], immersed boundary [Udaykumar et al., 1999, 2003] and level set [Gibou et al., 2003; Kim et al. 2000] methods. Due to the limits of current computing power, the above methods

only apply to small domains on a continuum scale from about 0.1 μm to 10 mm.

Fig. 1. Schematics of a coaxial laser metal deposition system with powder injection

In this study, the continuum model is adopted to develop the governing equations.

To treat the effects of transport phenomena at the process-scale (~ 1 m), a macroscopic model needs to be adopted, where a representative elementary volume (REV) is selected to include a representative and uniform sampling of the mushy region such that local scale solidification processes can be described by variables averaged over the REV [Voller et al., 2004]. Based on the REV concept, governing equations for the mass, momentum, energy and species conservation at the process scale are developed and solved. Two main approaches have been used for the derivation and solution of the macroscopic conservation equations. One approach is the two-phase model [Beckermann & Viskanta, 1988; Ganesan & Poirier, 1990; Ni & Beckermann, 1991], in which the two phases are treated as separate and separate volume-averaged conservation equations are derived for solid and liquid phases using a volume averaging technique. This approach gives the complete mathematical models for solidification developed today, which have the potential to build a strong linkage between physical phenomena occurring on macroscopic and microscopic scales [Ni & Incropera, 1995]. However, the numerical procedures of this model are fairly involved since two separate sets of conservation equations need to be solved and the interface between the two phases must be determined for each time step [Jaluria, 2006]. This places a great demand on computational capabilities. In addition, the lack of information about the microscopic configuration at the solid-liquid interface is still a serious obstacle in the implementation of this model for practical applications [Stefanescu, 2002]. An alternative approach to the development of macroscopic conservation equations is the continuum model [Bennon & Incropera, 1987; Hills et al., 1983; Prantil & Dawson, 1983; Prescott et al., 1991; Voller & Prakash, 1987; Voller et al., 1989]. This model uses the classical mixture theory [Muller, 1968] to develop a single set of mass, momentum, energy and species conservation equations, which concurrently apply to the solid, liquid and mushy regions. The numerical procedures for this model are much simpler since the same equations are employed over the entire computational domain, thereby facilitating use of standard, single-phase CFD procedures.

energy source and either metallic powder or wire as feedstock [Kobryn et al., 2006]. For the aerospace industry which is the biggest titanium market in the U.S. [Yu & Imam, 2007], fusion-based AM processes are more advantageous since they can produce 100% dense functional metal parts. This chapter will focus on fusion-based AM processes with application to titanium.

Numerical modeling and simulation is a very useful tool for assessing the impact of process parameters and predicting optimized conditions in AM processes. AM processes involve many process parameters, including total power and power intensity distribution of the energy source, travel speed, translation path, material feed rate and shielding gas pressure. These process parameters not only vary from part to part, but also frequently vary locally within a single part to attain the desired deposit shape [Kobryn et al., 2006]. Physical phenomena associated with AM processes are complex, including melting/solidification and vaporization phase changes, surface tension-dominated free-surface flow, heat and mass transfer, and moving heat source. The variable process parameters together with the interacting physical phenomena involved in AM complicate the development of processproperty relationships and appropriate process control. Thus, an effective numerical modeling of the processing is very useful for assessing the impact of process parameters and predicting optimized conditions.

Currently process-scale modeling mainly addresses transport phenomena such as heat transfer and fluid dynamics, which are closely related to the mechanical properties of the final structure. For example, the buoyancy-driven flow due to temperature and species gradients in the melt pool strongly influences the microstructure and thus the mechanical properties of the final products. The surface tension-driven free-surface flow determines the shape and smoothness of the clad. In this chapter, numerical modeling of transport phenomena in fusion-based AM processes will be presented, using the laser metal deposition process as an example. Coaxial laser deposition systems with blown powder as shown in Fig. 1 are considered for simulations and experiments. The material studied is Ti-6Al-4V for both the substrate and powder. As the main challenges in modeling of fusionbased AM processes are related to melting/solidification phase change and free-surface flow in the melt pool, modeling approaches for these physical phenomena will be introduced in Sections 1.2 and 1.3.

#### **1.2 Modeling of melting/solidification phase change**

Fusion-based AM processes involve a melting/solidification phase change. Numerical modeling of the solidification of metal alloys is very challenging because a general solidification of metal alloys involves a so-called "mushy region" over which both solid and liquid coexist and the transport phenomena occur across a wide range of time and length scales [Voller, 2006]. A rapidly developing approach that tries to resolve the smallest scales of the solid-liquid interface can be thought of as direct microstructure simulation. In order to simulate the microstructure development directly, the evolution of the interface between different phases or different microstructure constituents has to be calculated, coupled with the physical fields such as temperature and concentration [Pavlyk & Dilthey, 2004]. To this approach belong phase-field [Beckermann et al., 1999; Boettinger et al., 2002; Caginalp, 1989; Karma & Rappel, 1996,1998; Kobayashi,1993; Provatas et al., 1998; Steinbach et al., 1996; Warren & Boettinger, 1995; Wheeler et al., 1992], cellular-automaton [Boettinger et al., 2000;

energy source and either metallic powder or wire as feedstock [Kobryn et al., 2006]. For the aerospace industry which is the biggest titanium market in the U.S. [Yu & Imam, 2007], fusion-based AM processes are more advantageous since they can produce 100% dense functional metal parts. This chapter will focus on fusion-based AM processes with

Numerical modeling and simulation is a very useful tool for assessing the impact of process parameters and predicting optimized conditions in AM processes. AM processes involve many process parameters, including total power and power intensity distribution of the energy source, travel speed, translation path, material feed rate and shielding gas pressure. These process parameters not only vary from part to part, but also frequently vary locally within a single part to attain the desired deposit shape [Kobryn et al., 2006]. Physical phenomena associated with AM processes are complex, including melting/solidification and vaporization phase changes, surface tension-dominated free-surface flow, heat and mass transfer, and moving heat source. The variable process parameters together with the interacting physical phenomena involved in AM complicate the development of processproperty relationships and appropriate process control. Thus, an effective numerical modeling of the processing is very useful for assessing the impact of process parameters and

Currently process-scale modeling mainly addresses transport phenomena such as heat transfer and fluid dynamics, which are closely related to the mechanical properties of the final structure. For example, the buoyancy-driven flow due to temperature and species gradients in the melt pool strongly influences the microstructure and thus the mechanical properties of the final products. The surface tension-driven free-surface flow determines the shape and smoothness of the clad. In this chapter, numerical modeling of transport phenomena in fusion-based AM processes will be presented, using the laser metal deposition process as an example. Coaxial laser deposition systems with blown powder as shown in Fig. 1 are considered for simulations and experiments. The material studied is Ti-6Al-4V for both the substrate and powder. As the main challenges in modeling of fusionbased AM processes are related to melting/solidification phase change and free-surface flow in the melt pool, modeling approaches for these physical phenomena will be

Fusion-based AM processes involve a melting/solidification phase change. Numerical modeling of the solidification of metal alloys is very challenging because a general solidification of metal alloys involves a so-called "mushy region" over which both solid and liquid coexist and the transport phenomena occur across a wide range of time and length scales [Voller, 2006]. A rapidly developing approach that tries to resolve the smallest scales of the solid-liquid interface can be thought of as direct microstructure simulation. In order to simulate the microstructure development directly, the evolution of the interface between different phases or different microstructure constituents has to be calculated, coupled with the physical fields such as temperature and concentration [Pavlyk & Dilthey, 2004]. To this approach belong phase-field [Beckermann et al., 1999; Boettinger et al., 2002; Caginalp, 1989; Karma & Rappel, 1996,1998; Kobayashi,1993; Provatas et al., 1998; Steinbach et al., 1996; Warren & Boettinger, 1995; Wheeler et al., 1992], cellular-automaton [Boettinger et al., 2000;

application to titanium.

predicting optimized conditions.

introduced in Sections 1.2 and 1.3.

**1.2 Modeling of melting/solidification phase change** 

Fan et al., 2007a; Gandin & Rappaz, 1994; Grujicic et al. 2001; Rappaz & Gandin, 1993; Zhu et al., 2004], front tracking [Juric & Tryggvason, 1996; Sullivan et al., 1987; Tryggvason et al., 2001], immersed boundary [Udaykumar et al., 1999, 2003] and level set [Gibou et al., 2003; Kim et al. 2000] methods. Due to the limits of current computing power, the above methods only apply to small domains on a continuum scale from about 0.1 μm to 10 mm.

Fig. 1. Schematics of a coaxial laser metal deposition system with powder injection

To treat the effects of transport phenomena at the process-scale (~ 1 m), a macroscopic model needs to be adopted, where a representative elementary volume (REV) is selected to include a representative and uniform sampling of the mushy region such that local scale solidification processes can be described by variables averaged over the REV [Voller et al., 2004]. Based on the REV concept, governing equations for the mass, momentum, energy and species conservation at the process scale are developed and solved. Two main approaches have been used for the derivation and solution of the macroscopic conservation equations. One approach is the two-phase model [Beckermann & Viskanta, 1988; Ganesan & Poirier, 1990; Ni & Beckermann, 1991], in which the two phases are treated as separate and separate volume-averaged conservation equations are derived for solid and liquid phases using a volume averaging technique. This approach gives the complete mathematical models for solidification developed today, which have the potential to build a strong linkage between physical phenomena occurring on macroscopic and microscopic scales [Ni & Incropera, 1995]. However, the numerical procedures of this model are fairly involved since two separate sets of conservation equations need to be solved and the interface between the two phases must be determined for each time step [Jaluria, 2006]. This places a great demand on computational capabilities. In addition, the lack of information about the microscopic configuration at the solid-liquid interface is still a serious obstacle in the implementation of this model for practical applications [Stefanescu, 2002]. An alternative approach to the development of macroscopic conservation equations is the continuum model [Bennon & Incropera, 1987; Hills et al., 1983; Prantil & Dawson, 1983; Prescott et al., 1991; Voller & Prakash, 1987; Voller et al., 1989]. This model uses the classical mixture theory [Muller, 1968] to develop a single set of mass, momentum, energy and species conservation equations, which concurrently apply to the solid, liquid and mushy regions. The numerical procedures for this model are much simpler since the same equations are employed over the entire computational domain, thereby facilitating use of standard, single-phase CFD procedures. In this study, the continuum model is adopted to develop the governing equations.

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 7

In this study the calculation domain for a laser deposition system includes the substrate, melt pool, remelted zone, deposited layer and part of the gas region, as shown in Fig.2. The continuum model Bennon & Incropera, 1987; Prescott et al., 1991 is adopted to derive the governing equations for melting and solidification with the mushy zone. Some important terms for the melt pool have been added in the momentum equations, including the buoyancy force term and surface tension force term, while some minor terms in the original derivation in [Prescott et al., 1991 have been neglected. The molten metal is assumed to be Newtonian fluid, and the melt pool is assumed to be an incompressible, laminar flow. The laminar flow assumption can be relaxed if turbulence is considered by an appropriate turbulence model, such as a low-Reynolds-number *k*-ε model [Jones & Launder, 1973]. The solid and liquid phases in the mushy zone are assumed to be in local thermal equilibrium.

> Substrate Remelted Zone Deposited Layer

Fig. 2. Schematic diagram of the calculation domain for laser metal deposition process

( )0 *<sup>t</sup>*

<sup>0</sup> () ( ) ( ) ( ) [1 ( )] *<sup>l</sup>*

<sup>0</sup> () ( ) ( ) ( ) [1 ( )] *<sup>l</sup>*

*l y l <sup>p</sup> vv v v v T T*

*l x l <sup>p</sup> uu u u u T T*

 

> 

 

> 

*l s x Sx*

*l s y Sy*

**V g F** (3)

**V g F** (2)

**V** (1)

 

 

For the system of interest, the conservation equations are summarized as follows:

*t x K*

*t y K*

 

  Melt pool

Laser Beam

Powder

Shielding Gas

**2. Mathematical model 2.1 Governing equations** 

Mass conservation:

Momentum conservation:

Energy conservation:

#### **1.3 Modeling of free-surface flow**

In fusion-based AM processes, the melt pool created by the energy source on the substrate is usually modelled as a free-surface flow, in which the pressure of the lighter fluid is not dependent on space, and viscous stresses in the lighter fluid is negligible. The techniques to find the shape of the free surface can be classified into two major groups: Lagrangian (or moving grid) methods and Eulerian (or fixed grid) methods. In Lagrangian methods [Hansbo, 2000; Idelsohn et al., 2001; Ramaswany& Kahawara, 1987; Takizawa et al., 1992], every point of the liquid domain is moved with the liquid velocity. A continuous remeshing of the domain or part of it is required at each time step so as to follow the interface movement. A special procedure is needed to enforce volume conservation in the moving cells. All of this can lead to complex algorithms. They are mainly used if the deformation of the interface is small, for example, in fluid-structure interactions or small amplitude waves [Caboussat, 2005]. In Eulerian methods, the interface is moving within a fixed grid, and no re-meshing is needed. The interface is determined from a field variable, for example, a volume fraction [DeBar, 1974; Hirt & Nichols, 1981; Noh & Woodward, 1976], a level-set [ Sethian, 1996, 1999] or a phase-field [Boettinger et al., 2002; Jacqmin, 1999]. While Lagrangian techniques are superior for small deformations of the interfaces, Eulerian techniques are usually preferred for highly distorted, complex interfaces, which is the case for fusion-based additive manufacturing processes. For example, in AM processes with metallic powder as feedstock, powder injection causes intermittent mergers and breakups at the interface of the melt pool, which needs a robust Eulerian technique to handle.

Among the Eulerian methods, VOF (for Volume-Of-Fluid) [Hirt & Nichols, 1981] is probably the most widely used. It has been adopted by many in-house codes and built into commercial codes (SOLA-VOF [Nichols et al, 1980], NASA-VOF2D [Torrey et al 1985], NASA-VOF3D [Torrey et al 1987], RIPPLE [Kothe & Mjolsness 1991], and FLOW3D [Hirt & Nichols 1988], ANSYS Fluent, to name a few). In this method a scalar indicator function, F, is defined on the grid to indicate the liquid-volume fraction in each computational cell. Volume fraction values between zero and unity indicate the presence of the interface. The VOF method consists of an interface reconstruction algorithm and a volume fraction advection scheme. The features of these two steps are used to distinguish different VOF versions. For modeling of AM processes, an advantage of VOF is that it can be readily integrated with the techniques for simulation of the melting /solidification phase change. VOF methods have gone through a continuous process of development and improvement. Reviews of the historical development of VOF can be found in [Benson, 2002; Rider & Kothe, 1998; Rudman, 1997; Tang et al., 2004]. In earlier versions of VOF [Chorin, 1980; Debar, 1974; Hirt & Nichols, 1981; Noh & Woodward, 1976], reconstruction algorithms are based on a piecewise-constant or "stair-stepped" representation of the interface and advection schemes are at best first-order accurate. These first-order VOF methods are numerically unstable in the absence of surface tension, leading to the deterioration of the interface in the form of flotsam and jetsam [Scardovelli & Zaleski, 1999]. The current generation of VOF methods approximate the interface as a plane within a computational cell, and are commonly referred to as piecewise linear interface construction (PLIC) methods [Gueyffier et al., 1999; Rider & Kothe, 1998; Youngs, 1982, 1984]. PLIC-VOF is more accurate and avoids the numerical instability [Scardovelli & Zaleski, 1999].

#### **2. Mathematical model**

6 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

In fusion-based AM processes, the melt pool created by the energy source on the substrate is usually modelled as a free-surface flow, in which the pressure of the lighter fluid is not dependent on space, and viscous stresses in the lighter fluid is negligible. The techniques to find the shape of the free surface can be classified into two major groups: Lagrangian (or moving grid) methods and Eulerian (or fixed grid) methods. In Lagrangian methods [Hansbo, 2000; Idelsohn et al., 2001; Ramaswany& Kahawara, 1987; Takizawa et al., 1992], every point of the liquid domain is moved with the liquid velocity. A continuous remeshing of the domain or part of it is required at each time step so as to follow the interface movement. A special procedure is needed to enforce volume conservation in the moving cells. All of this can lead to complex algorithms. They are mainly used if the deformation of the interface is small, for example, in fluid-structure interactions or small amplitude waves [Caboussat, 2005]. In Eulerian methods, the interface is moving within a fixed grid, and no re-meshing is needed. The interface is determined from a field variable, for example, a volume fraction [DeBar, 1974; Hirt & Nichols, 1981; Noh & Woodward, 1976], a level-set [ Sethian, 1996, 1999] or a phase-field [Boettinger et al., 2002; Jacqmin, 1999]. While Lagrangian techniques are superior for small deformations of the interfaces, Eulerian techniques are usually preferred for highly distorted, complex interfaces, which is the case for fusion-based additive manufacturing processes. For example, in AM processes with metallic powder as feedstock, powder injection causes intermittent mergers and breakups at

the interface of the melt pool, which needs a robust Eulerian technique to handle.

Among the Eulerian methods, VOF (for Volume-Of-Fluid) [Hirt & Nichols, 1981] is probably the most widely used. It has been adopted by many in-house codes and built into commercial codes (SOLA-VOF [Nichols et al, 1980], NASA-VOF2D [Torrey et al 1985], NASA-VOF3D [Torrey et al 1987], RIPPLE [Kothe & Mjolsness 1991], and FLOW3D [Hirt & Nichols 1988], ANSYS Fluent, to name a few). In this method a scalar indicator function, F, is defined on the grid to indicate the liquid-volume fraction in each computational cell. Volume fraction values between zero and unity indicate the presence of the interface. The VOF method consists of an interface reconstruction algorithm and a volume fraction advection scheme. The features of these two steps are used to distinguish different VOF versions. For modeling of AM processes, an advantage of VOF is that it can be readily integrated with the techniques for simulation of the melting /solidification phase change. VOF methods have gone through a continuous process of development and improvement. Reviews of the historical development of VOF can be found in [Benson, 2002; Rider & Kothe, 1998; Rudman, 1997; Tang et al., 2004]. In earlier versions of VOF [Chorin, 1980; Debar, 1974; Hirt & Nichols, 1981; Noh & Woodward, 1976], reconstruction algorithms are based on a piecewise-constant or "stair-stepped" representation of the interface and advection schemes are at best first-order accurate. These first-order VOF methods are numerically unstable in the absence of surface tension, leading to the deterioration of the interface in the form of flotsam and jetsam [Scardovelli & Zaleski, 1999]. The current generation of VOF methods approximate the interface as a plane within a computational cell, and are commonly referred to as piecewise linear interface construction (PLIC) methods [Gueyffier et al., 1999; Rider & Kothe, 1998; Youngs, 1982, 1984]. PLIC-VOF is more accurate and avoids the numerical

**1.3 Modeling of free-surface flow** 

instability [Scardovelli & Zaleski, 1999].

#### **2.1 Governing equations**

In this study the calculation domain for a laser deposition system includes the substrate, melt pool, remelted zone, deposited layer and part of the gas region, as shown in Fig.2. The continuum model Bennon & Incropera, 1987; Prescott et al., 1991 is adopted to derive the governing equations for melting and solidification with the mushy zone. Some important terms for the melt pool have been added in the momentum equations, including the buoyancy force term and surface tension force term, while some minor terms in the original derivation in [Prescott et al., 1991 have been neglected. The molten metal is assumed to be Newtonian fluid, and the melt pool is assumed to be an incompressible, laminar flow. The laminar flow assumption can be relaxed if turbulence is considered by an appropriate turbulence model, such as a low-Reynolds-number *k*-ε model [Jones & Launder, 1973]. The solid and liquid phases in the mushy zone are assumed to be in local thermal equilibrium.

Fig. 2. Schematic diagram of the calculation domain for laser metal deposition process For the system of interest, the conservation equations are summarized as follows: Mass conservation:

$$\frac{\partial \rho}{\partial t} + \nabla \cdot (\rho \mathbf{V}) = 0 \tag{1}$$

Momentum conservation:

$$\frac{\partial}{\partial t}(\rho u) + \nabla \cdot (\rho \mathbf{V} u) = \nabla \cdot (\mu\_l \frac{\rho}{\rho\_l} \nabla u) - \frac{\partial p}{\partial x} - \frac{\mu\_l}{K\_x} \frac{\rho}{\rho\_l} (u - u\_s) + \rho \mathbf{g}\_x [1 - \alpha (T - T\_0)] + \mathbf{F}\_{\text{Sx}} \tag{2}$$

$$\frac{\partial}{\partial t}(\rho v) + \nabla \cdot (\rho \mathbf{V} v) = \nabla \cdot (\mu\_l \frac{\rho}{\rho\_l} \nabla v) - \frac{\partial p}{\partial y} - \frac{\mu\_l}{K\_y} \frac{\rho}{\rho\_l} (v - v\_s) + \rho \mathbf{g}\_y [1 - \alpha (T - T\_0)] + \mathbf{F}\_{\mathbb{S}y} \tag{3}$$

Energy conservation:

$$\frac{\partial}{\partial t}(\rho h) + \nabla \cdot (\rho \mathbf{V} h) = \nabla \cdot (k \nabla T) - \nabla \cdot [\rho (h\_l - h)(\mathbf{V} - \mathbf{V}\_s)] + \mathbf{S} \tag{4}$$

In equations (1)-(4), the subscripts *s* and *l* stand for solid and liquid phase, respectively. *t*, *μ*, and *T* are time, dynamics viscosity and temperature, respectively. *u* and *v* are x-direction and y-direction velocity components. The continuum density ρ, vector velocity V, enthalpy *h,* and thermal conductivity *k* are defined as follows:

$$
\rho = \mathbb{g}\_s \rho\_s + \mathbb{g}\_l \rho\_l \tag{5}
$$

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 9

1984; Ganesan et al., 1992; Poirier, 1987; West, 1985. Here the mushy zone is considered as rigid (i.e. a porous media). If the mushy zone is modeled as a slurry region, these two terms can be treated as in [Ni & Incropera, 1995]. In Eqs. (2) and (3), the fourth terms on the righthand side are the buoyancy force components due to temperature gradients. Here Boussinesq approximation is applied. is the thermal expansion coefficient. The fifth terms on the right-hand side of Eqs. (2) and (3) are surface tension force components, which will be

> ˆ *S S* **F n**

caused by spatial variations in the surface tension coefficient along the interface due to temperature and/or species gradients. It causes the fluid flow from regions of lower to

The conventional approach when dealing with surface tension is to use finite difference schemes to apply a pressure jump at a free-surface discontinuity. More recently, a general practice is to model surface tension as a volume force using a continuum model, either the Continuum Surface Force (CSF) model [Brackbill et al., 1992] or the Continuum Surface Stress (CSS) model [Lafaurie et al., 1994]. The volume force acts everywhere within a finite transition region between the two phases. In this study, the CSF model is adopted, which has been shown to make more accurate use of the free-surface VOF data [Brackbill et al.,

A well-known problem with VOF (and other Eularian methods) modeling of surface tension is so-called "parasitic currents" or "spurious currents", which is a flow induced solely by the discretization and by a lack of convergence with mesh refinement. Under some circumstances, this artificial flow can be strong enough to dominate the solution, and the resulting strong vortices at the interface may lead to catastrophic instability of the interface and may even break-up [Fuster et al., 2009; Gerlach et al. 2006]. Two measures can be taken to relieve or even resolve this problem. One measure is to use a force-balance flow algorithm in which the CSF model is applied in a way that is consistent with the calculation of the pressure gradient field. Thus, imbalance between discrete surface tension and pressuregradient terms can be avoided. Within a VOF framework, such force-balance flow algorithms can be found in [Francois et al., 2006; Y.Renardy & M. Renardy, 2002; Shirani et al., 2005]. In this study, the algorithm in [Shirani et al., 2005] is followed. The other measure is to get an accurate calculation of surface tension by accurately calculating interface normals and curvatures from volume fractions. For this purpose, many methods have been developed, such as those in [Cummins et al, 2005; Francois et al., 2006; López & Hernández, 2010; Meier et al., 2002; Pilliod Jr. & Puckett, 2004; Y.Renardy & M. Renardy, 2002]. The method we use here is the height function (HF) technique, which has been shown to be second-order accurate, and superior to those based on kernel derivatives of volume fractions

 

is surface tension coefficient, κ the curvature of the interface, **n**ˆ the unit normal to

(13)

represents the Marangoni effect

**n**ˆ is the normal

described in Section 2.2 below. The term S in Eq. (4) is the heat source.

the local surface, and *<sup>S</sup>* the surface gradient operator. The term

component of the surface tension force. The term *<sup>S</sup>*

**2.2 Surface tension** 

where 

1992].

The surface tension force, *FS*, is given by:

higher surface tension coefficient.

$$\mathbf{V} = f\_s \mathbf{V}\_s + f\_l \mathbf{V}\_l \tag{6}$$

$$\mathbf{h} = f\_s \mathbf{h}\_s + f\_l \mathbf{h}\_l \tag{7}$$

$$k = \mathcal{g}\_s k\_s + \mathcal{g}\_l k\_l \tag{8}$$

Here the subscripts *s* and *l* stand for solid and liquid phase, respectively. *fs* and *fl* refer to mass fractions of solid and liquid phases, and *gs* and *gl* are volume fractions of solid and liquid phases. To calculate these four quantities, a general practice is that *gl* (or *gs*) is calculated first and then the other three quantities are obtained according to the following relationships:

$$f\_l = \frac{\mathcal{g}\_l \mathcal{P}\_l}{\rho} \quad f\_s = \frac{\mathcal{g}\_s \mathcal{P}\_s}{\rho} \quad \mathcal{g}\_s + \mathcal{g}\_l = 1 \quad f\_s + f\_l = 1 \tag{9}$$

The volume fraction of liquid *gl* can be found using different models, such as the level rule, the Scheil model [Scheil, 1942], or the Clyne and Kurz model [Clyne & Kurz, 1981]. For the target material Ti-6Al-4V, it is assumed that *gl* is only dependent on temperature. The *gl* (T) function is given by Swaminathan & Voller, 1992:

$$\mathbf{g}\_{l} = \begin{cases} 0 & \text{if } T < T\_s \\ \frac{T - T\_s}{T\_l - T\_s} & \text{if } T\_s \le T \le T\_l \\ & \mathbf{1} & \text{if } T > T\_l \end{cases} \tag{10}$$

The phase enthalpy for the solid and the liquid can be expressed as:

$$h\_s = \int\_0^T c\_s(T)dT\tag{11}$$

$$h\_l = \int\_0^{T\_s} c\_s(T)dT + \int\_{T\_s}^T c\_l(T)dT + L\_m \tag{12}$$

where *L*m is the latent heat of melting. *cs* and *cl* are specific heat of solid and liquid phases.

In Eqs. (2) and (3), the third terms on the right-hand side are the drag interaction terms, and *Kx* and *Ky* are the permeability of the two-phase mushy zone in x- and y- directions, which can be calculated from various models Bhat et al., 1995; Carman, 1937; Drummond & Tahir, 1984; Ganesan et al., 1992; Poirier, 1987; West, 1985. Here the mushy zone is considered as rigid (i.e. a porous media). If the mushy zone is modeled as a slurry region, these two terms can be treated as in [Ni & Incropera, 1995]. In Eqs. (2) and (3), the fourth terms on the righthand side are the buoyancy force components due to temperature gradients. Here Boussinesq approximation is applied. is the thermal expansion coefficient. The fifth terms on the right-hand side of Eqs. (2) and (3) are surface tension force components, which will be described in Section 2.2 below. The term S in Eq. (4) is the heat source.

#### **2.2 Surface tension**

8 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

( ) ( ) ( ) [ ( )( )] *l s h h kT h h S*

In equations (1)-(4), the subscripts *s* and *l* stand for solid and liquid phase, respectively. *t*, *μ*, and *T* are time, dynamics viscosity and temperature, respectively. *u* and *v* are x-direction and y-direction velocity components. The continuum density ρ, vector velocity V, enthalpy

*ss ll*

 

 *g g* 

Here the subscripts *s* and *l* stand for solid and liquid phase, respectively. *fs* and *fl* refer to mass fractions of solid and liquid phases, and *gs* and *gl* are volume fractions of solid and liquid phases. To calculate these four quantities, a general practice is that *gl* (or *gs*) is calculated first and then the other three quantities are obtained according to the following

The volume fraction of liquid *gl* can be found using different models, such as the level rule, the Scheil model [Scheil, 1942], or the Clyne and Kurz model [Clyne & Kurz, 1981]. For the target material Ti-6Al-4V, it is assumed that *gl* is only dependent on temperature. The *gl* (T)

0 if

1 if

*l s*

*T T g T TT T T*

if

<sup>0</sup> ( ) *T*

<sup>0</sup> () () *<sup>s</sup> s*

where *L*m is the latent heat of melting. *cs* and *cl* are specific heat of solid and liquid phases. In Eqs. (2) and (3), the third terms on the right-hand side are the drag interaction terms, and *Kx* and *Ky* are the permeability of the two-phase mushy zone in x- and y- directions, which can be calculated from various models Bhat et al., 1995; Carman, 1937; Drummond & Tahir,

*ls l m <sup>T</sup>*

*T T*

*<sup>s</sup> l sl*

*s*

*T T*

*T T*

*l*

**V V V** (4)

*ss ll* **VVV** *f f* (6)

*ss ll h fh fh* (7)

*ss ll k gk gk* (8)

<sup>1</sup> *s l g g* <sup>1</sup> *s l f f* (9)

*s s h c T dT* (11)

*h c T dT c T dT L* (12)

(10)

(5)

*t* 

relationships:

 

*l l*

 *s s <sup>s</sup> <sup>g</sup> <sup>f</sup>* 

The phase enthalpy for the solid and the liquid can be expressed as:

*l <sup>g</sup> <sup>f</sup>* 

function is given by Swaminathan & Voller, 1992:

*h,* and thermal conductivity *k* are defined as follows:

The surface tension force, *FS*, is given by:

$$\mathbf{F}\_{\rm S} = \boldsymbol{\chi}\boldsymbol{\kappa}\hat{\mathbf{n}} + \nabla\_{\rm S}\boldsymbol{\gamma} \tag{13}$$

where is surface tension coefficient, κ the curvature of the interface, **n**ˆ the unit normal to the local surface, and *<sup>S</sup>* the surface gradient operator. The term**n**ˆ is the normal component of the surface tension force. The term *<sup>S</sup>* represents the Marangoni effect caused by spatial variations in the surface tension coefficient along the interface due to temperature and/or species gradients. It causes the fluid flow from regions of lower to higher surface tension coefficient.

The conventional approach when dealing with surface tension is to use finite difference schemes to apply a pressure jump at a free-surface discontinuity. More recently, a general practice is to model surface tension as a volume force using a continuum model, either the Continuum Surface Force (CSF) model [Brackbill et al., 1992] or the Continuum Surface Stress (CSS) model [Lafaurie et al., 1994]. The volume force acts everywhere within a finite transition region between the two phases. In this study, the CSF model is adopted, which has been shown to make more accurate use of the free-surface VOF data [Brackbill et al., 1992].

A well-known problem with VOF (and other Eularian methods) modeling of surface tension is so-called "parasitic currents" or "spurious currents", which is a flow induced solely by the discretization and by a lack of convergence with mesh refinement. Under some circumstances, this artificial flow can be strong enough to dominate the solution, and the resulting strong vortices at the interface may lead to catastrophic instability of the interface and may even break-up [Fuster et al., 2009; Gerlach et al. 2006]. Two measures can be taken to relieve or even resolve this problem. One measure is to use a force-balance flow algorithm in which the CSF model is applied in a way that is consistent with the calculation of the pressure gradient field. Thus, imbalance between discrete surface tension and pressuregradient terms can be avoided. Within a VOF framework, such force-balance flow algorithms can be found in [Francois et al., 2006; Y.Renardy & M. Renardy, 2002; Shirani et al., 2005]. In this study, the algorithm in [Shirani et al., 2005] is followed. The other measure is to get an accurate calculation of surface tension by accurately calculating interface normals and curvatures from volume fractions. For this purpose, many methods have been developed, such as those in [Cummins et al, 2005; Francois et al., 2006; López & Hernández, 2010; Meier et al., 2002; Pilliod Jr. & Puckett, 2004; Y.Renardy & M. Renardy, 2002]. The method we use here is the height function (HF) technique, which has been shown to be second-order accurate, and superior to those based on kernel derivatives of volume fractions or RDF distributions [Cummins et al, 2005; Francois et al., 2006; Liovic et al., 2010]. Specifically, we adopt the HF technique in [López & Hernández, 2010] that has many improvements over earlier versions (such as that in [Torrey et al., 1985]) of HF, including using an error correction procedure to minimize estimation error. Within the HF framework, suppose the absolute value of the y-direction component of the interface normal vector is larger than the x-direction component, interface curvature (in 2D) is given by

$$\kappa = \frac{H\_{\text{xx}}}{\left(1 + H\_{\text{x}}^{\text{-}2}\right)^{3/2}} \tag{14}$$

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 11

convective coefficient, ε emissivity, σ the Stefan-Boltzmann constant, and **n** the normal vector at the local interface. *me* can be evaluated according to the "overall evaporation model" in [Choi et al., 1987], and Patten can be calculated according to Frenk et al., 1997 with

( )0 *<sup>c</sup>*

(18)

**V** 0 (19)

*<sup>T</sup> k hT T*

**<sup>n</sup>**

Note that the radiation heat loss at these surfaces is neglected due to the fact that the

Finite difference and finite volume methods are used for spatial discretization of the governing equations. Staggered grids are employed where the temperatures, pressures and VOF function are located at the cell center and the velocities at the walls. In the numerical implementation, material properties play an important role. The material properties are generally dependent on temperature, concentration, and pressure. For fusion-based additive manufacturing processes, the material experiences a large variation from room temperature to above the melting temperature. For Ti-6Al-4V, many material properties experience large variations over this wide temperature range, as shown in Table 1. For example, the value of specific heat varies from 546 JK-1kg-1 at room temperature to 831 JK-1kg-1 at liquidus temperature. Thermal conductivity varies from 7 to 33.4 Wm-1K-1 over the same temperature range. Thus, the temperature dependence of the properties dominates, which necessitates a coupling of the momentum equations with the energy equation and

The variable properties have two effects on the numerical solution procedure [Ferziger & Peric, 2002]. First, although an incompressible flow assumption is made, the thermophysical properties need to be kept inside the differential operators. Thus, solution methods for incompressible flow can be used. Second, the momentum and energy conservation equations have to be solved in a coupled way. In this study, the coupling between

1. Eqs. (1) - (3) and the related boundary conditions are solved iteratively using a two-step projection method [Chorin, 1968] to obtain velocities and pressures. Thermo-physical properties used in this step are computed from the old temperature field. At each time step, the discretized momentum equations calculate new velocities in terms of an estimated pressure field. Then the pressure field is iteratively adjusted and velocity changes induced by each pressure correction are added to the previous velocities. This iterative process is repeated until the continuity equation is satisfied under an imposed tolerance by the newly computed velocities. This imposes a requirement for solving a linear system of equations. The preconditioned Bi-CGSTAB (for Bi-Conjugate Gradient Stabilized) method [Barrett et al., 1994] is used to solve the linear system of equations.

momentum and energy equations is achieved by the following iterative scheme:

On the bottom surface and side surfaces, boundary conditions are given by

temperature differences at these surfaces are not large.

gives rise to strong nonlinearity in the conservation equations.

a minor modification.

**2.5 Numerical Implementation** 

where *H* is the height function, *Hx* and *Hxx* are first-order and second-order derivatives of *H*, respectively. *Hx* and *Hxx* are obtained by using a finite difference formula. Interface normals are calculated based on the Least-Squares Fit method from [Aulisa et al., 2007].

#### **2.3 Tracking of the free surface**

The free surface of the melt pool is tracked using the PLIC-VOF [Gueyffier et al., 1999; Scardovelli & Zaleski, 2000, 2003]. The Volume of Fluid function, F, satisfies the following conservation equation:

$$\frac{\partial F}{\partial t} + (\mathbf{V} \cdot \nabla)F = 0\tag{15}$$

The PLIC-VOF method consists of two steps: interface reconstruction and interface advection. In 2D calculation, a reconstructed planar surface becomes a straight line which satisfies the following equation:

$$m\_x \mathbf{x} + n\_y \mathbf{y} = d\tag{16}$$

where *nx* and *ny* are x and y components of the interface normal vector. *d* is a parameter related to the distance between the line and the coordinate origin of the reference cell. In the interface reconstruction step, *nx* and *ny* of each cell are calculated based on volume fraction data, using the Least-Squares Fit method from [Aulisa et al., 2007]. Then the parameter *d* is determined to match the given volume fraction. Finally given the velocity field, the reconstructed interface is advected according to the combined Eulerian-Lagrangian scheme in [Aulisa et al., 2007].

#### **2.4 Boundary conditions**

Energy balance at the free surface satisfies the following equation:

$$k\frac{\partial T}{\partial \mathbf{n}} = \frac{\eta \left(P\_{\text{laser}} - P\_{\text{atten}}\right)}{\pi R^2} - h\_c \left(T - T\_w\right) - \varepsilon \sigma \left(T^4 - T\_w^4\right) - \dot{m}\_c L\_v \tag{17}$$

where terms on the right-hand side are laser irradiation, convective heat loss, radiation heat loss and evaporation heat loss, respectively. *P*laser is the power of laser beam, *P*atten the power attenuated by the powder cloud, *R* the radius of laser beam spot, the laser absorption coefficient, *me* the evaporation mass flux, *Lv* the latent heat of evaporation, *hc* the heat convective coefficient, ε emissivity, σ the Stefan-Boltzmann constant, and **n** the normal vector at the local interface. *me* can be evaluated according to the "overall evaporation model" in [Choi et al., 1987], and Patten can be calculated according to Frenk et al., 1997 with a minor modification.

On the bottom surface and side surfaces, boundary conditions are given by

$$k\frac{\partial T}{\partial \mathbf{n}} + h\_c(T - T\_{\text{ov}}) = 0\tag{18}$$

$$\mathbf{V} = \mathbf{0} \tag{19}$$

Note that the radiation heat loss at these surfaces is neglected due to the fact that the temperature differences at these surfaces are not large.

#### **2.5 Numerical Implementation**

10 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

or RDF distributions [Cummins et al, 2005; Francois et al., 2006; Liovic et al., 2010]. Specifically, we adopt the HF technique in [López & Hernández, 2010] that has many improvements over earlier versions (such as that in [Torrey et al., 1985]) of HF, including using an error correction procedure to minimize estimation error. Within the HF framework, suppose the absolute value of the y-direction component of the interface normal vector is

> 2 3/2 (1 ) *xx x H H*

where *H* is the height function, *Hx* and *Hxx* are first-order and second-order derivatives of *H*, respectively. *Hx* and *Hxx* are obtained by using a finite difference formula. Interface normals

The free surface of the melt pool is tracked using the PLIC-VOF [Gueyffier et al., 1999; Scardovelli & Zaleski, 2000, 2003]. The Volume of Fluid function, F, satisfies the following

( )0 *<sup>F</sup> <sup>F</sup>*

The PLIC-VOF method consists of two steps: interface reconstruction and interface advection. In 2D calculation, a reconstructed planar surface becomes a straight line which

*nx n x y y d*

where *nx* and *ny* are x and y components of the interface normal vector. *d* is a parameter related to the distance between the line and the coordinate origin of the reference cell. In the interface reconstruction step, *nx* and *ny* of each cell are calculated based on volume fraction data, using the Least-Squares Fit method from [Aulisa et al., 2007]. Then the parameter *d* is determined to match the given volume fraction. Finally given the velocity field, the reconstructed interface is advected according to the combined Eulerian-Lagrangian scheme

> ( ) ( )( ) *laser atten <sup>c</sup> e v <sup>T</sup> P P <sup>k</sup> h T T T T mL*

where terms on the right-hand side are laser irradiation, convective heat loss, radiation heat loss and evaporation heat loss, respectively. *P*laser is the power of laser beam, *P*atten the power attenuated by the powder cloud, *R* the radius of laser beam spot, the laser absorption coefficient, *me* the evaporation mass flux, *Lv* the latent heat of evaporation, *hc* the heat

**<sup>n</sup>** (17)

(14)

**V** (15)

4 4

(16)

larger than the x-direction component, interface curvature (in 2D) is given by

are calculated based on the Least-Squares Fit method from [Aulisa et al., 2007].

*t* 

Energy balance at the free surface satisfies the following equation:

2

*R*

**2.3 Tracking of the free surface** 

satisfies the following equation:

conservation equation:

in [Aulisa et al., 2007].

**2.4 Boundary conditions** 

Finite difference and finite volume methods are used for spatial discretization of the governing equations. Staggered grids are employed where the temperatures, pressures and VOF function are located at the cell center and the velocities at the walls. In the numerical implementation, material properties play an important role. The material properties are generally dependent on temperature, concentration, and pressure. For fusion-based additive manufacturing processes, the material experiences a large variation from room temperature to above the melting temperature. For Ti-6Al-4V, many material properties experience large variations over this wide temperature range, as shown in Table 1. For example, the value of specific heat varies from 546 JK-1kg-1 at room temperature to 831 JK-1kg-1 at liquidus temperature. Thermal conductivity varies from 7 to 33.4 Wm-1K-1 over the same temperature range. Thus, the temperature dependence of the properties dominates, which necessitates a coupling of the momentum equations with the energy equation and gives rise to strong nonlinearity in the conservation equations.

The variable properties have two effects on the numerical solution procedure [Ferziger & Peric, 2002]. First, although an incompressible flow assumption is made, the thermophysical properties need to be kept inside the differential operators. Thus, solution methods for incompressible flow can be used. Second, the momentum and energy conservation equations have to be solved in a coupled way. In this study, the coupling between momentum and energy equations is achieved by the following iterative scheme:

1. Eqs. (1) - (3) and the related boundary conditions are solved iteratively using a two-step projection method [Chorin, 1968] to obtain velocities and pressures. Thermo-physical properties used in this step are computed from the old temperature field. At each time step, the discretized momentum equations calculate new velocities in terms of an estimated pressure field. Then the pressure field is iteratively adjusted and velocity changes induced by each pressure correction are added to the previous velocities. This iterative process is repeated until the continuity equation is satisfied under an imposed tolerance by the newly computed velocities. This imposes a requirement for solving a linear system of equations. The preconditioned Bi-CGSTAB (for Bi-Conjugate Gradient Stabilized) method [Barrett et al., 1994] is used to solve the linear system of equations.

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 13

The parameters for the simulation were chosen based on the capability of our experimental facilities to compare the simulation results with the experimental measurements. A diode laser deposition system (the LAMP system of Missouri S&T) and a YAG laser deposition system at South Dakota School of Mines and Technology (SDSMT) were used for simulations and experiments. Ti-6Al-V4 plates with a thickness of 0.25 inch were selected as substrates. Ti-6Al-V4 powder particles with a diameter from 40 to 140 m were used as deposit material. Fig. 3 shows the typical simulation results for temperature, velocity and

The numerical model was validated from different aspects. First, it was validated in terms of melt pool peak temperature and melt pool length. The experiments were performed on the LAMP system as shown in Fig. 4. The system consists of a diode laser, powder delivery unit, 5-axis CNC machine, and monitoring subsystem. The laser system used was a Nuvonyx ISL-1000M Diode Laser that is rated for 1 kW of output power. The laser emits at 808 nm and operates in the continuous wave (CW) mode. The laser spot size is 2.5 mm. To protect oxidization of Ti-6Al-V4, the system is covered in an environmental chamber to supply argon gas. The melt pool peak temperature is measured by a non-contact optical pyrometer that is designed for rough conditions, such as high ambient temperatures or electromagnetic interferences. A laser sight within the pyrometer allows for perfect alignment and focal length positioning; the spot size is 2.6 mm which encompasses the melt pool. The pyrometer senses the maximum temperature between 400 and 2500 (degrees C) and correlates the emissivity of the object to the resulting measurement. Temperature measurements are taken in real-time at 500 or 1000 Hz using a National Instruments real-time control system. A 4-20 mA signal is sent to the real-time system which is converted to degrees Celsius, displayed to the user and simultaneously recorded to be analyzed at a later date. Due to the collimator, the pyrometer is mounted to the Zaxis of the CNC at 42 (degrees) and is aligned with the center of the nozzle. Temperature measurements recorded the rise and steady state temperatures and the cooling rates of the melt pool. A complementary metal oxide semiconductor (CMOS) camera was installed right above the nozzle head for a better view in dynamically acquiring the melt pool image. The melt pool dimensions can be calculated from the image by the image process

Fig. 5 and Fig. 6 show the measured and predicted melt pool peak temperatures at different laser power levels and at different travel speeds, respectively. It can be seen from the plot that the general trend between simulation and experiment is consistent. At different power intensity level, there is a different error from 10 K (about 0.5%) to 121 K (about 5%). Fig. 7 shows measured and predicted melt pool length at different laser power levels. The biggest disagreement between measured and simulated values is about 7%. It can be seen that the differences between measured and predicted values at higher power intensities (higher power levels or slower travel speeds) are generally bigger than those at lower power intensities. This can be explained by the two-dimensional nature of the numerical model. A 2D model does not consider the heat transfer in the third direction. At a higher power level,

heat transfer in the third dimension is more significant.

**3. Simulation results and model validation** 

VOF function.

software.



Table 1. Material properties for Ti-6Al-4V and main process parameters used in simulations. aValue for commercially pure titanium was used.

The time step is taken at the level of 10-6 s initially and adapted subsequently according to the convergence and stability requirements of the Courant–Friedrichs–Lewy (CFL) condition, the explicit differencing of the Newtonian viscous stress tensor, and the explicit treatment of the surface tension force.

#### **3. Simulation results and model validation**

12 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

2. Eq. (4) is solved by a method [Knoll et al., 1999] based on a finite volume discretization of the enthalpy formulation of Eq. (4). The finite volume approach ensures that the numerical scheme is locally and globally conservative, while the enthalpy formulation can treat phase change in a straightforward and unified manner. Once new temperature

3. Equation (15) is solved using the PLIC-VOF [Gueyffier et al., 1999; Scardovelli & Zaleski, 2000, 2003] to obtain the updated free surface and geometry of the melt pool. 4. Advance to the next time step and back to step 1 until the desired process time is

412.7 0.1801 1268 1923

1.2595 0.0157 1268 3.5127 0.0127 1268 1923

 


*T T*

2.66 10-3 (2073K) 2.36 10-3 (2173K)

Table 1. Material properties for Ti-6Al-4V and main process parameters used in simulations.

The time step is taken at the level of 10-6 s initially and adapted subsequently according to the convergence and stability requirements of the Courant–Friedrichs–Lewy (CFL) condition, the explicit differencing of the Newtonian viscous stress tensor, and the explicit

Surface tension (N m-1) 1.525 – 0.28×10-3(T – 1941K)a [Mills, 2002]

Liquid specific heat (*J kg-1 K-1*) 831.0 [Mills, 2002]

Solid density (kg m-3) 4420 – 0.154 (T – 298 K) [Mills, 2002] Liquid density (kg m-3) 3920 – 0.68 (T – 1923 K) [Mills, 2002] Latent heat of fusion (J kg-1) 2.86 105 [Mills, 2002]

*TT K T T*

*TT K T T*

9.83 106 [Mills, 2002]

1.1 × 10-5 [Mills, 2002]

[Mills, 2002]

[Mills, 2002]

[Mills, 2002]

*10-4 (T -300.0 K)* [Lips&Fritsche, 2005]

Physical Properties Value Reference Liquidus temperature (*K*) 1923.0 [Mills, 2002] Solidus temperature (*K*) 1877.0 [Boyer et al., 1994] Evaporation temperature (*K*) 3533.0 [Boyer et al., 1994]

field is obtained, the thermo-physical properties are updated.

Solid specific heat (*J kg-1 K-1*) 483.04 0.215 1268

 

Dynamic viscosity (N m-1 s-1) 3.25 10-3 (1923K) 3.03 10-3 (1973K)

10

reached.

Thermal conductivity

Latent heat of evaporation

Thermal expansion coefficient

Convective coefficient

Laser absorption coefficient 0.4 Ambient temperature (K) 300

treatment of the surface tension force.

Radiation emissivity *0.1536+1.8377*

aValue for commercially pure titanium was used.

(W m-1 K-1)

(J kg-1)

(K-1)

(W m-2 K-1)

The parameters for the simulation were chosen based on the capability of our experimental facilities to compare the simulation results with the experimental measurements. A diode laser deposition system (the LAMP system of Missouri S&T) and a YAG laser deposition system at South Dakota School of Mines and Technology (SDSMT) were used for simulations and experiments. Ti-6Al-V4 plates with a thickness of 0.25 inch were selected as substrates. Ti-6Al-V4 powder particles with a diameter from 40 to 140 m were used as deposit material. Fig. 3 shows the typical simulation results for temperature, velocity and VOF function.

The numerical model was validated from different aspects. First, it was validated in terms of melt pool peak temperature and melt pool length. The experiments were performed on the LAMP system as shown in Fig. 4. The system consists of a diode laser, powder delivery unit, 5-axis CNC machine, and monitoring subsystem. The laser system used was a Nuvonyx ISL-1000M Diode Laser that is rated for 1 kW of output power. The laser emits at 808 nm and operates in the continuous wave (CW) mode. The laser spot size is 2.5 mm. To protect oxidization of Ti-6Al-V4, the system is covered in an environmental chamber to supply argon gas. The melt pool peak temperature is measured by a non-contact optical pyrometer that is designed for rough conditions, such as high ambient temperatures or electromagnetic interferences. A laser sight within the pyrometer allows for perfect alignment and focal length positioning; the spot size is 2.6 mm which encompasses the melt pool. The pyrometer senses the maximum temperature between 400 and 2500 (degrees C) and correlates the emissivity of the object to the resulting measurement. Temperature measurements are taken in real-time at 500 or 1000 Hz using a National Instruments real-time control system. A 4-20 mA signal is sent to the real-time system which is converted to degrees Celsius, displayed to the user and simultaneously recorded to be analyzed at a later date. Due to the collimator, the pyrometer is mounted to the Zaxis of the CNC at 42 (degrees) and is aligned with the center of the nozzle. Temperature measurements recorded the rise and steady state temperatures and the cooling rates of the melt pool. A complementary metal oxide semiconductor (CMOS) camera was installed right above the nozzle head for a better view in dynamically acquiring the melt pool image. The melt pool dimensions can be calculated from the image by the image process software.

Fig. 5 and Fig. 6 show the measured and predicted melt pool peak temperatures at different laser power levels and at different travel speeds, respectively. It can be seen from the plot that the general trend between simulation and experiment is consistent. At different power intensity level, there is a different error from 10 K (about 0.5%) to 121 K (about 5%). Fig. 7 shows measured and predicted melt pool length at different laser power levels. The biggest disagreement between measured and simulated values is about 7%. It can be seen that the differences between measured and predicted values at higher power intensities (higher power levels or slower travel speeds) are generally bigger than those at lower power intensities. This can be explained by the two-dimensional nature of the numerical model. A 2D model does not consider the heat transfer in the third direction. At a higher power level, heat transfer in the third dimension is more significant.

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 15

Fig. 5. Melt pool peak temperature comparison between simulation and experiment at different laser power levels (powder mass flow rate = 4.68g/min., travel speed = 20

Fig. 4. Schematic of experimental setup

inch/min.)

(a) Temperature field of the region around the melt pool

(b) Velocity field of the melt pool and falling powder particles

(c) VOF field of part of the region around the melt pool

Fig. 3. Simulation results of laser deposition of Ti-6Al-4V

Fig. 4. Schematic of experimental setup

(a) Temperature field of the region around the melt pool

(b) Velocity field of the melt pool and falling powder particles

(c) VOF field of part of the region around the melt pool Fig. 3. Simulation results of laser deposition of Ti-6Al-4V

Fig. 5. Melt pool peak temperature comparison between simulation and experiment at different laser power levels (powder mass flow rate = 4.68g/min., travel speed = 20 inch/min.)

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 17

Fig. 8. Comparison of dilution depth between simulation and experiment at different power

Fig. 9. Comparison of dilution depth between simulation and experiment at different travel

speeds and different laser power levels (powder mass flow rate = 4.68 g/min.)

levels (powder mass flow rate = 4.68 g/min., travel speed = 20 inch/min.)

Fig. 6. Melt pool peak temperature comparison between simulation and experiment at different travel speeds (powder mass flow rate = 4.68g/min., laser power = 700 W)

Fig. 7. Melt pool length comparisons between simulation and experiment at different power levels (powder mass flow rate = 4. 68 g/min., travel speed = 20 inch/min.)

The samples were cross-sectioned using a Wire-EDM machine to measure dilution depth. An SEM (Scanning Electron Microscope) line trace was used to determine the dilution of the clad layer. The deposited Ti-6Al-4V is of Widmansttaten structure. The substrate has a rolled equi-axed alpha plus beta structure. Even though these two structures are are easily distinguishable, the HAZ is large and has a martensitic structure that can be associated with it. Hence a small quantity of tool steel in the order of 5% was mixed with Ti-6Al-4V. The small quantity makes sure that it does not drastically change the deposit features of a 100% Ti-6Al-4V deposit. At the same time, the presence of Cr in tool steel makes it easily identifiable by means of EDS scans using SEM. Simulation and experimental results of dilution depth are shown in Figs. 8 - 10.

Fig. 6. Melt pool peak temperature comparison between simulation and experiment at different travel speeds (powder mass flow rate = 4.68g/min., laser power = 700 W)

Fig. 7. Melt pool length comparisons between simulation and experiment at different power

The samples were cross-sectioned using a Wire-EDM machine to measure dilution depth. An SEM (Scanning Electron Microscope) line trace was used to determine the dilution of the clad layer. The deposited Ti-6Al-4V is of Widmansttaten structure. The substrate has a rolled equi-axed alpha plus beta structure. Even though these two structures are are easily distinguishable, the HAZ is large and has a martensitic structure that can be associated with it. Hence a small quantity of tool steel in the order of 5% was mixed with Ti-6Al-4V. The small quantity makes sure that it does not drastically change the deposit features of a 100% Ti-6Al-4V deposit. At the same time, the presence of Cr in tool steel makes it easily identifiable by means of EDS scans using SEM. Simulation and experimental results of

levels (powder mass flow rate = 4. 68 g/min., travel speed = 20 inch/min.)

dilution depth are shown in Figs. 8 - 10.

Fig. 8. Comparison of dilution depth between simulation and experiment at different power levels (powder mass flow rate = 4.68 g/min., travel speed = 20 inch/min.)

Fig. 9. Comparison of dilution depth between simulation and experiment at different travel speeds and different laser power levels (powder mass flow rate = 4.68 g/min.)

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 19

ultrasonic graph, the distance of which is the height of the deposition and the thickness of the substrate. Fig. 12 shows an ultrasonic graph of a deposited specimen with a very good deposition. The ultrasonic result indicates there is not lack of fusion occurring between layers and the interface. The distance between two peaks is the height of the deposition and the thickness of the substrate. For the deposition as shown in Fig. 13, the lack of fusion occurs as the small peak (in circle) appears between two high peaks. The results revealed that no lack-of-fusion was detected in specimens deposited using 1,200 watts and higher energy levels. However, lack-of-fusion was detected in specimens deposited from lower energy levels (minus 10% up to minus 30% of 1,200 watts.). The test results validated the

Fig. 12. Ultrasonic graph of a laser deposited Ti-6Al-4V specimen without lack of fusion

simulation model.

Fig. 11. Ultrasonic graph of a Ti-6Al-4V substrate

Fig. 10. Comparison of dilution depth between simulation and experiment at different powder mass flow rates and different laser power levels (travel speed = 20 inch/min.)

Good agreements between measured and simulated dilution depths can be found in Figs. 8- 10. The differences are from about 4.8% to 15.1%. It can be seen that an increase in the laser power will increase the dilution depth. An increase in the laser travel speed will decrease the dilution depth. It is clear that the dilution depth has a linear dependence on the laser power and the laser travel speed. This is easy to understand. As the laser power increases, more power is available for melting the substrate. As travel speed decreases, the laser material interaction time is extended. From Fig. 10, it can be seen that an increase in powder mass flow rate will decrease the dilution depth. But this effect is more significant at a higher level of laser power. It is likely that at a lower level of laser power, a significant portion of laser energy is consumed to melt the powder. Hence the energy available is barely enough to melt the substrate. Detailed discussion can be found in [Fan et al, 2006, 2007b; Fan, 2007].

Finally, the numerical model was validated in terms of its capability for predicting the lackof-fusion defect. The test was performed using the YAG laser deposition system at South Dakota School of Mines and Technology (SDSMT). The simulation model determined that 1,200 watts would be the nominal energy level for the test. This means that based on the model, lack of fusion should occur when the laser power is below 1200 W. In accordance with the test matrix, seven energy levels were tested: nominal, nominal ± 10%, nominal ± 20%, and nominal ± 30%. Based on the predicted nominal value of 1,200 watts, the seven energy levels in the test matrix are 840, 960, 1080, 1200, 1320, 1440, and 1540 watts. The deposited Ti-6Al-4V specimens were inspected at Quality Testing Services Co. using ultrasonic and radiographic inspections to determine the extent of lack-of-fusion in the specimens. The determination of whether or not there exists lack of fusion in a deposited specimen can be explained using Figs. 11 - 13. First a substrate without deposit on it was inspected as shown in Fig. 11. Notice that the distance between two peaks are the thickness of the substrate. Then laser deposited specimens were inspected. If there is lack of fusion in a deposited specimen, some form of peaks can be found between the two high peaks in the

Fig. 10. Comparison of dilution depth between simulation and experiment at different powder mass flow rates and different laser power levels (travel speed = 20 inch/min.)

Good agreements between measured and simulated dilution depths can be found in Figs. 8- 10. The differences are from about 4.8% to 15.1%. It can be seen that an increase in the laser power will increase the dilution depth. An increase in the laser travel speed will decrease the dilution depth. It is clear that the dilution depth has a linear dependence on the laser power and the laser travel speed. This is easy to understand. As the laser power increases, more power is available for melting the substrate. As travel speed decreases, the laser material interaction time is extended. From Fig. 10, it can be seen that an increase in powder mass flow rate will decrease the dilution depth. But this effect is more significant at a higher level of laser power. It is likely that at a lower level of laser power, a significant portion of laser energy is consumed to melt the powder. Hence the energy available is barely enough to melt the substrate. Detailed discussion can be found in [Fan et al, 2006, 2007b; Fan, 2007]. Finally, the numerical model was validated in terms of its capability for predicting the lackof-fusion defect. The test was performed using the YAG laser deposition system at South Dakota School of Mines and Technology (SDSMT). The simulation model determined that 1,200 watts would be the nominal energy level for the test. This means that based on the model, lack of fusion should occur when the laser power is below 1200 W. In accordance with the test matrix, seven energy levels were tested: nominal, nominal ± 10%, nominal ± 20%, and nominal ± 30%. Based on the predicted nominal value of 1,200 watts, the seven energy levels in the test matrix are 840, 960, 1080, 1200, 1320, 1440, and 1540 watts. The deposited Ti-6Al-4V specimens were inspected at Quality Testing Services Co. using ultrasonic and radiographic inspections to determine the extent of lack-of-fusion in the specimens. The determination of whether or not there exists lack of fusion in a deposited specimen can be explained using Figs. 11 - 13. First a substrate without deposit on it was inspected as shown in Fig. 11. Notice that the distance between two peaks are the thickness of the substrate. Then laser deposited specimens were inspected. If there is lack of fusion in a deposited specimen, some form of peaks can be found between the two high peaks in the ultrasonic graph, the distance of which is the height of the deposition and the thickness of the substrate. Fig. 12 shows an ultrasonic graph of a deposited specimen with a very good deposition. The ultrasonic result indicates there is not lack of fusion occurring between layers and the interface. The distance between two peaks is the height of the deposition and the thickness of the substrate. For the deposition as shown in Fig. 13, the lack of fusion occurs as the small peak (in circle) appears between two high peaks. The results revealed that no lack-of-fusion was detected in specimens deposited using 1,200 watts and higher energy levels. However, lack-of-fusion was detected in specimens deposited from lower energy levels (minus 10% up to minus 30% of 1,200 watts.). The test results validated the simulation model.

Fig. 11. Ultrasonic graph of a Ti-6Al-4V substrate

Fig. 12. Ultrasonic graph of a laser deposited Ti-6Al-4V specimen without lack of fusion

Numerical Modeling of the Additive Manufacturing (AM) Processes of Titanium Alloy 21

Aulisa, E.; Manservisi, S.; Scardovelli, R. & Zaleski, S. (2007). Interface reconstruction with

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Fig. 13. Ultrasonic graph of a laser deposited Ti-6Al-4V specimen with lack of fusion

#### **4. Conclusion**

This chapter has outlined the approach for mathematical and numerical modeling of fusionbased additive manufacturing of titanium. The emphasis is put on modeling of transport phenomena associated with the process, including heat transfer and fluid flow dynamics. Of particular interest are the modeling approaches for solidification and free surface flow with surface tension. The advantages and disadvantages of the main modeling approaches are briefly discussed. Based on the comparisons, the continuum model is adopted for modeling of melting/solidification phase change, and the VOF method for modeling of free-surface flow in the melt pool.

The laser deposition process is selected as an example of fusion-based additive manufacturing processes. The governing equations, auxiliary relationships, and boundary conditions for the solidification system and free-surface flow are presented. The main challenge for modeling of the surface tension-dominant free surface flow is discussed and the measures to overcome the challenge are given. The numerical implementation procedures are outlined, with a focus on the effects of variable material property on the discretization schemes and solution algorithms. Finally the simulation results are presented and compared with experimental measurements. A good agreement has been obtained and thus the numerical model is validated. The modeling approach can be applied to other fusion-based manufacturing processes, such as casting and welding.

#### **5. Acknowledgment**

This research was partially supported by the National Aeronautics and Space Administration Grant Number NNX11AI73A, the grant from the U.S. Air Force Research Laboratory, and Missouri S&T's Intelligent Systems Center and Manufacturing Engineering program. Their support is greatly appreciated. The help from Dr. Kevin Slattery and Mr. Hsin Nan Chou at Boeing-St Louis and Dr. James Sears at South Dakota School of Mines and Technology are also acknowledged.

#### **6. References**

20 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Fig. 13. Ultrasonic graph of a laser deposited Ti-6Al-4V specimen with lack of fusion

This chapter has outlined the approach for mathematical and numerical modeling of fusionbased additive manufacturing of titanium. The emphasis is put on modeling of transport phenomena associated with the process, including heat transfer and fluid flow dynamics. Of particular interest are the modeling approaches for solidification and free surface flow with surface tension. The advantages and disadvantages of the main modeling approaches are briefly discussed. Based on the comparisons, the continuum model is adopted for modeling of melting/solidification phase change, and the VOF method for modeling of free-surface

The laser deposition process is selected as an example of fusion-based additive manufacturing processes. The governing equations, auxiliary relationships, and boundary conditions for the solidification system and free-surface flow are presented. The main challenge for modeling of the surface tension-dominant free surface flow is discussed and the measures to overcome the challenge are given. The numerical implementation procedures are outlined, with a focus on the effects of variable material property on the discretization schemes and solution algorithms. Finally the simulation results are presented and compared with experimental measurements. A good agreement has been obtained and thus the numerical model is validated. The modeling approach can be applied to other

This research was partially supported by the National Aeronautics and Space Administration Grant Number NNX11AI73A, the grant from the U.S. Air Force Research Laboratory, and Missouri S&T's Intelligent Systems Center and Manufacturing Engineering program. Their support is greatly appreciated. The help from Dr. Kevin Slattery and Mr. Hsin Nan Chou at Boeing-St Louis and Dr. James Sears at South Dakota School of Mines and

fusion-based manufacturing processes, such as casting and welding.

**4. Conclusion** 

flow in the melt pool.

**5. Acknowledgment** 

Technology are also acknowledged.


http://www.americanmachinist.com/304/Issue/Article/False/77297/


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**2** 

*Korea* 

**Formation of Alpha Case Mechanism** 

**on Titanium Investment Cast Parts** 

Si-Young Sung1, Beom-Suck Han1 and Young-Jig Kim2

At present the practical utilization of titanium (Ti) and its alloys ranges from sports equipment to the aerospace, power generation and chemical processing industries, and automotive, marine and medical engineering [1]. Yet, Ti alloy forming processes are much more laborious than one thinks, which is mainly due to the fact that in solid solution Ti combines with carbon, nitrogen and especially oxygen and additionally Ti is constrained to be melted in inert environments or in a vacuum [2]. For the intended applications, Ti and its alloy forming processes should incorporate forging, powder metallurgy and casting. A center bulkhead of the F-22 Raptor fighter aircraft is a good example of the difficulty in Ti alloy forging. Although the final Ti6Al4V alloy component weighs only about 150 kg, it should be forged initially from a single cast ingot of nearly 3,000 kg, which clearly exhibits extremely high machining losses in the Ti alloy forging process [2]. By the way, the powder metallurgy of Ti alloys offers a practicable way of producing complex components with less machining losses than in forging processes. However, the applications of the powder metallurgy have been limited by the freedom of size and shape. In addition, in order to ensure the soundness of the powder metallurgy products without porosity defects, hot isostatic pressing (HIP) is essentially required, which makes the powder metallurgy even

Casting is a typical net shape forming technique in which molten metals are poured into a mold to produce an object of desired shape. However, the casting of titanium alloys is considered as only a near net shape forming technique. That is because regardless of mold type, titanium alloys casting has a drawback, called alpha-case (α-case), which makes it difficult to machine them and can lend themselves to crack initiation and propagation, due to their enormous reactivity in molten states [4]. Thus, the depth of α-case must be taken into consideration in the initial design for casting since the brittle α-case must be removed by chemical milling. For the reason, the wide use of titanium alloys casting has been limited, although titanium alloys castings are comparable, and quite often superior, to wrought

In order to avoid the -case problem, the expensive ceramics, such as CaO, ZrO2 and Y2O3 have been adopted as mold materials since the standard free energy changes of the formation of their oxides are more negative than that of TiO2. Regardless of thermodynamic

**1. Introduction** 

more expensive [3].

products in all respects [3].

*1KATECH(Korea Automotive Technology Institute),* 

*2Sungkyunkwan University* 


### **Formation of Alpha Case Mechanism on Titanium Investment Cast Parts**

Si-Young Sung1, Beom-Suck Han1 and Young-Jig Kim2 *1KATECH(Korea Automotive Technology Institute), 2Sungkyunkwan University Korea* 

#### **1. Introduction**

28 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

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prediction of dendritic growth with melt convection. *Physical Review E*, Vol. 69,

At present the practical utilization of titanium (Ti) and its alloys ranges from sports equipment to the aerospace, power generation and chemical processing industries, and automotive, marine and medical engineering [1]. Yet, Ti alloy forming processes are much more laborious than one thinks, which is mainly due to the fact that in solid solution Ti combines with carbon, nitrogen and especially oxygen and additionally Ti is constrained to be melted in inert environments or in a vacuum [2]. For the intended applications, Ti and its alloy forming processes should incorporate forging, powder metallurgy and casting. A center bulkhead of the F-22 Raptor fighter aircraft is a good example of the difficulty in Ti alloy forging. Although the final Ti6Al4V alloy component weighs only about 150 kg, it should be forged initially from a single cast ingot of nearly 3,000 kg, which clearly exhibits extremely high machining losses in the Ti alloy forging process [2]. By the way, the powder metallurgy of Ti alloys offers a practicable way of producing complex components with less machining losses than in forging processes. However, the applications of the powder metallurgy have been limited by the freedom of size and shape. In addition, in order to ensure the soundness of the powder metallurgy products without porosity defects, hot isostatic pressing (HIP) is essentially required, which makes the powder metallurgy even more expensive [3].

Casting is a typical net shape forming technique in which molten metals are poured into a mold to produce an object of desired shape. However, the casting of titanium alloys is considered as only a near net shape forming technique. That is because regardless of mold type, titanium alloys casting has a drawback, called alpha-case (α-case), which makes it difficult to machine them and can lend themselves to crack initiation and propagation, due to their enormous reactivity in molten states [4]. Thus, the depth of α-case must be taken into consideration in the initial design for casting since the brittle α-case must be removed by chemical milling. For the reason, the wide use of titanium alloys casting has been limited, although titanium alloys castings are comparable, and quite often superior, to wrought products in all respects [3].

In order to avoid the -case problem, the expensive ceramics, such as CaO, ZrO2 and Y2O3 have been adopted as mold materials since the standard free energy changes of the formation of their oxides are more negative than that of TiO2. Regardless of thermodynamic

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 31

were inspected and dressed to eliminate any imperfection or contamination, and coated with Al2O3 slurry. The Al2O3 shell molds were dried at a controlled temperature (298 K ± 1 K) and a relative humidity (40% ±1%) for 4 hrs. Dipping, stuccoing and drying procedures were repeated three times. After the primary layer coating, the patterns were coated with the back-up layers by the chamotte. To prevent the shell cracks, the dewaxing process of the shell molds were carried out at around 423 K and 0.5 MPa in a steam autoclave. Finally, the

Fig. 2. Schematic diagram of drop casting procedure of titanium with a plasma arc melting

In order to prevent any contamination from refractory crucibles, the specimens for the of case formation examination were prepared in a plasma arc melting (PAM) furnace with drop casting procedure. The pressure in the PAM was controlled at about 1.33×10-1Pa by a rotary pump and then in order to minimize the effect of oxygen contamination, high purity

After the drop casting, Ti castings were taken out of the molds, sectioned, polished and etched using a Keller solution [5]. The microstructure in the reaction region of the castings was observed using Olympus PME3 microscopy, and its hardness was measured using a Mitutoyo MVK-H2 microvickers hardness tester with the condition of 100 g load and 50 µm intervals. For a closer examination of the alpha-case formation, the distribution state of the elements analysis at the alpha-case region was performed by SHIMADZU EPMA 1600 and

In this research, Al2O3 was selected for the mold material because the standard free energy changes of the formation of its oxides are more negative than that of TiO2. In addition, Al2O3

the phase structures of reaction products were identified by JEOL JEM-3011 TEM.

**3. Thermodynamic calculation of alpha-case formation mechanism** 

shell molds were fired at 1,223 K for 2 hrs.

argon was backfilled to the pressure of 4.9×103 Pa.

**3.1 Evaluation of the alpha-case reaction** 

furnace.

approaches, the -case formation reaction still remains to be eliminated in the chemical milling processes. In order to develop the economic net shape forming technique of titanium alloys by casting process, it is necessary to take a much closer look into the -case formation mechanism. Therefore, the exact -case formation mechanism must be examined and the development of -case controlled mold materials is required for practical applications of Ti alloys.

#### **2. Evaluation of the alpha-case formation mechanism**

#### **2.1 Conventional alpha-case formation mechanism**

The investment casting of titanium has a drawback, called '-case', which makes it difficult to machine and can lead to crack initiation and propagation, due to their enormous reactivity in molten states [4]. The -case is generally known to be developed by the interstitials such as carbon, nitrogen and especially oxygen dissolved from mold materials. In order to avoid this problem, the expensive ceramics, such as ZrO2, ZrSiO4, CaZrO3, CaO and Y2O3 have been adopted for mold materials because their standard free energy changes of the formation of oxides are more negative than that of TiO2 as shown in figure 1.

Fig. 1. Standard free energy changes of the formation of various oxides.

Regardless of the thermodynamic approaches, some amount of -case still remains to be eliminated with the complex chemical milling processes except for Y2O3 mold. However, in the case of Y2O3 mold, there is not enough strength to handling due to the silica-free binder and there occur metal-mold reactions to some extent, too.

#### **2.2 Evaluation of alpha-case reaction**

The wax patterns for the examination of -case reactions were made by pouring molten wax into a simple cylindrical silicon rubber mold (Ø15 × 70 mm). Subsequently, the patterns

approaches, the -case formation reaction still remains to be eliminated in the chemical milling processes. In order to develop the economic net shape forming technique of titanium alloys by casting process, it is necessary to take a much closer look into the -case formation mechanism. Therefore, the exact -case formation mechanism must be examined and the development of

The investment casting of titanium has a drawback, called '-case', which makes it difficult to machine and can lead to crack initiation and propagation, due to their enormous reactivity in molten states [4]. The -case is generally known to be developed by the interstitials such as carbon, nitrogen and especially oxygen dissolved from mold materials. In order to avoid this problem, the expensive ceramics, such as ZrO2, ZrSiO4, CaZrO3, CaO and Y2O3 have been adopted for mold materials because their standard free energy changes


of the formation of oxides are more negative than that of TiO2 as shown in figure 1.

Fig. 1. Standard free energy changes of the formation of various oxides.

and there occur metal-mold reactions to some extent, too.

**2.2 Evaluation of alpha-case reaction** 

Regardless of the thermodynamic approaches, some amount of -case still remains to be eliminated with the complex chemical milling processes except for Y2O3 mold. However, in the case of Y2O3 mold, there is not enough strength to handling due to the silica-free binder

The wax patterns for the examination of -case reactions were made by pouring molten wax into a simple cylindrical silicon rubber mold (Ø15 × 70 mm). Subsequently, the patterns

**2. Evaluation of the alpha-case formation mechanism** 

**2.1 Conventional alpha-case formation mechanism** 

were inspected and dressed to eliminate any imperfection or contamination, and coated with Al2O3 slurry. The Al2O3 shell molds were dried at a controlled temperature (298 K ± 1 K) and a relative humidity (40% ±1%) for 4 hrs. Dipping, stuccoing and drying procedures were repeated three times. After the primary layer coating, the patterns were coated with the back-up layers by the chamotte. To prevent the shell cracks, the dewaxing process of the shell molds were carried out at around 423 K and 0.5 MPa in a steam autoclave. Finally, the shell molds were fired at 1,223 K for 2 hrs.

Fig. 2. Schematic diagram of drop casting procedure of titanium with a plasma arc melting furnace.

In order to prevent any contamination from refractory crucibles, the specimens for the of case formation examination were prepared in a plasma arc melting (PAM) furnace with drop casting procedure. The pressure in the PAM was controlled at about 1.33×10-1Pa by a rotary pump and then in order to minimize the effect of oxygen contamination, high purity argon was backfilled to the pressure of 4.9×103 Pa.

After the drop casting, Ti castings were taken out of the molds, sectioned, polished and etched using a Keller solution [5]. The microstructure in the reaction region of the castings was observed using Olympus PME3 microscopy, and its hardness was measured using a Mitutoyo MVK-H2 microvickers hardness tester with the condition of 100 g load and 50 µm intervals. For a closer examination of the alpha-case formation, the distribution state of the elements analysis at the alpha-case region was performed by SHIMADZU EPMA 1600 and the phase structures of reaction products were identified by JEOL JEM-3011 TEM.

#### **3. Thermodynamic calculation of alpha-case formation mechanism**

#### **3.1 Evaluation of the alpha-case reaction**

In this research, Al2O3 was selected for the mold material because the standard free energy changes of the formation of its oxides are more negative than that of TiO2. In addition, Al2O3

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 33

 Ti(*l*)+2/3Al2O3(s)=TiO2(s)+4/3Al(*l*) ΔGoF= +99.748 kJ/mol (3) The above calculations utilized the joint of army-navy-air force (JANAF) thermochemical tables [6]. According to the equation (3), the -case formation by interstitial oxygen cannot occur spontaneously. Therefore, the reason why the -case reaction is generated cannot be

For the clear examination of the -case formation mechanism, the distribution of the elements of the -case and its chemical composition were investigated by EPMA elemental mapping as shown in Fig. 4. On the -case region, the oxygen element was uniformly distributed. Also, the Si element originating from the colloidal silica binder was scarcely

Fig. 4. Comparison of elemental mapping images of O, Al and Si in Ti castings into Al2O3

the effect of metallic elements cannot be overlooked any more.

However, the concentrated Al contamination layer about 30 µm thick was detected on the interface. The EPMA mapping result shows that not only the interstitial oxygen elements but also the substitutional Al elements dissolved from the mold material affect the metalmold reactions. Until recently, the effect of substitutional metallic element dissolved from mold materials has been ignored as negligibly small [7]. However, Fig. 3 and 4 indicate that

The phase identification of the detected Al element will be the very core of -case formation mechanism. The phases of the detected Al on the interface were examined by transmission

explained by the conventional -case formation mechanism.

**3.2 Alpha-case formation mechanism** 

detected on the surface.

mold and BEI image by EPMA.

features suitable strength, permeability and collapse-ability, which can ensure dimensional accuracy of castings. Fig. 3 shows the microstructure and hardness profile on the surface of pure titanium investment castings poured into Al2O3 mold.

Fig. 3. Microstructure of the interface between Ti and Al2O3 mold, and hardness profile.

The microstructure of a distinct reaction layer is about 200 µm thick on the interface between Ti and Al2O3 mold. In addition to the reaction layer, there shows a hardened layer about 300 µm thick. The reaction layer and the hardened layer together are called the -case of titanium castings. However, when the molten titanium (around 2,000 K) is poured into the Al2O3 investment mold, and if the -case results between Ti and the interstitial oxygen, the reaction could be described as follows:

$$\text{Ti(l)} + \text{O}\_2(\text{g}) = \text{TiO}\_2(\text{s}) \quad \Delta\text{G} \\ \text{TiO}\_2 = \text{-585.830 kJ/mol} \tag{1}$$

$$\text{4/3Al(l)} + \text{O}\_2(\text{g}) = 2/3Al \cdot \text{O}\_2(\text{s}) \text{ } \Delta \text{G}^\circ \text{AlZnO} = 1,028.367 \text{ kJ/mol} \tag{2}$$

$$\text{Ti(l)} + 2/3 \text{Al}\_2\text{O}\_3(s) = \text{TiO}\_2(s) + 4/3 \text{Al(l)} \ \Delta \text{Ge}\_{\overline{\mathbb{F}}} = +99.748 \text{ kJ/mol} \tag{3}$$

The above calculations utilized the joint of army-navy-air force (JANAF) thermochemical tables [6]. According to the equation (3), the -case formation by interstitial oxygen cannot occur spontaneously. Therefore, the reason why the -case reaction is generated cannot be explained by the conventional -case formation mechanism.

#### **3.2 Alpha-case formation mechanism**

32 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

features suitable strength, permeability and collapse-ability, which can ensure dimensional accuracy of castings. Fig. 3 shows the microstructure and hardness profile on the surface of

Fig. 3. Microstructure of the interface between Ti and Al2O3 mold, and hardness profile.

reaction could be described as follows:

The microstructure of a distinct reaction layer is about 200 µm thick on the interface between Ti and Al2O3 mold. In addition to the reaction layer, there shows a hardened layer about 300 µm thick. The reaction layer and the hardened layer together are called the -case of titanium castings. However, when the molten titanium (around 2,000 K) is poured into the Al2O3 investment mold, and if the -case results between Ti and the interstitial oxygen, the

Ti(*l*)+O2(g)=TiO2(s) ΔGoTiO2= -585.830 kJ/mol (1)

4/3Al(*l*)+O2(g)=2/3Al2O3(s) ΔGoAl2O3= 1,028.367 kJ/mol (2)

pure titanium investment castings poured into Al2O3 mold.

For the clear examination of the -case formation mechanism, the distribution of the elements of the -case and its chemical composition were investigated by EPMA elemental mapping as shown in Fig. 4. On the -case region, the oxygen element was uniformly distributed. Also, the Si element originating from the colloidal silica binder was scarcely detected on the surface.

Fig. 4. Comparison of elemental mapping images of O, Al and Si in Ti castings into Al2O3 mold and BEI image by EPMA.

However, the concentrated Al contamination layer about 30 µm thick was detected on the interface. The EPMA mapping result shows that not only the interstitial oxygen elements but also the substitutional Al elements dissolved from the mold material affect the metalmold reactions. Until recently, the effect of substitutional metallic element dissolved from mold materials has been ignored as negligibly small [7]. However, Fig. 3 and 4 indicate that the effect of metallic elements cannot be overlooked any more.

The phase identification of the detected Al element will be the very core of -case formation mechanism. The phases of the detected Al on the interface were examined by transmission

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 35

primitive cell volume is characteristic material constant. So, the phase of interest can be easily identified by comparing the measured primitive cell volume of an unknown phase using CBED patterns with the known values of the possible phases. This method is accurate (error range of <10%) for the phase identification. In the CBED pattern where zero order Laue zone (ZOLZ) disk and high order Laue zone (HOLZ) ring are present, by measuring the distance of ZOLZ disk from the transmitted beam to the diffracted beam (D1, D2), an internal angle (ANG) between D1 and D2, and HOLZ ring's radius (CRAD), the primitive cell volume of unknown phase can be easily determined using the equation (4). Camera

1 2 sin 1 cos tan

Fig. 7 (a) shows the CBED pattern of the HCP phase. There are two possible HCP phases of Ti and Ti3Al. The measured primitive cell volume (61.08 ang3, CL=521.4 mm, D1=D2=6.5 mm, ANG=60º, CRAD=23 mm) from CBED pattern corresponded to the theoretic value of Ti3Al (66.60 ang3). This phase identification was in accordance with EDS spot analysis as shown in Fig. 6 (b). The phase of the detected Al from EPMA mapping was Ti3Al phase. Thus, considering the microstructure, hardness profile, EPMA mapping and TEM, the -

2 3

*CRAD D D ANG*

1

*CL*

(4)

Length (CL) was calibrated with Au standard sample at 200 keV.

*CL Cell volume*

case is not wholly TiO2, but TiO2 and Ti3Al between Ti and Al2O3 mold.

(a) (b)

mm, ANG=60º, CRAD=23 mm, λ=0.0251 Å) and (b) EDS spot analysis result.

Fig. 8.

Fig. 7. Convergent beam electron diffraction analysis with the primitive unit cell volume method (a) The measured primitive unit cell volume, 61.08 ang3 (CL=521.4 mm, D1=D2=6.5

In this study, from the synthesis of the microstructure, hardness profile, EPMA mapping and TEM, it could be confirmed that the -case is formed by not only interstitial oxygen element but also substitutional metallic elements dissolved from mold materials as shown in

electron microscopy (TEM). In order to observe -case region precisely, the titanium castings specimen was grinded from inside to the surface until about 30 µm and final thinning was carried out using ion milling. Fig. 5 is the cross-sectional bright field TEM image of the -case region.

Fig. 5. Bright field TEM image of the -case region.

To examine what kinds of phases are presented in the -case region, C2 aperture was temporarily removed. Fig. 6 (a) shows ring and spot patterns on the TEM image without C2 aperture. The definite contrast of continuous ring pattern could not be found on the TEM image since the ring pattern was an extremely small size polycrystalline TiO2 phase. In the case of spot patterns on Fig. 6 (b), the contrast could be distinguished on the TEM image. And the indexed pattern was a hexagonal close-packed (HCP) phase in the [2īī0] beam direction as shown Fig. 6 (b).

Fig. 6. Results of analytical transmission electron microscopy of (a) ring and spot patterns on the TEM image without C2 aperture on the bright field image, and (b) spot pattern on the TEM image was HCP phase in the [2īī0] beam direction.

The convergent beam electron diffraction (CBED) analysis was carried out for the verification of the phase of spot patterns with the primitive cell volume method [7]. The

electron microscopy (TEM). In order to observe -case region precisely, the titanium castings specimen was grinded from inside to the surface until about 30 µm and final thinning was carried out using ion milling. Fig. 5 is the cross-sectional bright field TEM

To examine what kinds of phases are presented in the -case region, C2 aperture was temporarily removed. Fig. 6 (a) shows ring and spot patterns on the TEM image without C2 aperture. The definite contrast of continuous ring pattern could not be found on the TEM image since the ring pattern was an extremely small size polycrystalline TiO2 phase. In the case of spot patterns on Fig. 6 (b), the contrast could be distinguished on the TEM image. And the indexed pattern was a hexagonal close-packed (HCP) phase in the [2īī0] beam

Fig. 6. Results of analytical transmission electron microscopy of (a) ring and spot patterns on the TEM image without C2 aperture on the bright field image, and (b) spot pattern on the

The convergent beam electron diffraction (CBED) analysis was carried out for the verification of the phase of spot patterns with the primitive cell volume method [7]. The

image of the -case region.

direction as shown Fig. 6 (b).

Fig. 5. Bright field TEM image of the -case region.

TEM image was HCP phase in the [2īī0] beam direction.

(a) (b)

primitive cell volume is characteristic material constant. So, the phase of interest can be easily identified by comparing the measured primitive cell volume of an unknown phase using CBED patterns with the known values of the possible phases. This method is accurate (error range of <10%) for the phase identification. In the CBED pattern where zero order Laue zone (ZOLZ) disk and high order Laue zone (HOLZ) ring are present, by measuring the distance of ZOLZ disk from the transmitted beam to the diffracted beam (D1, D2), an internal angle (ANG) between D1 and D2, and HOLZ ring's radius (CRAD), the primitive cell volume of unknown phase can be easily determined using the equation (4). Camera Length (CL) was calibrated with Au standard sample at 200 keV.

$$\text{Cell volume} = \frac{\text{CL}^2 \cdot \text{\AA}^3}{D\_1 \cdot D\_2 \cdot \sin\left(\text{ANG}\right) \left[1 - \cos\left\{\tan^{-1}\left(\frac{\text{CR}AD}{CL}\right)\right\}\right]} \tag{4}$$

Fig. 7 (a) shows the CBED pattern of the HCP phase. There are two possible HCP phases of Ti and Ti3Al. The measured primitive cell volume (61.08 ang3, CL=521.4 mm, D1=D2=6.5 mm, ANG=60º, CRAD=23 mm) from CBED pattern corresponded to the theoretic value of Ti3Al (66.60 ang3). This phase identification was in accordance with EDS spot analysis as shown in Fig. 6 (b). The phase of the detected Al from EPMA mapping was Ti3Al phase. Thus, considering the microstructure, hardness profile, EPMA mapping and TEM, the case is not wholly TiO2, but TiO2 and Ti3Al between Ti and Al2O3 mold.

Fig. 7. Convergent beam electron diffraction analysis with the primitive unit cell volume method (a) The measured primitive unit cell volume, 61.08 ang3 (CL=521.4 mm, D1=D2=6.5 mm, ANG=60º, CRAD=23 mm, λ=0.0251 Å) and (b) EDS spot analysis result.

In this study, from the synthesis of the microstructure, hardness profile, EPMA mapping and TEM, it could be confirmed that the -case is formed by not only interstitial oxygen element but also substitutional metallic elements dissolved from mold materials as shown in Fig. 8.

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 37

Fig. 9. Illustration of the diffraction patterns and replots for identifying phases in 50 wt%

Fig. 10. X-ray diffraction pattern of synthesized -case controlled stable mold materials.

The scan angles were from 20° to 80° and the pattern was verified with JCPDS card numbers 88-0826 (Al2O3), 84-1284 (TiO2) and 14-0451 (Ti3Al) [8]. No trace of the starting titanium reagent was detected by XRD analysis. Fig. 11 shows the variation of XRD patterns with the amount of titanium powders. The component ratio of TiO2 and Ti3Al phases increased

titanium and Al2O3 mixed with colloidal silica before curing.

Fig. 8. Schematic diagram of the interstitial and substitutional -case formation mechanism of titanium castings.

#### **4. New alpha-case formation mechanism with alpha-case controlled mold**

#### **4.1 Alpha-case controlled stable mold fabrication**

The -case formation mechanism was applied to the development of -case controlled stable mold and verified by titanium alloy castings with the -case controlled mold materials. To synthesize the -case formation products, 10 wt% to 50 wt% of about 45 µm titanium powders were added into Al2O3 for the purpose of eliminating the -case. After blending, the slurry was prepared with the Al2O3 based powder. The -case controlled stable mold manufacturing and titanium casting procedures were carried out at the same condition as mentioned above. The synthesized phases between titanium and Al2O3 were identified by RIGAKU X-ray diffractometer.

#### **4.2 New alpha-case formation mechanism**

Although the above mentioned Al2O3 has the proper strength, permeability and collapseability enough to ensure dimensional accuracy of castings, it could not be applied for mold due to its severe -case formation. However, in this study, for the maximization of performance and cost-effectiveness of the investment mold, the -case controlled mold was designed on Al2O3 base.

In consideration of the interstitial and substitutional elements in -case formation mechanism, 10 wt% to 50 wt% of titanium powders were blended with Al2O3 for the previous synthesis of -case reaction products into the mold. Then the blended powders were mixed with the colloidal silica and agitated.

Fig. 9 shows the X-ray diffraction result of 50 wt% titanium mixed Al2O3 powders before curing. The composition of the starting titanium and Al2O3 remained unchanged after being mixed with colloidal silica. After the mixing and agitation, the mold materials were cured at 1,223 K for 2 hrs. According to XRD analysis, TiO2 and Ti3Al phases were in-situ synthesized on the Al2O3 base between titanium and Al2O3 powders as shown in Fig. 10.

Fig. 8. Schematic diagram of the interstitial and substitutional -case formation mechanism

**4. New alpha-case formation mechanism with alpha-case controlled mold** 

The -case formation mechanism was applied to the development of -case controlled stable mold and verified by titanium alloy castings with the -case controlled mold materials. To synthesize the -case formation products, 10 wt% to 50 wt% of about 45 µm titanium powders were added into Al2O3 for the purpose of eliminating the -case. After blending, the slurry was prepared with the Al2O3 based powder. The -case controlled stable mold manufacturing and titanium casting procedures were carried out at the same condition as mentioned above. The synthesized phases between titanium and Al2O3 were

Although the above mentioned Al2O3 has the proper strength, permeability and collapseability enough to ensure dimensional accuracy of castings, it could not be applied for mold due to its severe -case formation. However, in this study, for the maximization of performance and cost-effectiveness of the investment mold, the -case controlled mold was

In consideration of the interstitial and substitutional elements in -case formation mechanism, 10 wt% to 50 wt% of titanium powders were blended with Al2O3 for the previous synthesis of -case reaction products into the mold. Then the blended powders

Fig. 9 shows the X-ray diffraction result of 50 wt% titanium mixed Al2O3 powders before curing. The composition of the starting titanium and Al2O3 remained unchanged after being mixed with colloidal silica. After the mixing and agitation, the mold materials were cured at 1,223 K for 2 hrs. According to XRD analysis, TiO2 and Ti3Al phases were in-situ synthesized

on the Al2O3 base between titanium and Al2O3 powders as shown in Fig. 10.

**4.1 Alpha-case controlled stable mold fabrication** 

identified by RIGAKU X-ray diffractometer.

**4.2 New alpha-case formation mechanism** 

were mixed with the colloidal silica and agitated.

designed on Al2O3 base.

of titanium castings.

Fig. 9. Illustration of the diffraction patterns and replots for identifying phases in 50 wt% titanium and Al2O3 mixed with colloidal silica before curing.

Fig. 10. X-ray diffraction pattern of synthesized -case controlled stable mold materials.

The scan angles were from 20° to 80° and the pattern was verified with JCPDS card numbers 88-0826 (Al2O3), 84-1284 (TiO2) and 14-0451 (Ti3Al) [8]. No trace of the starting titanium reagent was detected by XRD analysis. Fig. 11 shows the variation of XRD patterns with the amount of titanium powders. The component ratio of TiO2 and Ti3Al phases increased

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 39

Fig. 12. Comparison of externals of titanium castings (a) CaO stabilized ZrO2 mold and (b)

 Fig. 13. Microstructure and hardness profile between pure titanium and -case controlled

This external examination result of the -case controlled stable mold is in good accordance with the microstructure and hardness profile as shown in Fig. 13. The effect of prevention of

The -case controlled stable mold is much less expensive than the conventional mold materials for titanium such as ZrO2, CaO stabilized ZrO2 and Y2O3, and the complete control of -case is possible. However, in this study, the noble synthesis route was conceived, which is more cost-

The homogeneous powders containing 30 mol% TiO2 powders (anatase, 45 µm, 99% pure) and 70 mol% aluminum powders were prepared by blending. After blending, the powders were mixed with colloidal silica and agitated. The -case controlled stable mold


effective than the above route by the inexpensive TiO2 and aluminum raw materials.

**4.3 Noble synthesis of the alpha-case controlled stable mold** 

the -case controlled stable mold.

stable mold.

gradually with the addition of titanium. Consequently, the synthesis of the -case formation products TiO2 and Ti3Al phases into Al2O3 based mold could be possible by the addition of titanium powders.

The effect of -case controlled stable mold which contained the interstitial metal-mold reaction product TiO2 and the substitutional metal-mold reaction product Ti3Al was verified with titanium casting.

Fig. 11. Illustration of the diffraction patterns and replots for identifying phases from 10 wt% to 50 wt% titanium and Al2O3 after curing at 1,223 K for 2 hrs.

Fig. 12 (a) indicates that in the case of CaO stabilized ZrO2, the expensive and thermally stable mold, the externals of titanium castings lost metallic luster as a result of the -case formation reactions. However, in the -case controlled stable mold, its characteristic luster of titanium was well preserved as shown in Fig. 12 (b).

gradually with the addition of titanium. Consequently, the synthesis of the -case formation products TiO2 and Ti3Al phases into Al2O3 based mold could be possible by the addition of

The effect of -case controlled stable mold which contained the interstitial metal-mold reaction product TiO2 and the substitutional metal-mold reaction product Ti3Al was verified

Fig. 11. Illustration of the diffraction patterns and replots for identifying phases from 10

Fig. 12 (a) indicates that in the case of CaO stabilized ZrO2, the expensive and thermally stable mold, the externals of titanium castings lost metallic luster as a result of the -case formation reactions. However, in the -case controlled stable mold, its characteristic luster

wt% to 50 wt% titanium and Al2O3 after curing at 1,223 K for 2 hrs.

of titanium was well preserved as shown in Fig. 12 (b).

titanium powders.

with titanium casting.

Fig. 12. Comparison of externals of titanium castings (a) CaO stabilized ZrO2 mold and (b) the -case controlled stable mold.

Fig. 13. Microstructure and hardness profile between pure titanium and -case controlled stable mold.

This external examination result of the -case controlled stable mold is in good accordance with the microstructure and hardness profile as shown in Fig. 13. The effect of prevention of -case formation can be obtained after addition of more than 2 wt% titanium.

#### **4.3 Noble synthesis of the alpha-case controlled stable mold**

The -case controlled stable mold is much less expensive than the conventional mold materials for titanium such as ZrO2, CaO stabilized ZrO2 and Y2O3, and the complete control of -case is possible. However, in this study, the noble synthesis route was conceived, which is more costeffective than the above route by the inexpensive TiO2 and aluminum raw materials.

The homogeneous powders containing 30 mol% TiO2 powders (anatase, 45 µm, 99% pure) and 70 mol% aluminum powders were prepared by blending. After blending, the powders were mixed with colloidal silica and agitated. The -case controlled stable mold manufacturing and titanium casting procedures were carried out at the same condition as mentioned above. The 3:7 molar ratio of TiO2 and aluminum was chosen to yield the Al2O3 and TiAl final product after in-situ synthesis according the following reaction.

$$\text{3TiO}\_2 + 7\text{Al} \rightarrow 2\text{Al}\_2\text{O}\_3 + 3\text{TiAl} \tag{5}$$

Formation of Alpha Case Mechanism on Titanium Investment Cast Parts 41

According to the XRD analysis as shown in Fig. 14, the major phases are anatase TiO2 and rutile TiO2 as well as the in-situ synthesized Al2O3, Ti3Al and TiAl phases since the partially non-reacted anatase TiO2 remained and the TiAl, Ti3Al intermetallic compounds and rutile TiO2 were synthesized between anatase TiO2 and aluminum powders. As a result, the synthesis which is similar to the titanium powders added -case controlled stable mold can be obtained by the economical route. The -case free titanium casting can be possible by the noble route synthesized -case controlled stable mold which contained the -case formation

By the conventional -case formation mechanism regardless of thermodynamic approaches, the alpha-case generation cannot be explained. In order to ascertain the reason, -case formation mechanism was closely examined and from the mechanism, two kinds of -case

1. Regardless of thermodynamic approaches, about 500 µm thick -case was generated between titanium and Al2O3 mold. The reason why the -case generated cannot be explained by the conventional -case formation mechanism, which is known to be

3. Considering the microstructure, hardness profile, EPMA mapping and TEM analysis, it could be confirmed that the -case is formed not only by interstitial oxygen element but

4. The synthesis of the TiO2 and Ti3Al phases which are the -case reaction products between titanium and Al2O3 mold can be obtained by the simple curing of Al2O3 mold added titanium powders. The complete control of -case formation can be possible with

5. The -case free titanium casting can be possible by the noble route synthesized -case controlled stable mold which is in-situ synthesized between TiO2 and aluminum powders to obtain the -case formation product such as Al2O3, TiO2, TiAl and Ti3Al

6. Consequently, the economical net-shape forming of titanium alloys without -case formation and the verification of a newly established -case formation mechanism, can

[1] Matthew JD.; *Titanium a Technical Guide*, ASM International, ISBN 978-0871706867, OHIO

[2] Christoph L. & Manfred P.; *Ttitanium and Titanium Alloys* : Wiley-VCH, ISBN 3-527-

[3] Joseph RD. (editor) ; *Metals Handbook Vol. 2,* ASM International, ISBN 0871703785, OHIO

formed by the interstitials, especially oxygen dissolved from mold materials. 2. In spite of having used pure titanium, the concentrated aluminum contamination layer about 30 µm thick was detected on the interface by the EPMA elemental mapping. From the results of the TEM phase identification, the phase of the detected aluminum from

also by substitutional metallic elements dissolved from mold materials.

products such as Al2O3, TiO2, TiAl and Ti3Al phases as shown in Fig. 15.

**5. Conclusion** 

phases.

**6. References** 

USA

USA

controlled stable molds were developed.

the -case controlled stable mold.

30534-3, Weinheim, Germany

EPMA mapping was identified as Ti3Al phase.

be possible using the -case controlled stable molds.

In order to synthesize the Al2O3 and TiAl phases, the mixed TiO2 and aluminum powders were cured at the 1,323 K for 2 hrs. However, the equation (5) is an ideal reaction, and thus it is very difficult to synthesize the ideal product by the simple blending and after agitation.

Fig. 14. Illustration of the diffraction patterns and replots for identifying the synthesized phases between TiO2 and aluminum powders.

Fig. 15. Microstructure and hardness profile between pure titanium and noble route -case controlled stable mold.

According to the XRD analysis as shown in Fig. 14, the major phases are anatase TiO2 and rutile TiO2 as well as the in-situ synthesized Al2O3, Ti3Al and TiAl phases since the partially non-reacted anatase TiO2 remained and the TiAl, Ti3Al intermetallic compounds and rutile TiO2 were synthesized between anatase TiO2 and aluminum powders. As a result, the synthesis which is similar to the titanium powders added -case controlled stable mold can be obtained by the economical route. The -case free titanium casting can be possible by the noble route synthesized -case controlled stable mold which contained the -case formation products such as Al2O3, TiO2, TiAl and Ti3Al phases as shown in Fig. 15.

#### **5. Conclusion**

40 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

manufacturing and titanium casting procedures were carried out at the same condition as mentioned above. The 3:7 molar ratio of TiO2 and aluminum was chosen to yield the Al2O3

 3TiO2 + 7Al → 2Al2O3 + 3TiAl (5) In order to synthesize the Al2O3 and TiAl phases, the mixed TiO2 and aluminum powders were cured at the 1,323 K for 2 hrs. However, the equation (5) is an ideal reaction, and thus it is very difficult to synthesize the ideal product by the simple blending and after agitation.

Fig. 14. Illustration of the diffraction patterns and replots for identifying the synthesized

Fig. 15. Microstructure and hardness profile between pure titanium and noble route -case

phases between TiO2 and aluminum powders.

controlled stable mold.

and TiAl final product after in-situ synthesis according the following reaction.

By the conventional -case formation mechanism regardless of thermodynamic approaches, the alpha-case generation cannot be explained. In order to ascertain the reason, -case formation mechanism was closely examined and from the mechanism, two kinds of -case controlled stable molds were developed.


#### **6. References**


**3** 

*Russia* 

**Genesis of Gas Containing** 

*2VSMPO-AVISMA Corporation,* 

**Defects in Cast Titanium Parts** 

Vladimir Vykhodets1, Tatiana Kurennykh2 and Nataliya Tarenkova2

At present, gas-containing inclusions of metallurgical origin like Ti-O, Ti-N, and Ti-O-N are one of the main problems of titanium production. These defects negatively influence the mechanical properties of titanium alloys, as the mechanical properties (hardness and plasticity) of gas-containing inclusions differ appreciably from the corresponding characteristics of alloy matrices. Therefore, cracks nucleate in defect-matrix contact zone upon loading, which results in catastrophic fracture of heavily loaded parts. For example, the tragic consequences of failures of aircraft engine parts owing to the presence of similar effects are well known. The technology of vacuum arc remelting (VAR) of titanium alloys is most vulnerable from the view point of the formation of gas-containing defects, and, at the same time, it dominates in the overall production of titanium alloys. Efforts over many years have not resulted in the development of a technology for the defect-free production of titanium alloys by the VAR method and the titanium defect problem known also as "hard alpha" remains a troublesome aspect of the use of titanium alloys. Attempts to remove or modify the hard-alpha inclusions by diffusion homogenization of ingots had failed due to the long time required for the appropriate heat treatment at high temperature. Therefore, the genesis of defects under the conditions of industrial processes, their identification in each specific case, classification, and the elaboration of measures for reducing the formation

A generally accepted concept has been developed (Bellot & Mitchel, 1994), according to which the sources of defects are particles of titanium and its alloys with a high concentration of oxygen and nitrogen atoms (the predominant element is nitrogen) and liquidus temperatures exceeding the smelting process temperature. Such particles can be formed at different stages of the preparation of batch materials for smelting. It may be supposed that all the components of a batch, such as titanium sponge, titanium return production wastes, and titanium containing alloy additions, are potential sources of defects. Smelting products with rather high concentrations of oxygen and nitrogen atoms may also be sources of defects. In melting oxygen and nitrogen atoms pass into a melt due to diffusion, the melting temperature of particles enriched in light elements decreases, the inclusion size decreases too, and conditions for the their dissolution are created. This scenario takes place for most particles enriched in oxygen and nitrogen atoms. Nevertheless, some of them remain solid by the end of melting

**1. Introduction** 

of defects are topical questions.

*1Institute of Metal Physics, Ural Division, Russian Academy of Sciences,* 


### **Genesis of Gas Containing Defects in Cast Titanium Parts**

Vladimir Vykhodets1, Tatiana Kurennykh2 and Nataliya Tarenkova2 *1Institute of Metal Physics, Ural Division, Russian Academy of Sciences, 2VSMPO-AVISMA Corporation, Russia* 

#### **1. Introduction**

42 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

[4] Doru. S. (editor); *Metals Handbook Vol. 15,* ASM International, ISBN 0871700212, OHIO

[7] Williams D.B. & Carter C.B., *Transmission Electron Microscopy*, Plenum Press, ISBN 978-

[5] Petzow G.; *Metallographic Etching*, ASM International, ISBN 0871706334, OHIO USA [6] Chase M.W., Davis C.A., Downey J.R., Frurip D.J., McDonald R.A. & Syverud A.N.; *JANAF Thermochemical Tables*, American Chemical Society and American Institute of

[8] *PCPDFWIN Version 2.1*, JCPDS-International Centre for Diffraction Data, 2000.

Physics, ISBN 1-56396-831-2, New York, USA

0306453243, New York, USA

USA

At present, gas-containing inclusions of metallurgical origin like Ti-O, Ti-N, and Ti-O-N are one of the main problems of titanium production. These defects negatively influence the mechanical properties of titanium alloys, as the mechanical properties (hardness and plasticity) of gas-containing inclusions differ appreciably from the corresponding characteristics of alloy matrices. Therefore, cracks nucleate in defect-matrix contact zone upon loading, which results in catastrophic fracture of heavily loaded parts. For example, the tragic consequences of failures of aircraft engine parts owing to the presence of similar effects are well known. The technology of vacuum arc remelting (VAR) of titanium alloys is most vulnerable from the view point of the formation of gas-containing defects, and, at the same time, it dominates in the overall production of titanium alloys. Efforts over many years have not resulted in the development of a technology for the defect-free production of titanium alloys by the VAR method and the titanium defect problem known also as "hard alpha" remains a troublesome aspect of the use of titanium alloys. Attempts to remove or modify the hard-alpha inclusions by diffusion homogenization of ingots had failed due to the long time required for the appropriate heat treatment at high temperature. Therefore, the genesis of defects under the conditions of industrial processes, their identification in each specific case, classification, and the elaboration of measures for reducing the formation of defects are topical questions.

A generally accepted concept has been developed (Bellot & Mitchel, 1994), according to which the sources of defects are particles of titanium and its alloys with a high concentration of oxygen and nitrogen atoms (the predominant element is nitrogen) and liquidus temperatures exceeding the smelting process temperature. Such particles can be formed at different stages of the preparation of batch materials for smelting. It may be supposed that all the components of a batch, such as titanium sponge, titanium return production wastes, and titanium containing alloy additions, are potential sources of defects. Smelting products with rather high concentrations of oxygen and nitrogen atoms may also be sources of defects. In melting oxygen and nitrogen atoms pass into a melt due to diffusion, the melting temperature of particles enriched in light elements decreases, the inclusion size decreases too, and conditions for the their dissolution are created. This scenario takes place for most particles enriched in oxygen and nitrogen atoms. Nevertheless, some of them remain solid by the end of melting

Genesis of Gas Containing Defects in Cast Titanium Parts 45

concentrations of less than 5.0% (hereinafter, all concentrations are given in weight percents). For higher aluminum concentrations, the error was 0.5% and the error in the

The concentrations of oxygen and nitrogen atoms in specimens were determined by nuclear microanalysis technique (NRA) on a 2 MV Van de Graaf accelerator with the use of the 16O(*d*,*p*1)17O and 14N(*d*,α1)12C reactions at an incident beam particle energy of 0.9 MeV. The method in its traditional variant (Vykhodets et al., 1987), in which the diameter of the incident beam of the accelerator was 1.0 or 2.0 mm, was used for the measurement of the average concentrations of light element atoms in a matrix of the titanium alloys. Defects usually had smaller sizes. In this connection, we developed a variant of the NRA method with enhanced locality (Vykhodets et al., 2006) to study them. Using this method, we measured the average concentrations of oxygen and nitrogen in defects 0.1–2.0 mm in size. In this case, each specimen that was measured was equipped with an individual collimator. Its dimension was chosen with consideration for the size of an analyzed zone, i.e., a defect,

measurement of the vanadium concentration was approximately 0.5%.

in a certain specimen. A scheme of NRA experiment is shown in Fig. 1.

Fig. 1. Scheme of NRA experiment: 1 – incident beam; 2 – specimen; 3 – defect to be investigated; 4,6 – silicon surface-barrier detectors of the system of monitoring incident beam; 5 - silicon surface-barrier detector, which register spectrum of the products of nuclear reactions for the studied specimen; 7 – propeller of the system of monitoring incident beam;

Using the NRA method, we measured the concentrations without destruction of a specimen down to a depth of nearly 2.0 μm. This allowed us to distinguish the effects of the volume and surface (uncontrollable) alloying of alloys with the atoms of light elements in the nondestructive profile analysis. These appeared, in particular, at the stage of the preparing

8 - individual collimator; 9 – plate with specimens.

although titanium sponge (plus titanium return production wastes, and alloy additions) are melted twice for many applications and even three times for critical aeroengine parts.

Studies that have been conducted on the defects in titanium alloys may be conditionally classified into two groups: works with model defect sources and model smelting processes and works based on the analysis of the characteristics of defects revealed for industrial processes. At present, the model approach is predominant. However, the patterns established in model experiments may be not adequate for the processes that occur during smelting under industrial conditions, owing to the difference between the characteristics of real and model defect sources, industrial and experimental smelting modes, and also of the corresponding equipment.

In our opinion, the possibilities of the model approaches have been almost exhausted. This is shown by the fact that the model approaches has not led to substantial progress in the understanding of the genesis of defect formation under industrial conditions and the development of technologies for the defect-free production of titanium alloys. In this connection, we place our emphasis on gaining information on the processes of defect formation exclusively on the basis of the data characterizing industrial processes in the present work. In this case, it is impossible to study a statistically representative sample of defect sources, as, on average, a single defect is formed in several millions of particles of batch materials, so the study of a statistically significant ensemble of defects of industrial origin was used as the basis of this approach. This was achieved by the long term monitoring of a process in which the VAR technology was used. We know only one work of a similar type in the literature. This is a report that was published by the Material Development Department of the TRW Company (Grala, 1968). Metallographic and micro X-ray spectral studies of seventeen metallurgical defects in a Ti–6Al–4V alloy were performed in this work. According to the results of this work, the formation of defects may be caused by an increased concentration of atoms of both light elements (oxygen and nitrogen) and aluminum.

#### **2. Materials and methods**

The detection of inclusions in metal and their localization in semi products (rods, billets, plates, and sheets) made from titanium alloys was conducted by the method of ultrasonic testing. After the device received an echo signal indicating the presence of an internal defect we marked and cut out a specimen containing the zone of the signal with increased amplitude, whereupon specimens were subjected to mechanical processing (layer by layer metal stripping) accompanied by metallographic studies. As a result, the zone of the maximum echo signal was brought into the visible plane of a metallographic polished specimen. These polished specimens were the objects of metallographic studies; measurements of the microhardness and the concentration of oxygen, nitrogen, and alloying metal atoms in the matrix and the material of defects were also made.

The measurements of the microhardness in the zone of defects were performed with the use of a device with a Vickers diamond tip. The load of the indentor was 200 g. The concentrations of the atoms of alloying metals (mainly, aluminum and vanadium) in revealed defects were determined by X-ray spectral analysis using the energy dispersion method on an EDAX scanning electron microscope. The statistical error in the measurement of the aluminum concentration was 10% of the measured value for aluminum

although titanium sponge (plus titanium return production wastes, and alloy additions) are

Studies that have been conducted on the defects in titanium alloys may be conditionally classified into two groups: works with model defect sources and model smelting processes and works based on the analysis of the characteristics of defects revealed for industrial processes. At present, the model approach is predominant. However, the patterns established in model experiments may be not adequate for the processes that occur during smelting under industrial conditions, owing to the difference between the characteristics of real and model defect sources, industrial and experimental smelting modes, and also of the

In our opinion, the possibilities of the model approaches have been almost exhausted. This is shown by the fact that the model approaches has not led to substantial progress in the understanding of the genesis of defect formation under industrial conditions and the development of technologies for the defect-free production of titanium alloys. In this connection, we place our emphasis on gaining information on the processes of defect formation exclusively on the basis of the data characterizing industrial processes in the present work. In this case, it is impossible to study a statistically representative sample of defect sources, as, on average, a single defect is formed in several millions of particles of batch materials, so the study of a statistically significant ensemble of defects of industrial origin was used as the basis of this approach. This was achieved by the long term monitoring of a process in which the VAR technology was used. We know only one work of a similar type in the literature. This is a report that was published by the Material Development Department of the TRW Company (Grala, 1968). Metallographic and micro X-ray spectral studies of seventeen metallurgical defects in a Ti–6Al–4V alloy were performed in this work. According to the results of this work, the formation of defects may be caused by an increased concentration of

The detection of inclusions in metal and their localization in semi products (rods, billets, plates, and sheets) made from titanium alloys was conducted by the method of ultrasonic testing. After the device received an echo signal indicating the presence of an internal defect we marked and cut out a specimen containing the zone of the signal with increased amplitude, whereupon specimens were subjected to mechanical processing (layer by layer metal stripping) accompanied by metallographic studies. As a result, the zone of the maximum echo signal was brought into the visible plane of a metallographic polished specimen. These polished specimens were the objects of metallographic studies; measurements of the microhardness and the concentration of oxygen, nitrogen, and alloying

The measurements of the microhardness in the zone of defects were performed with the use of a device with a Vickers diamond tip. The load of the indentor was 200 g. The concentrations of the atoms of alloying metals (mainly, aluminum and vanadium) in revealed defects were determined by X-ray spectral analysis using the energy dispersion method on an EDAX scanning electron microscope. The statistical error in the measurement of the aluminum concentration was 10% of the measured value for aluminum

atoms of both light elements (oxygen and nitrogen) and aluminum.

metal atoms in the matrix and the material of defects were also made.

melted twice for many applications and even three times for critical aeroengine parts.

corresponding equipment.

**2. Materials and methods** 

concentrations of less than 5.0% (hereinafter, all concentrations are given in weight percents). For higher aluminum concentrations, the error was 0.5% and the error in the measurement of the vanadium concentration was approximately 0.5%.

The concentrations of oxygen and nitrogen atoms in specimens were determined by nuclear microanalysis technique (NRA) on a 2 MV Van de Graaf accelerator with the use of the 16O(*d*,*p*1)17O and 14N(*d*,α1)12C reactions at an incident beam particle energy of 0.9 MeV. The method in its traditional variant (Vykhodets et al., 1987), in which the diameter of the incident beam of the accelerator was 1.0 or 2.0 mm, was used for the measurement of the average concentrations of light element atoms in a matrix of the titanium alloys. Defects usually had smaller sizes. In this connection, we developed a variant of the NRA method with enhanced locality (Vykhodets et al., 2006) to study them. Using this method, we measured the average concentrations of oxygen and nitrogen in defects 0.1–2.0 mm in size. In this case, each specimen that was measured was equipped with an individual collimator. Its dimension was chosen with consideration for the size of an analyzed zone, i.e., a defect, in a certain specimen. A scheme of NRA experiment is shown in Fig. 1.

Fig. 1. Scheme of NRA experiment: 1 – incident beam; 2 – specimen; 3 – defect to be investigated; 4,6 – silicon surface-barrier detectors of the system of monitoring incident beam; 5 - silicon surface-barrier detector, which register spectrum of the products of nuclear reactions for the studied specimen; 7 – propeller of the system of monitoring incident beam; 8 - individual collimator; 9 – plate with specimens.

Using the NRA method, we measured the concentrations without destruction of a specimen down to a depth of nearly 2.0 μm. This allowed us to distinguish the effects of the volume and surface (uncontrollable) alloying of alloys with the atoms of light elements in the nondestructive profile analysis. These appeared, in particular, at the stage of the preparing

Genesis of Gas Containing Defects in Cast Titanium Parts 47

defects. Some of the results represented in this work were previously published in (Tarenkova et al., 2006; Vykhodets et al., 2007; Vykhodets et al., 2011), but the corresponding

The above mentioned characteristics of defects (concentrations of gaseous impurities and aluminum, microhardness etc.) are traditional in the description of defects in titanium alloys. In this work, we included also the data on defect coordinates in ingots into the defect database. They are also important for the understanding of the genesis of gas-containing defects since in the industrial VAR technology the rate of dissolution of defect sources in different ingot parts may differ essentially. This is connected with different temperature of the liquid pool in different parts of the ingot and with the use of magnetic stirring in the VAR technology. This operation results in different liquid flow velocity around solid defect sources in the melt. Earlier, in model experiments (Bellot et al., 1997; Reddy, 1990; Schwartz, 1993; Mitchell, 1984) it was found that the dissolution rate of defect sources strongly depends both on the temperature of the liquid pool and the degree of stirring. For example, the measured dissolution rate of nitrogen-containing inclusions for intense stirring was 10 times faster than those without stirring and the dissolution rate doubled for the temperature rise of about 1000С. In our study, the data on the distribution of defects in ingots were obtained by means of the following procedure. Under production conditions, for each metal volume we made a flow chart of technological operations. After detecting the defect in semiproduct or finished product, the flow chart allowed finding the position of the metal volume in the ingot from which the semiproduct or finished product was produced. Further calculations allowed determining two coordinates of a defect in the ingot: distances from a defect to the ingot base and the ingot axis. Accuracy of coordinate determination in the

Histograms of the distributions of defects over the concentrations of nitrogen and oxygen are shown in Figs. 3 and 4. Almost all the revealed defects contained excess concentrations of oxygen and nitrogen in comparison with the volume of titanium alloys; the excess ranged from several times to several tens of times larger. In almost of cases, the concentration of oxygen was slightly higher than that of nitrogen. At the registered concentration level of gaseous impurities, the alloying of inclusions by nitrogen is the predominant factor in the formation of defects. This is associated with the fact that the alloying of titanium by nitrogen leads to an appreciably higher increase in the liquidus temperature of a material in comparison with the alloying of titanium by oxygen. This result, which indicates the predominant role of nitrogen in the formation of gas-containing defects, agrees with the data from study (Wood, 1969). The data obtained on the concentration of gaseous impurities in defects turned out to be insufficient for the verification of the concept, according to which all sources of defects did not dissolve in the liquid pool because of a high concentration of nitrogen and oxygen in them. There are two main reasons for this. Firstly, the range of possible values of smelting temperatures for an industrial process is not known with certainty. Secondly, the concentration of gaseous impurities could not be measured strictly at the end of smelting since during ingot cooling or other high-temperature treatment processes the concentration of light element atoms decreased due to their diffusion from a

data were obtained on an appreciably smaller sample of defects.

ingots was about several millimeters.

**3. Experimental results and discussion** 

**3.1 Oxygen and nitrogen concentration in defects** 

polished specimens and could propagate in a specimen down to a depth of nearly 0.5 μm. The effects of the surface alloying of specimens with atoms of oxygen and nitrogen are excluded from consideration in all the results below. In this work, the sensitivity in the measurements of the concentrations of oxygen and nitrogen was approximately 0.01%, the statistical error in the measurements of concentrations was at a level of several percent of a measured value. The mounting of the individual collimators 8 and the specimens 2 with the defects onto the plate 9 was performed under an optical microscope before the plate was placed into the vacuum chamber of the accelerator unit. As a result, the precision with which the beam hit the defect was 0.01 mm. The typical spectrum of nuclear reactions is shown in Fig. 2.

Fig. 2. The typical spectrum of nuclear reactions.

In total, more than 100 metallurgical defects revealed in the Ti–6Al–4V alloy were studied. This grade dominates in the overall production of titanium alloys. Defects revealed in other grades of titanium alloys were also investigated. In all the cases, the alloys were obtained by the VAR. Inclusions varied in their shapes and sizes and, to a first approximation, it was possible to distinguish spherical and extended defects. The causes of such variations have not been established as yet and this question is not discussed in this work. The minimum size of the studied effect was predominantly within the interval from 0.1 to 2.0 mm. The length of extended defects could amount to several centimeters. The full range of studies was conducted not only for defects, but also for the matrix surrounding a revealed defect. These investigations were performed for verification, and their results did not reveal any particular features. The microstructure, microhardness, and the concentrations of gaseous impurities and metal atoms near defects were identical to those for the zones remote from

polished specimens and could propagate in a specimen down to a depth of nearly 0.5 μm. The effects of the surface alloying of specimens with atoms of oxygen and nitrogen are excluded from consideration in all the results below. In this work, the sensitivity in the measurements of the concentrations of oxygen and nitrogen was approximately 0.01%, the statistical error in the measurements of concentrations was at a level of several percent of a measured value. The mounting of the individual collimators 8 and the specimens 2 with the defects onto the plate 9 was performed under an optical microscope before the plate was placed into the vacuum chamber of the accelerator unit. As a result, the precision with which the beam hit the defect

100 200 300 400 500 700 800

channel number

In total, more than 100 metallurgical defects revealed in the Ti–6Al–4V alloy were studied. This grade dominates in the overall production of titanium alloys. Defects revealed in other grades of titanium alloys were also investigated. In all the cases, the alloys were obtained by the VAR. Inclusions varied in their shapes and sizes and, to a first approximation, it was possible to distinguish spherical and extended defects. The causes of such variations have not been established as yet and this question is not discussed in this work. The minimum size of the studied effect was predominantly within the interval from 0.1 to 2.0 mm. The length of extended defects could amount to several centimeters. The full range of studies was conducted not only for defects, but also for the matrix surrounding a revealed defect. These investigations were performed for verification, and their results did not reveal any particular features. The microstructure, microhardness, and the concentrations of gaseous impurities and metal atoms near defects were identical to those for the zones remote from

**14N(d,<sup>1</sup>**

**) 12C**

**x 5**

**) 12C**

**14N(d,<sup>0</sup>**

was 0.01 mm. The typical spectrum of nuclear reactions is shown in Fig. 2.

**x 0.01**

**12C(d,p)13C**

**14N(d,p)**

**) 17O**

**16O(d,p1**

Fig. 2. The typical spectrum of nuclear reactions.

0

100

200

300

400

500

yield

600

700

800

900

1000

defects. Some of the results represented in this work were previously published in (Tarenkova et al., 2006; Vykhodets et al., 2007; Vykhodets et al., 2011), but the corresponding data were obtained on an appreciably smaller sample of defects.

The above mentioned characteristics of defects (concentrations of gaseous impurities and aluminum, microhardness etc.) are traditional in the description of defects in titanium alloys. In this work, we included also the data on defect coordinates in ingots into the defect database. They are also important for the understanding of the genesis of gas-containing defects since in the industrial VAR technology the rate of dissolution of defect sources in different ingot parts may differ essentially. This is connected with different temperature of the liquid pool in different parts of the ingot and with the use of magnetic stirring in the VAR technology. This operation results in different liquid flow velocity around solid defect sources in the melt. Earlier, in model experiments (Bellot et al., 1997; Reddy, 1990; Schwartz, 1993; Mitchell, 1984) it was found that the dissolution rate of defect sources strongly depends both on the temperature of the liquid pool and the degree of stirring. For example, the measured dissolution rate of nitrogen-containing inclusions for intense stirring was 10 times faster than those without stirring and the dissolution rate doubled for the temperature rise of about 1000С. In our study, the data on the distribution of defects in ingots were obtained by means of the following procedure. Under production conditions, for each metal volume we made a flow chart of technological operations. After detecting the defect in semiproduct or finished product, the flow chart allowed finding the position of the metal volume in the ingot from which the semiproduct or finished product was produced. Further calculations allowed determining two coordinates of a defect in the ingot: distances from a defect to the ingot base and the ingot axis. Accuracy of coordinate determination in the ingots was about several millimeters.

#### **3. Experimental results and discussion**

#### **3.1 Oxygen and nitrogen concentration in defects**

Histograms of the distributions of defects over the concentrations of nitrogen and oxygen are shown in Figs. 3 and 4. Almost all the revealed defects contained excess concentrations of oxygen and nitrogen in comparison with the volume of titanium alloys; the excess ranged from several times to several tens of times larger. In almost of cases, the concentration of oxygen was slightly higher than that of nitrogen. At the registered concentration level of gaseous impurities, the alloying of inclusions by nitrogen is the predominant factor in the formation of defects. This is associated with the fact that the alloying of titanium by nitrogen leads to an appreciably higher increase in the liquidus temperature of a material in comparison with the alloying of titanium by oxygen. This result, which indicates the predominant role of nitrogen in the formation of gas-containing defects, agrees with the data from study (Wood, 1969). The data obtained on the concentration of gaseous impurities in defects turned out to be insufficient for the verification of the concept, according to which all sources of defects did not dissolve in the liquid pool because of a high concentration of nitrogen and oxygen in them. There are two main reasons for this. Firstly, the range of possible values of smelting temperatures for an industrial process is not known with certainty. Secondly, the concentration of gaseous impurities could not be measured strictly at the end of smelting since during ingot cooling or other high-temperature treatment processes the concentration of light element atoms decreased due to their diffusion from a

Genesis of Gas Containing Defects in Cast Titanium Parts 49

was assumed that in Ti-N and Ti-O-N systems nitrogen atoms raise the liquidus temperature of the defect material by the same value if nitrogen concentrations are equal in these systems. The same additive approximation was used for Ti-O and Ti-O-N systems. The effect of vanadium and aluminum on the liquidus temperature of the defect material

1600 1700 1800 1900 2000 2100 2200 2300 2400

*TL ,***°C**

Fig. 5. Distribution of defects over the liquidus temperature of the defect material: *N* is the

The temperature of smelting is usually assumed to be close to 19000С. In study (Bellot et al., 1997), a lower value (17800С) is reported. Anyway, it can be stated that defects were revealed, the material of which has the liquidus temperature both higher and lower than that of the smelting process. This conclusion can hardly change the rough approximations used for the calculation of the liquidus temperature of the defect material. The existence of defects with liquidus temperatures below the temperature of smelting in principle can be

On the whole, in most cases the concentration of nitrogen in defects was higher than in the alloy matrix. For such defects, the data on nitrogen and oxygen concentrations are not in conflict with the existing viewpoint that the sources of defects are titanium particles and its alloys with high concentrations of oxygen and nitrogen. At the same time, prolonged monitoring of the industrial technology revealed about ten defects, in which the concentration of nitrogen did not exceed that in the alloy matrix. Naturally, such defects cannot be considered as gas-containing defects; they turned out to be abnormal also in some

Under the conditions of industrial production, the microhardness is traditionally measured in each case that a defect is revealed. These data are usually used as an

number of defects; *TL* is the calculated liquidus temperature of the defect material.

due to diffusion processes occurring in the system upon completion of smelting.

other respects. We shall dwell on this question in greater detail in section 3.3.

was not considered since it is much weaker than that of nitrogen and oxygen.

0

**3.2 Microhardness** 

2

4

*N*

6

8

Fig. 3. Distribution of defects over the concentration of nitrogen in them: *N* is the number of defects; *CN* is the average concentration of nitrogen in a defect.

Fig. 4. Distribution of defects over the concentration of oxygen in them: *N* is the number of defects; *CO* is the average concentration of oxygen in a defect.

defect to a solid titanium alloy. Fig. 5 presents a calculated histogram of defect distribution on the liquidus temperature of the defect material. The liquidus temperatures *TL* were found using the data of binary diagrams of states for the systems Ti-O and Ti-N (Bellot mitchell, (1994); Fromm & Gebhardt, 1976). Here, an additive approximation was used, namely, it

0 1 2 3 4 5 6 15

Fig. 3. Distribution of defects over the concentration of nitrogen in them: *N* is the number of

*CN*, **%**

0 2 4 6 810

*CO ,* **%**

Fig. 4. Distribution of defects over the concentration of oxygen in them: *N* is the number of

defect to a solid titanium alloy. Fig. 5 presents a calculated histogram of defect distribution on the liquidus temperature of the defect material. The liquidus temperatures *TL* were found using the data of binary diagrams of states for the systems Ti-O and Ti-N (Bellot mitchell, (1994); Fromm & Gebhardt, 1976). Here, an additive approximation was used, namely, it

0

0

5

10

*N*

15

20

defects; *CN* is the average concentration of nitrogen in a defect.

defects; *CO* is the average concentration of oxygen in a defect.

2

4

6

8

*N*

10

12

14

was assumed that in Ti-N and Ti-O-N systems nitrogen atoms raise the liquidus temperature of the defect material by the same value if nitrogen concentrations are equal in these systems. The same additive approximation was used for Ti-O and Ti-O-N systems. The effect of vanadium and aluminum on the liquidus temperature of the defect material was not considered since it is much weaker than that of nitrogen and oxygen.

Fig. 5. Distribution of defects over the liquidus temperature of the defect material: *N* is the number of defects; *TL* is the calculated liquidus temperature of the defect material.

The temperature of smelting is usually assumed to be close to 19000С. In study (Bellot et al., 1997), a lower value (17800С) is reported. Anyway, it can be stated that defects were revealed, the material of which has the liquidus temperature both higher and lower than that of the smelting process. This conclusion can hardly change the rough approximations used for the calculation of the liquidus temperature of the defect material. The existence of defects with liquidus temperatures below the temperature of smelting in principle can be due to diffusion processes occurring in the system upon completion of smelting.

On the whole, in most cases the concentration of nitrogen in defects was higher than in the alloy matrix. For such defects, the data on nitrogen and oxygen concentrations are not in conflict with the existing viewpoint that the sources of defects are titanium particles and its alloys with high concentrations of oxygen and nitrogen. At the same time, prolonged monitoring of the industrial technology revealed about ten defects, in which the concentration of nitrogen did not exceed that in the alloy matrix. Naturally, such defects cannot be considered as gas-containing defects; they turned out to be abnormal also in some other respects. We shall dwell on this question in greater detail in section 3.3.

#### **3.2 Microhardness**

Under the conditions of industrial production, the microhardness is traditionally measured in each case that a defect is revealed. These data are usually used as an

Genesis of Gas Containing Defects in Cast Titanium Parts 51

As different batch materials before smelting and smelted alloys contain various quantities of alloying elements, the patterns of the concentrations of alloying element atoms in defects may be used for the identification of defect sources. A similar idea was suggested in study (Bewlay & Gigliotti, 1997); this approach is widely used in industry. Let us estimate the prospects of this approach for the Ti–6Al–4V alloy. This alloy mainly contains atoms of

0 2 4 6 8 10 12 14

*CAl ,* **%**

Fig. 7. Distribution of defects revealed in the Ti–6Al–4V alloy over the concentration of aluminum: *N* is the number of defects; *CAl* is the aluminum concentration at the center of a

Aluminum and vanadium are also present in the same quantities in titanium production wastes. Titanium sponge does not contain appreciable amounts of alloying elements; only atoms of aluminum (20.0%), oxygen, and titanium are present in the oxygen containing addition alloy before smelting. A much more difficult situation with the concentration of aluminum and vanadium takes place in the smelting products, whose particles can pass into a melt in the process of smelting. According to the data of several reports of the VSMPO– AVISMA Corporation, the concentration of vanadium and aluminum in smelting products obtained in the smelting of the Ti–6Al–4V alloy may be 2.0–4.2 and 4.7–20.0%, respectively. On the whole, it is possible to talk of higher aluminum concentrations in smelting products

The histogram of the distribution of defects over the aluminum concentration at their centers is shown in Fig. 7. The presence of a pronounced peak in the distribution may be

**3.3 Aluminum concentration at the centre of defects** 

in comparison with those in a melted alloy.

defect.

*N*

aluminum (from 5.50 to 6.75%) and vanadium (from 3.50 to 4.50%).

indicator of the presence of light element atoms in defects. In this study, we did not observe any univocal correlation between the microhardness and the concentration of gaseous impurities in defects. This can be seen from Fig. 6, which illustrates the dependence of the microhardness on the total concentration of oxygen and nitrogen in defects. For definiteness, in Fig. 6, we give only the minimum values of the microhardness in the body of a defect because of a slight scatter observed in the measurements of the microhardness.

Fig. 6. Microhardness of defects versus total concentration of nitrogen and oxygen (*CO + CN*) in them.

We also detected several defects with a very high microhardness from ~ 900 to 2600 kg/mm2, the corresponding data are not shown in Fig. 6. These defects proved to be abnormal in several other respects. As in the case of the results on the concentration of nitrogen in defects, this circumstance will be used in this work for the classification of revealed defects. The result shown in Fig. 6 is not trivial. A linear or close to linear dependence of the microhardness on the concentration of gaseous impurities is usually observed (David et al., 1979). From Fig. 6, it follows that in the case of industrial defects, the value of microhardness cannot serve as a criterion of the presence of gaseous impurities in them. On the whole, the only pattern that can be noted with respect to the microhardness is that its value in defects is usually higher that that in the matrix (nearly 350 kg/mm2).

indicator of the presence of light element atoms in defects. In this study, we did not observe any univocal correlation between the microhardness and the concentration of gaseous impurities in defects. This can be seen from Fig. 6, which illustrates the dependence of the microhardness on the total concentration of oxygen and nitrogen in defects. For definiteness, in Fig. 6, we give only the minimum values of the microhardness in the body of a defect because of a slight scatter observed in the measurements of the

0 2 4 6 8 10 12 14

*CO+ CN ,* **%**

Fig. 6. Microhardness of defects versus total concentration of nitrogen and oxygen (*CO + CN*)

We also detected several defects with a very high microhardness from ~ 900 to 2600 kg/mm2, the corresponding data are not shown in Fig. 6. These defects proved to be abnormal in several other respects. As in the case of the results on the concentration of nitrogen in defects, this circumstance will be used in this work for the classification of revealed defects. The result shown in Fig. 6 is not trivial. A linear or close to linear dependence of the microhardness on the concentration of gaseous impurities is usually observed (David et al., 1979). From Fig. 6, it follows that in the case of industrial defects, the value of microhardness cannot serve as a criterion of the presence of gaseous impurities in them. On the whole, the only pattern that can be noted with respect to the microhardness is that its value in defects is usually higher that that in the matrix (nearly

microhardness.

300

400

500

600

microhardness, kg/mm2

in them.

350 kg/mm2).

700

800

900

#### **3.3 Aluminum concentration at the centre of defects**

As different batch materials before smelting and smelted alloys contain various quantities of alloying elements, the patterns of the concentrations of alloying element atoms in defects may be used for the identification of defect sources. A similar idea was suggested in study (Bewlay & Gigliotti, 1997); this approach is widely used in industry. Let us estimate the prospects of this approach for the Ti–6Al–4V alloy. This alloy mainly contains atoms of aluminum (from 5.50 to 6.75%) and vanadium (from 3.50 to 4.50%).

Fig. 7. Distribution of defects revealed in the Ti–6Al–4V alloy over the concentration of aluminum: *N* is the number of defects; *CAl* is the aluminum concentration at the center of a defect.

Aluminum and vanadium are also present in the same quantities in titanium production wastes. Titanium sponge does not contain appreciable amounts of alloying elements; only atoms of aluminum (20.0%), oxygen, and titanium are present in the oxygen containing addition alloy before smelting. A much more difficult situation with the concentration of aluminum and vanadium takes place in the smelting products, whose particles can pass into a melt in the process of smelting. According to the data of several reports of the VSMPO– AVISMA Corporation, the concentration of vanadium and aluminum in smelting products obtained in the smelting of the Ti–6Al–4V alloy may be 2.0–4.2 and 4.7–20.0%, respectively. On the whole, it is possible to talk of higher aluminum concentrations in smelting products in comparison with those in a melted alloy.

The histogram of the distribution of defects over the aluminum concentration at their centers is shown in Fig. 7. The presence of a pronounced peak in the distribution may be

Genesis of Gas Containing Defects in Cast Titanium Parts 53

3

1.0 1.5 2.0 2.5 3.0 3.5 4.0

4

*l ,* mm

0.0 0.1 0.2 0.3 0.4

*Dt / l 2*

Fig. 8. Calculated dependences *A(Dt/l2)* and *A*(*l*) for cylindrical specimens. The dependences *A*(*l*) were calculated at *Dt* of 3.44· 10–5, 6.88· 10–5, 13.8· 10–5, and 27.5· 10–5 mm2 (respectively:

Below, we will use the following notation: *l* is the radius of a cylindrical defect source before smelting, *d* is the diameter of a revealed defect, and *Δ* is the thickness of the defect source

*<sup>d</sup> l= + <sup>Δ</sup>*

*0 0S 0 S*

*2 2*

*1 1 C =C - C -C A 0 - ηΔ + C -C <sup>η</sup><sup>d</sup>*

Hence, according to expression (4), the dependence of the aluminum concentration *CAl* at the center of a revealed defect on its linear size *d* must be linear. It is self evident that this issue concerns the situation where the concentrations *C0* and *CS* are equal for all defect sources. Otherwise, a band of the values of the concentration *CAl* must be observed in experiments. Equation (4) may be used for calculating the thickness *Δ* of the defect source layer dissolved in the process of smelting. No fitted parameters are required to perform the corresponding calculations. As shown by the calculations, the value of *A*(0) does not depend on the parameter *Dt* and is determined only by the shape of a defect source. For example, *A*(0) ≈ 1.49 for a cylindrical specimen and *A*(0) ≈ 1.31 for a plate (*l* was accepted as the radius of a

*<sup>2</sup>* , (3)

(4)

layer dissolved in the process of smelting. Taking into consideration that

*Al*

0.0

1, 2, 3, 4).

we obtain

0.2

0.4

1

2

*A*

0.6

0.8

associated with the fact that all the probable defect sources fall within the corresponding concentration interval. It can be seen that the genesis of only 26 defects can reliably be established on the basis of the measurement results on aluminum concentration. These were defects with an aluminum concentration of less than 4.0%; their origination from titanium sponge does not raise any peculiar doubts. The remaining revealed defects cannot be identified with the use of the data shown in Fig. 7. The vanadium concentration at the center of defects was also measured. As shown by analysis, the data of vanadium do not clarify the situation at all, so they are not given here. Hence, we can state that the approach that is used in practice for the identification of defect sources and based on the measurement of the concentration of aluminum and vanadium in defects has no any substantial grounds. It is not inconceivable that progress in this matter may be made with the use of additional criteria. First of all, we may use the analysis of the dependence of the aluminum concentration at the center of a revealed defect on its linear size.

To analyze the diffusion processes that occur during the smelting of titanium alloys in the subsystem of aluminum atoms we shall use the solution of the second Fick equation for the diffusion from a finite sized body with conjunctive boundaries in the form (Fromm & Gebhardt, 1976)

$$\frac{\text{C} \cdot \text{C}^{0}}{\text{C}^{S} \cdot \text{C}^{0}} = A \left( Dt/l^{2} \right), \tag{1}$$

where *C* is the concentration of an impurity (aluminum in our case) at the center of an inclusion at the end of smelting; *C0* is the initial concentration of an impurity in an inclusion; *CS* is the concentration of an impurity at the interface between an inclusion and melt; *A* is the function of the parameter *<sup>2</sup> Dt l* ; *D*, *t* and *l* are the diffusion coefficient of an impurity, the annealing time and the linear size of a specimen, respectively.

The precise dependence *A(Dt/l2)* for a cylindrical specimen (defect in our case) is shown in Fig. 8 (curve). This dependence also has a similar shape for other regular bodies. Certainly, in the case of defect sources that appear in the smelting of titanium alloys, it is necessary to take into account the fact that their sizes are reduced in the process of smelting. Because of the absence of a precise solution of the second Fick equation in the case of variable dimensions of an inclusion, we shall obtain an approximate solution of the problem for the dependence of the aluminum concentration at the center of a defect, whose dimensions are varied in the process of smelting. The calculated dependences *A*(*l*) obtained at several certain values of the parameter *Dt* are illustrated in Fig. 8 (straight lines). It can be seen that the dependences *A*(*l*) are close to linear ones within the interval of the most appreciable changes, i.e.,

$$A(l) = A(0) \text{- } \eta l \text{ .} \tag{2}$$

The points in one of the straight lines indicate the precise values of the function *A*(*l*). It can be seen that the deviation of the points from a linear dependence is slight. In this case, expression (1) may be applied also for the variable size of a defect with the substitution of *A* for its average value obtained for the entire interval of changes during smelting.

Fig. 8. Calculated dependences *A(Dt/l2)* and *A*(*l*) for cylindrical specimens. The dependences *A*(*l*) were calculated at *Dt* of 3.44· 10–5, 6.88· 10–5, 13.8· 10–5, and 27.5· 10–5 mm2 (respectively: 1, 2, 3, 4).

Below, we will use the following notation: *l* is the radius of a cylindrical defect source before smelting, *d* is the diameter of a revealed defect, and *Δ* is the thickness of the defect source layer dissolved in the process of smelting. Taking into consideration that

$$
\Delta l = \frac{d}{2} + \Delta l\_{\text{\textquotedblleft}g} \tag{3}
$$

we obtain

52 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

associated with the fact that all the probable defect sources fall within the corresponding concentration interval. It can be seen that the genesis of only 26 defects can reliably be established on the basis of the measurement results on aluminum concentration. These were defects with an aluminum concentration of less than 4.0%; their origination from titanium sponge does not raise any peculiar doubts. The remaining revealed defects cannot be identified with the use of the data shown in Fig. 7. The vanadium concentration at the center of defects was also measured. As shown by analysis, the data of vanadium do not clarify the situation at all, so they are not given here. Hence, we can state that the approach that is used in practice for the identification of defect sources and based on the measurement of the concentration of aluminum and vanadium in defects has no any substantial grounds. It is not inconceivable that progress in this matter may be made with the use of additional criteria. First of all, we may use the analysis of the dependence of the aluminum

To analyze the diffusion processes that occur during the smelting of titanium alloys in the subsystem of aluminum atoms we shall use the solution of the second Fick equation for the diffusion from a finite sized body with conjunctive boundaries in the form (Fromm &

*<sup>0</sup>*

where *C* is the concentration of an impurity (aluminum in our case) at the center of an inclusion at the end of smelting; *C0* is the initial concentration of an impurity in an inclusion; *CS* is the concentration of an impurity at the interface between an inclusion and melt; *A* is the function of the parameter *<sup>2</sup> Dt l* ; *D*, *t* and *l* are the diffusion coefficient of an impurity,

The precise dependence *A(Dt/l2)* for a cylindrical specimen (defect in our case) is shown in Fig. 8 (curve). This dependence also has a similar shape for other regular bodies. Certainly, in the case of defect sources that appear in the smelting of titanium alloys, it is necessary to take into account the fact that their sizes are reduced in the process of smelting. Because of the absence of a precise solution of the second Fick equation in the case of variable dimensions of an inclusion, we shall obtain an approximate solution of the problem for the dependence of the aluminum concentration at the center of a defect, whose dimensions are varied in the process of smelting. The calculated dependences *A*(*l*) obtained at several certain values of the parameter *Dt* are illustrated in Fig. 8 (straight lines). It can be seen that the dependences *A*(*l*) are close to linear ones within the interval of the most appreciable

The points in one of the straight lines indicate the precise values of the function *A*(*l*). It can be seen that the deviation of the points from a linear dependence is slight. In this case, expression (1) may be applied also for the variable size of a defect with the substitution of *A*

for its average value obtained for the entire interval of changes during smelting.

*S 0 C -C = A Dt l*

*2*

*C -C* , (1)

*A l =A 0 - ηl* . (2)

concentration at the center of a revealed defect on its linear size.

the annealing time and the linear size of a specimen, respectively.

Gebhardt, 1976)

changes, i.e.,

$$\mathbf{C}\_{Al} = \mathbf{C}^0 \cdot \left(\mathbf{C}^0 \cdot \mathbf{C}^S\right) \left[A\left(\mathbf{O}\right) \cdot \frac{1}{2}\eta\Delta\right] + \frac{1}{2}\left(\mathbf{C}^0 \cdot \mathbf{C}^S\right)\eta d\tag{4}$$

Hence, according to expression (4), the dependence of the aluminum concentration *CAl* at the center of a revealed defect on its linear size *d* must be linear. It is self evident that this issue concerns the situation where the concentrations *C0* and *CS* are equal for all defect sources. Otherwise, a band of the values of the concentration *CAl* must be observed in experiments. Equation (4) may be used for calculating the thickness *Δ* of the defect source layer dissolved in the process of smelting. No fitted parameters are required to perform the corresponding calculations. As shown by the calculations, the value of *A*(0) does not depend on the parameter *Dt* and is determined only by the shape of a defect source. For example, *A*(0) ≈ 1.49 for a cylindrical specimen and *A*(0) ≈ 1.31 for a plate (*l* was accepted as the radius of a cylinder or the half thickness of a plate). In this connection, the results of the calculations for the dissolved layer thickness *Δ* depend on the shape of a defect. This creates certain problems, as defect sources can actually have any shape.

Note that the approach for obtaining the experimental data on the thickness *Δ* of the defect source layer dissolved in the process of smelting and on the size *l* of a defect source before smelting for the industrial technology was first found in this work. Information of this kind is important for the development of methods that reveal the probable defect sources in batch materials. Moreover, the dissolution rate *V* of defect sources in the process of industrial smelting can be estimated in a linear approximation with the use of the value *Δ* as

$$V = \frac{\Delta}{t} \,. \tag{5}$$

Genesis of Gas Containing Defects in Cast Titanium Parts 55

0.0 0.5 1.0 1.5 2.0

*d* **,** mm

Fig. 9. The aluminum concentration *CAl* at the centre of defects that were revealed in the Ti-6Al-4V alloy and originated from titanium sponge as a function of the minimum linear size of the defect: triangles are data for defects with microhardness above 900 kg/mm2; circles

The concentration intervals for defects with ordinary and high microhardness are denoted in Fig. 9 as 1-2 and 3-4, respectively. It may be supposed that these two groups of defect sources differ from each other in the diffusion coefficients *D* of aluminum atoms in them (these coefficients are lower for defects with high microhardness). In turn, the difference between the values of *D* may be caused by different phase states of inclusions with high and

The results displayed in Fig. 9 are of interest mainly for two reasons. Firstly, they showed that the value of microhardness should be considered in the classification of defect sources. Secondly, the accuracy with which the theoretical linear dependence *CAl*(*d*) was fulfilled proved to be high. This occurred in spite of a large number of factors determining the scatter of *CAl*(*d*) data. The width of the concentration intervals denoted in Fig. 9 as 1-2 and 3-4 was 1.25%. This value agrees in particular with the technical specifications for the concentration range of aluminum in the Ti-6Al-4V alloy (from 5.50 to 6.75%). Note however that a single point in Fig. 8 falls outside of the established dependence *CAl*(*d*). The microhardness of this defect was about 800 kg/mm2, *CAl* = 2.3%, and *d* = 0.25 mm. This shows that small and rare

In Fig. 10, the data on *CAl*(*d*) are shown for almost all the defects revealed in the Ti–6Al–4V alloy. Only a single defect is not represented in this figure. Its minimum size (3.0 mm)

deviations from the theoretical dependence *CAl*(*d*) can nevertheless take place.

4

3

0.0

ordinary microhardness values.

0.5

1.0

1.5

1

2.0

*CAl ,* **%**

2

are data for defects with microhardness below 800 kg/mm2.

2.5

3.0

3.5

4.0

Previously, the defect dissolution rate, as mentioned in section 2, was determined only in model experiment (Bewlay & Gigliotti, 1997; Bellot et al., 1997; Reddy, 1990; Schwartz, 1993; Mitchel, 1984), and its strong dependence on the temperature of the liquid pool and the liquid flow velocity around solid particle was established.

Since reliable identification of the sources of 26 defects proved to be feasible, the additional identification method based on the analysis of the linear dependence of the aluminum concentration *CAl* at the center of a defect on its size *d* (expression (4)) might be approved on the same defects. Actually, linear dependences (4) are expected to be fulfilled only qualitatively. This is connected with the following circumstances. To obtain such dependences, it was necessary to clarify the notion "linear size of a defect," which is especially topical for extended defects. In specimens with complex shapes, the rate of diffusion processes is generally determined by the minimum linear size of a specimen, so it is chosen as the parameter *d*. Besides, the specificity of experiments performed in this work consisted in the fact that the concentrations of light elements and aluminum could not be measured at the end of smelting. During the cooling of ingots and other high temperature treatment processes, the atoms of impurities diffuse from defects into the solid titanium alloy and on the contrary, thus leading to a decrease or increase in the concentration of impurities in defects. In section 3.1 it was pointed out that the concentrations of oxygen and nitrogen in the revealed defects were apparently much smaller than those at the end of smelting. However, for aluminum this effect is expected to be appreciably weaker or negligible since the diffusion coefficients of aluminum in solid titanium are almost five orders of magnitude smaller than those of oxygen and nitrogen (Le Claire & Neumann, 1990; Le Claire, 1990). Besides, there are other reasons for the scatter of experimental data on *CAl*(*d*). In particular, parameters *A*(0), *CS*, and *C0* in expression (4) will vary within certain intervals for different defects and different smelting processes. For example, as already noted above, the concentration of aluminum in the melt varied in different smelting processes from 5.50 to 6.75%.

The experimental data on *CAl*(*d*) for these 26 defects are illustrated in Fig. 9. It can be seen that there is no pattern in the corresponding data. The situation changes radically if defects with ordinary (below 800 kg/mm2) and high (above 900 kg/mm2) microhardness values are considered separately. As seen from Fig. 9, a linear dependence *CAl*(*d*), which results from the diffusion model, is observed for each individual group of defects in this case.

cylinder or the half thickness of a plate). In this connection, the results of the calculations for the dissolved layer thickness *Δ* depend on the shape of a defect. This creates certain

Note that the approach for obtaining the experimental data on the thickness *Δ* of the defect source layer dissolved in the process of smelting and on the size *l* of a defect source before smelting for the industrial technology was first found in this work. Information of this kind is important for the development of methods that reveal the probable defect sources in batch materials. Moreover, the dissolution rate *V* of defect sources in the process of industrial

*<sup>Δ</sup> V =*

Previously, the defect dissolution rate, as mentioned in section 2, was determined only in model experiment (Bewlay & Gigliotti, 1997; Bellot et al., 1997; Reddy, 1990; Schwartz, 1993; Mitchel, 1984), and its strong dependence on the temperature of the liquid pool and the

Since reliable identification of the sources of 26 defects proved to be feasible, the additional identification method based on the analysis of the linear dependence of the aluminum concentration *CAl* at the center of a defect on its size *d* (expression (4)) might be approved on the same defects. Actually, linear dependences (4) are expected to be fulfilled only qualitatively. This is connected with the following circumstances. To obtain such dependences, it was necessary to clarify the notion "linear size of a defect," which is especially topical for extended defects. In specimens with complex shapes, the rate of diffusion processes is generally determined by the minimum linear size of a specimen, so it is chosen as the parameter *d*. Besides, the specificity of experiments performed in this work consisted in the fact that the concentrations of light elements and aluminum could not be measured at the end of smelting. During the cooling of ingots and other high temperature treatment processes, the atoms of impurities diffuse from defects into the solid titanium alloy and on the contrary, thus leading to a decrease or increase in the concentration of impurities in defects. In section 3.1 it was pointed out that the concentrations of oxygen and nitrogen in the revealed defects were apparently much smaller than those at the end of smelting. However, for aluminum this effect is expected to be appreciably weaker or negligible since the diffusion coefficients of aluminum in solid titanium are almost five orders of magnitude smaller than those of oxygen and nitrogen (Le Claire & Neumann, 1990; Le Claire, 1990). Besides, there are other reasons for the scatter of experimental data on *CAl*(*d*). In particular, parameters *A*(0), *CS*, and *C0* in expression (4) will vary within certain intervals for different defects and different smelting processes. For example, as already noted above, the concentration of aluminum in the melt varied in different smelting

The experimental data on *CAl*(*d*) for these 26 defects are illustrated in Fig. 9. It can be seen that there is no pattern in the corresponding data. The situation changes radically if defects with ordinary (below 800 kg/mm2) and high (above 900 kg/mm2) microhardness values are considered separately. As seen from Fig. 9, a linear dependence *CAl*(*d*), which results from

the diffusion model, is observed for each individual group of defects in this case.

*<sup>t</sup>* . (5)

smelting can be estimated in a linear approximation with the use of the value *Δ* as

problems, as defect sources can actually have any shape.

liquid flow velocity around solid particle was established.

processes from 5.50 to 6.75%.

Fig. 9. The aluminum concentration *CAl* at the centre of defects that were revealed in the Ti-6Al-4V alloy and originated from titanium sponge as a function of the minimum linear size of the defect: triangles are data for defects with microhardness above 900 kg/mm2; circles are data for defects with microhardness below 800 kg/mm2.

The concentration intervals for defects with ordinary and high microhardness are denoted in Fig. 9 as 1-2 and 3-4, respectively. It may be supposed that these two groups of defect sources differ from each other in the diffusion coefficients *D* of aluminum atoms in them (these coefficients are lower for defects with high microhardness). In turn, the difference between the values of *D* may be caused by different phase states of inclusions with high and ordinary microhardness values.

The results displayed in Fig. 9 are of interest mainly for two reasons. Firstly, they showed that the value of microhardness should be considered in the classification of defect sources. Secondly, the accuracy with which the theoretical linear dependence *CAl*(*d*) was fulfilled proved to be high. This occurred in spite of a large number of factors determining the scatter of *CAl*(*d*) data. The width of the concentration intervals denoted in Fig. 9 as 1-2 and 3-4 was 1.25%. This value agrees in particular with the technical specifications for the concentration range of aluminum in the Ti-6Al-4V alloy (from 5.50 to 6.75%). Note however that a single point in Fig. 8 falls outside of the established dependence *CAl*(*d*). The microhardness of this defect was about 800 kg/mm2, *CAl* = 2.3%, and *d* = 0.25 mm. This shows that small and rare deviations from the theoretical dependence *CAl*(*d*) can nevertheless take place.

In Fig. 10, the data on *CAl*(*d*) are shown for almost all the defects revealed in the Ti–6Al–4V alloy. Only a single defect is not represented in this figure. Its minimum size (3.0 mm)

Genesis of Gas Containing Defects in Cast Titanium Parts 57

7

6

5

4

3

<sup>12</sup> 8

0.0 0.5 1.0 1.5 2.0

Fig. 10. The aluminum concentration *CAl* at the centre of defects that were revealed in the Ti-6Al-4V alloy as a function of the minimum linear size of the defect: triangles are data for

So, a method for the identification of the sources of defects in the industrial VAR technology was developed in this work. Diffusion laws form the basis of this method. This scheme has proven to be very simple for practical application. The identification of the source of a defect revealed in the smelting of the Ti–6Al–4V alloy requires a very limited data set containing the aluminum concentration at the center of the defect, its minimum size, the microhardness of the material of the defect, and the average nitrogen concentration in the body of the defect. Under industrial conditions, this information is usually available, so this method can be recommended for practical application. Note also that during the prolonged monitoring of an industrial process we did not observe defects whose data contradict this method.

Overall, the identification proved to be unfeasible for approximately 20% of the defects. According to the results of the identification, return production wastes were the source of approximately half of all the revealed defects, and nearly a quarter of them originated from titanium sponge. The remaining defects were formed from the particles of addition alloy

Further, using expressions (3)–(5) and the data shown in Fig. 10, the thickness of the defect source layer dissolved in the process of smelting and the dissolution rate *V* of defect sources were estimated. The average value of the thickness Δ was about 2 – 3 mm. At the same time,

defects with microhardness above 900 kg/mm2; circles are data for defects with microhardness below 800 kg/mm2; asterisks are data for defects in which the nitrogen

*d* **,** mm

0

and smelting products.

2

1

2

concentration does not exceed that in the matrix.

4

6

8

*CAl* **, %**

10

appreciably exceeded the sizes of the other defects, so it was unreasonable to provide the data on this defect. In connection with this exclusion, it should be noted that this defect was not inconsistent with the common patterns, which will be considered below. From Fig. 10, it can be seen that almost all the defects fit into the determined intervals of aluminum concentrations. In Fig. 10, they are denoted as 1–2, 3–4, 5–6, and 7–8. The first two intervals have already been discussed above. Interval 7–8, as well as intervals 1–2 and 3–4, correspond to the slanting linear dependence *CAl*(*d*). It makes sense to relate this to the sources of defects, in which the concentration of aluminum before smelting was higher in comparison with that in a melt. They might be the particles of smelting products and addition alloy. Interval 5–6 is horizontal. According to the diffusion model, it corresponds to the sources of defects in which the concentration of aluminum before smelting was close to that in the liquid pool. These may be particles of smelting products and return production wastes. The width of the concentration intervals was equal for all the corridors (1–2, 3–4, 5– 6, and 7–8).

Let us discuss the sense of these results in connection with the problem of the identification of defect sources. First, we consider the case of small defects. From Fig. 10, it can be seen that such defects can be equally related to the intervals denoted as 3–4, 5–6, and 7–8. This means that in the case of small defects, their identification with the use of the discussed scheme is impossible, i.e., these defects may be formed from the particles of titanium sponge, return production wastes, smelting products, and addition alloy.

Further, let us consider medium and large sized defects. Here, the situation looks to be much more unambiguous. First, it is possible to note that only the defect sources that are characterized by a nearly equal and high aluminum concentration before smelting fitted in the interval 7–8*.* Such a condition is satisfied by the particles of addition alloy with an aluminum concentration of 20.0%. Hence, we can state that the defects that originate from the addition alloy are presented within interval 7–8. The absence of defects that formed from smelting products within the interval 7–8 and their presence within interval 5–6 seems to be unnatural. This allows us to think that defects that formed from smelting products are not represented within the interval 5–6, i.e., in the case of large defects, the 5–6 interval is caused by the formation of defects from return production wastes.

Let us consider the defects that did not fit in any of the intervals distinguished in Fig. 10. Only four similar defects were found. One of them had microhardness higher than 900 kg/mm2. The abnormal behavior of such defects has already been mentioned. The other three defects, whose data are denoted by asterisks in Fig. 10, have nitrogen concentrations that are lower than that in the matrix. It was previously mentioned that such inclusions cannot be reliably classified as gas-containing defects. We suppose that these defects were formed from the particles of smelting products and that their residence in a liquid pool might be very short, for which reason they fall outside of the diffusion patterns. Another defect, which is within the interval 5–6 in Fig. 10, should be classified with inclusions of the same type. The results for this defect are denoted by an asterisk. Consequently, it can be seen that all the cases in which the data on defects lie outside the distinguished intervals 1– 2, 3–4, 5–6, and 7–8, are accompanied by abnormal properties of the defects (very high microhardness or very low nitrogen concentration).

appreciably exceeded the sizes of the other defects, so it was unreasonable to provide the data on this defect. In connection with this exclusion, it should be noted that this defect was not inconsistent with the common patterns, which will be considered below. From Fig. 10, it can be seen that almost all the defects fit into the determined intervals of aluminum concentrations. In Fig. 10, they are denoted as 1–2, 3–4, 5–6, and 7–8. The first two intervals have already been discussed above. Interval 7–8, as well as intervals 1–2 and 3–4, correspond to the slanting linear dependence *CAl*(*d*). It makes sense to relate this to the sources of defects, in which the concentration of aluminum before smelting was higher in comparison with that in a melt. They might be the particles of smelting products and addition alloy. Interval 5–6 is horizontal. According to the diffusion model, it corresponds to the sources of defects in which the concentration of aluminum before smelting was close to that in the liquid pool. These may be particles of smelting products and return production wastes. The width of the concentration intervals was equal for all the corridors (1–2, 3–4, 5–

Let us discuss the sense of these results in connection with the problem of the identification of defect sources. First, we consider the case of small defects. From Fig. 10, it can be seen that such defects can be equally related to the intervals denoted as 3–4, 5–6, and 7–8. This means that in the case of small defects, their identification with the use of the discussed scheme is impossible, i.e., these defects may be formed from the particles of titanium sponge, return

Further, let us consider medium and large sized defects. Here, the situation looks to be much more unambiguous. First, it is possible to note that only the defect sources that are characterized by a nearly equal and high aluminum concentration before smelting fitted in the interval 7–8*.* Such a condition is satisfied by the particles of addition alloy with an aluminum concentration of 20.0%. Hence, we can state that the defects that originate from the addition alloy are presented within interval 7–8. The absence of defects that formed from smelting products within the interval 7–8 and their presence within interval 5–6 seems to be unnatural. This allows us to think that defects that formed from smelting products are not represented within the interval 5–6, i.e., in the case of large defects, the 5–6 interval is caused

Let us consider the defects that did not fit in any of the intervals distinguished in Fig. 10. Only four similar defects were found. One of them had microhardness higher than 900 kg/mm2. The abnormal behavior of such defects has already been mentioned. The other three defects, whose data are denoted by asterisks in Fig. 10, have nitrogen concentrations that are lower than that in the matrix. It was previously mentioned that such inclusions cannot be reliably classified as gas-containing defects. We suppose that these defects were formed from the particles of smelting products and that their residence in a liquid pool might be very short, for which reason they fall outside of the diffusion patterns. Another defect, which is within the interval 5–6 in Fig. 10, should be classified with inclusions of the same type. The results for this defect are denoted by an asterisk. Consequently, it can be seen that all the cases in which the data on defects lie outside the distinguished intervals 1– 2, 3–4, 5–6, and 7–8, are accompanied by abnormal properties of the defects (very high

production wastes, smelting products, and addition alloy.

by the formation of defects from return production wastes.

microhardness or very low nitrogen concentration).

6, and 7–8).

Fig. 10. The aluminum concentration *CAl* at the centre of defects that were revealed in the Ti-6Al-4V alloy as a function of the minimum linear size of the defect: triangles are data for defects with microhardness above 900 kg/mm2; circles are data for defects with microhardness below 800 kg/mm2; asterisks are data for defects in which the nitrogen concentration does not exceed that in the matrix.

So, a method for the identification of the sources of defects in the industrial VAR technology was developed in this work. Diffusion laws form the basis of this method. This scheme has proven to be very simple for practical application. The identification of the source of a defect revealed in the smelting of the Ti–6Al–4V alloy requires a very limited data set containing the aluminum concentration at the center of the defect, its minimum size, the microhardness of the material of the defect, and the average nitrogen concentration in the body of the defect. Under industrial conditions, this information is usually available, so this method can be recommended for practical application. Note also that during the prolonged monitoring of an industrial process we did not observe defects whose data contradict this method.

Overall, the identification proved to be unfeasible for approximately 20% of the defects. According to the results of the identification, return production wastes were the source of approximately half of all the revealed defects, and nearly a quarter of them originated from titanium sponge. The remaining defects were formed from the particles of addition alloy and smelting products.

Further, using expressions (3)–(5) and the data shown in Fig. 10, the thickness of the defect source layer dissolved in the process of smelting and the dissolution rate *V* of defect sources were estimated. The average value of the thickness Δ was about 2 – 3 mm. At the same time,

Genesis of Gas Containing Defects in Cast Titanium Parts 59

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0

*h/H*

Fig. 11. Distribution of defects over their non-dimensional coordinate *h/H*: *N* is the number

For those cases when two defects were revealed in an ingot simultaneously, we analyzed the differences in *h* and *r* coordinates of these defects. In the majority of cases, such defects were very close to each other (the difference in *h* and *r* coordinates usually did not exceed several millimeters). Only in one case the difference in the *h* coordinates was very large (more than 600 mm). Based on this result we can assert that the sources of defects usually do not move

From Figs. 11 and 14 it is seen that the distribution of defects over the ingot height and radius is strongly inhomogeneous. The defects are located chiefly near the ingot base and axis. Since the sources of defects do not move over large distances during smelting as mentioned above, it may be stated that the dissolution rate of defect sources near the ingot base and the axis is much smaller than the average dissolution rate in the ingot. This may be due to a decreased temperature of the liquid pool and (or) a decreased liquid flow velocity around solid particle in the melt. The ingot volumes with *h/H* values from 0 to 0.2 and *r/R* from 0 to 0.1 turned out to be most critical for defect formation. In this range of *h/H* values (20% of the metal volume) 68% of defects were revealed, while in the above range of *r/R* values (1% of the metal volume) there were revealed 17% of defects. These findings are indicative of strong inhomogeneity of defect distribution in the ingots. The number of defects in the unit volume of the metal for the volume bounded by *h/H* values from 0 to 0.2 and *r/R* values from 0 to 0.1 (0.2% of the ingot volume) turned out to be almost 60 times

Relying on these findings we can propose two mechanisms of defect formation. One of them suggests fluctuations in the characteristics of defect sources, i.e. the sources of defects

of defects; *h* is the distance from a defect to the ingot base; *H* is the ingot height.

0

over large distances during smelting.

higher than the average value in the ingots.

20

40

*N*

60

80

the data obtained do not exclude a broad interval of Δ values during the industrial process: from 0 to 8.5 mm. With the corresponding estimates, the shape of the defects and the concentration of aluminum in the liquid pool varied from 5.50 to 6.75%. Besides, it was taken into consideration that the concentration of aluminum *CS* at the boundary between the defect source and the melt can be smaller than the corresponding equilibrium concentration. With this approach, the interval of Δ values was obviously overestimated. In reality, it will be smaller.

The average value of the dissolution rate *V* was approximately 1.0 μm/s. The estimation was performed under the assumption that the residence time of any melt volume in the liquid phase was 2500 s. This value of the average dissolution rate is very low. Model experiments (Bellot et al., 1997) showed that the dissolution rate of the nitrogen-containing source of defects is 2.2, 4.0, and 6.7 μm/s at the temperature of the liquid pool of 1700, 1800, and 19000С, respectively, when the liquid flow velocity around solid particle in the liquid pool was 1 cm/s. Even when the liquid flow velocity around solid particle was 0.1 cm/s and the temperature was 18000С, it was 2.0 μm/s. One of the possible directions for defect formation reduction is enhancement of the dissolution rate of defect sources. The prospects of this trend will be discussed in section 3.4.

#### **3.4 Location of defects in ingots**

The experimental data on the location of defects in ingots were obtained for the ingots from 650 to 1000 mm in diameter and from 1200 to 2750 mm in height. Hereafter we use the following non-dimensional coordinates of defects: *r/R* (*r* is the distance from a defect to the ingot axis, *R* is the ingot radius) and *h/H* (*h* is the distance from a defect to the ingot base, *H*  is the ingot height). The subjects of investigation were the distributions of defects over *r/R* and *h/H* coordinates, as well as the probability of formation of one, two, and more defects in an ingot. The experimental data on the location of defects in the ingots are presented in Figs. 11, 12 and 13. In Fig. 14, the number of defects revealed in the unit of the ingot volume is shown as a function of the non-dimensional radius *r/R*. The data for Fig. 14 were obtained by calculation from primarily data displayed in Fig. 13.

In a real process, the defect formation probability is *p* = *N*/*Nh* « 1, where *N* and *Nh* are the numbers of defects and heats, respectively. If the formation of a defect is an independent event the probability of formation of several defects in one ingot *p*(*n*) is

$$p(n) \;= p^n,\tag{6}$$

where *n* is the number of defects in the ingot. Since the formation of a defect is a very rare event (*p* « 1), the formation of two or more defects in one ingot in this mechanism is almost impossible. However, experimental results indicated that one, two and even more defects can exist in ingots at the close probability.

Altogether we have discovered 205 defects in 86 ingots. 42 ingots contained one defect, 27 ingots contained two defects, 7 ingots had three defects, and in 10 ingots there were more than three defects. These results show that generally the formation of each defect cannot be considered as an independent event .

the data obtained do not exclude a broad interval of Δ values during the industrial process: from 0 to 8.5 mm. With the corresponding estimates, the shape of the defects and the concentration of aluminum in the liquid pool varied from 5.50 to 6.75%. Besides, it was taken into consideration that the concentration of aluminum *CS* at the boundary between the defect source and the melt can be smaller than the corresponding equilibrium concentration. With this approach, the interval of Δ values was obviously overestimated. In reality, it will

The average value of the dissolution rate *V* was approximately 1.0 μm/s. The estimation was performed under the assumption that the residence time of any melt volume in the liquid phase was 2500 s. This value of the average dissolution rate is very low. Model experiments (Bellot et al., 1997) showed that the dissolution rate of the nitrogen-containing source of defects is 2.2, 4.0, and 6.7 μm/s at the temperature of the liquid pool of 1700, 1800, and 19000С, respectively, when the liquid flow velocity around solid particle in the liquid pool was 1 cm/s. Even when the liquid flow velocity around solid particle was 0.1 cm/s and the temperature was 18000С, it was 2.0 μm/s. One of the possible directions for defect formation reduction is enhancement of the dissolution rate of defect sources. The prospects

The experimental data on the location of defects in ingots were obtained for the ingots from 650 to 1000 mm in diameter and from 1200 to 2750 mm in height. Hereafter we use the following non-dimensional coordinates of defects: *r/R* (*r* is the distance from a defect to the ingot axis, *R* is the ingot radius) and *h/H* (*h* is the distance from a defect to the ingot base, *H*  is the ingot height). The subjects of investigation were the distributions of defects over *r/R* and *h/H* coordinates, as well as the probability of formation of one, two, and more defects in an ingot. The experimental data on the location of defects in the ingots are presented in Figs. 11, 12 and 13. In Fig. 14, the number of defects revealed in the unit of the ingot volume is shown as a function of the non-dimensional radius *r/R*. The data for Fig. 14 were obtained

In a real process, the defect formation probability is *p* = *N*/*Nh* « 1, where *N* and *Nh* are the numbers of defects and heats, respectively. If the formation of a defect is an independent

where *n* is the number of defects in the ingot. Since the formation of a defect is a very rare event (*p* « 1), the formation of two or more defects in one ingot in this mechanism is almost impossible. However, experimental results indicated that one, two and even more defects

Altogether we have discovered 205 defects in 86 ingots. 42 ingots contained one defect, 27 ingots contained two defects, 7 ingots had three defects, and in 10 ingots there were more than three defects. These results show that generally the formation of each defect cannot be

*<sup>n</sup> p n p* , (6)

be smaller.

of this trend will be discussed in section 3.4.

by calculation from primarily data displayed in Fig. 13.

can exist in ingots at the close probability.

considered as an independent event .

event the probability of formation of several defects in one ingot *p*(*n*) is

**3.4 Location of defects in ingots** 

Fig. 11. Distribution of defects over their non-dimensional coordinate *h/H*: *N* is the number of defects; *h* is the distance from a defect to the ingot base; *H* is the ingot height.

For those cases when two defects were revealed in an ingot simultaneously, we analyzed the differences in *h* and *r* coordinates of these defects. In the majority of cases, such defects were very close to each other (the difference in *h* and *r* coordinates usually did not exceed several millimeters). Only in one case the difference in the *h* coordinates was very large (more than 600 mm). Based on this result we can assert that the sources of defects usually do not move over large distances during smelting.

From Figs. 11 and 14 it is seen that the distribution of defects over the ingot height and radius is strongly inhomogeneous. The defects are located chiefly near the ingot base and axis. Since the sources of defects do not move over large distances during smelting as mentioned above, it may be stated that the dissolution rate of defect sources near the ingot base and the axis is much smaller than the average dissolution rate in the ingot. This may be due to a decreased temperature of the liquid pool and (or) a decreased liquid flow velocity around solid particle in the melt. The ingot volumes with *h/H* values from 0 to 0.2 and *r/R* from 0 to 0.1 turned out to be most critical for defect formation. In this range of *h/H* values (20% of the metal volume) 68% of defects were revealed, while in the above range of *r/R* values (1% of the metal volume) there were revealed 17% of defects. These findings are indicative of strong inhomogeneity of defect distribution in the ingots. The number of defects in the unit volume of the metal for the volume bounded by *h/H* values from 0 to 0.2 and *r/R* values from 0 to 0.1 (0.2% of the ingot volume) turned out to be almost 60 times higher than the average value in the ingots.

Relying on these findings we can propose two mechanisms of defect formation. One of them suggests fluctuations in the characteristics of defect sources, i.e. the sources of defects

Genesis of Gas Containing Defects in Cast Titanium Parts 61

with the model experiments (Bellot et al., 1997) on the examination of the dissolution rate of gas-containing inclusions as a function of liquid flow velocity around solid particles in the melt. In the existing technology, the rotation axis of the magnetic field coincides with the axis of the ingot during the whole period of smelting. This is likely to be one of the reasons why more defects are formed near the ingot axis than at the ingot periphery. Therefore, a possible measure for defect formation reduction is variation of the position of the magnetic field rotation axis during smelting. The estimates showed that with the use of this technique the defect formation can be reduced not more than in 2.5 times. In the estimation we used the data of Fig. 13 and Fig. 14. Besides, the following concepts were

First, different liquid flow velocities around solid particles in the liquid pool was considered to be the only reason for inhomogeneous distribution of defects over the ingot radius. Second, it was postulated that variation of the position of the magnetic field rotation axis can provide a high liquid flow velocity around solid particles in all parts of

0.0 0.2 0.4 0.6 0.8 1.0

*r/R*

Fig. 13. Distribution of defects over their non-dimensional coordinate *r/R*: *N* is the number of

defects; *r* is the distance from a defect to the ingot axis; *R* is the ingot radius.

accepted.

the ingot.

0

5

10

15

*N*

20

25

30

Fig. 12. Distribution of defects over their coordinate *h/H* near the ingot base: *N* is the number of defects; *h* is the distance from a defect to the ingot base; *H* is the ingot height.

contained in the charge materials have unfavorable characteristics for complete dissolution of defect sources during smelting. Such characteristics can include large sizes of defect sources and/or high concentrations of gaseous impurities in them. According to this mechanism two or more defects could form in one ingot and in one place of the ingot because some defect sources fell into two, three or more fragments during smelting. For this mechanism it is difficult to estimate the probabilities *p*(*n*) theoretically.

The second mechanism suggests fluctuations in the dissolution rate of defect sources during smelting. Such fluctuations can be associated with the existence of microvolumes with a decreased temperature of the liquid pool and/or microvolumes with a decreased liquid flow velocity around solid particles in the melt during smelting. According to this mechanism the formation of several defects in one ingot and in one place of the ingot is due to the value of fluctuation for the dissolution rate of defect sources. For this mechanism it is also difficult to determine theoretically the probabilities *p*(*n*). Thus, there are no available data at present to give preference to one of the considered fluctuation mechanisms. At the same time, this question is a topical problem. Its solution determines the direction of the main efforts for reducing defect formation in the VAR technology. The first mechanism suggests scientific and technological measures for enhancing the quality of charge materials. With the second mechanism, efforts should be focused on the stabilization of the smelting regime. The results obtained on the distribution of defects over the ingot height (Fig. 11) agree with the reasoning (Bellot & Mitchel, 1994) that the temperature of the liquid pool is minimal near the ingot base. Probably, this cause for enhanced defect formation in the VAR technology cannot be eliminated. The results obtained on the distribution of defects over the ingot radius agree

0.00 0.02 0.04 0.06 0.08 0.10 0.12

*h/H*

Fig. 12. Distribution of defects over their coordinate *h/H* near the ingot base: *N* is the number

contained in the charge materials have unfavorable characteristics for complete dissolution of defect sources during smelting. Such characteristics can include large sizes of defect sources and/or high concentrations of gaseous impurities in them. According to this mechanism two or more defects could form in one ingot and in one place of the ingot because some defect sources fell into two, three or more fragments during smelting. For this

The second mechanism suggests fluctuations in the dissolution rate of defect sources during smelting. Such fluctuations can be associated with the existence of microvolumes with a decreased temperature of the liquid pool and/or microvolumes with a decreased liquid flow velocity around solid particles in the melt during smelting. According to this mechanism the formation of several defects in one ingot and in one place of the ingot is due to the value of fluctuation for the dissolution rate of defect sources. For this mechanism it is also difficult to determine theoretically the probabilities *p*(*n*). Thus, there are no available data at present to give preference to one of the considered fluctuation mechanisms. At the same time, this question is a topical problem. Its solution determines the direction of the main efforts for reducing defect formation in the VAR technology. The first mechanism suggests scientific and technological measures for enhancing the quality of charge materials. With the second mechanism, efforts should be focused on the stabilization of the smelting regime. The results obtained on the distribution of defects over the ingot height (Fig. 11) agree with the reasoning (Bellot & Mitchel, 1994) that the temperature of the liquid pool is minimal near the ingot base. Probably, this cause for enhanced defect formation in the VAR technology cannot be eliminated. The results obtained on the distribution of defects over the ingot radius agree

of defects; *h* is the distance from a defect to the ingot base; *H* is the ingot height.

mechanism it is difficult to estimate the probabilities *p*(*n*) theoretically.

0

2

4

6

8

*N*

10

12

14

with the model experiments (Bellot et al., 1997) on the examination of the dissolution rate of gas-containing inclusions as a function of liquid flow velocity around solid particles in the melt. In the existing technology, the rotation axis of the magnetic field coincides with the axis of the ingot during the whole period of smelting. This is likely to be one of the reasons why more defects are formed near the ingot axis than at the ingot periphery. Therefore, a possible measure for defect formation reduction is variation of the position of the magnetic field rotation axis during smelting. The estimates showed that with the use of this technique the defect formation can be reduced not more than in 2.5 times. In the estimation we used the data of Fig. 13 and Fig. 14. Besides, the following concepts were accepted.

First, different liquid flow velocities around solid particles in the liquid pool was considered to be the only reason for inhomogeneous distribution of defects over the ingot radius. Second, it was postulated that variation of the position of the magnetic field rotation axis can provide a high liquid flow velocity around solid particles in all parts of the ingot.

Fig. 13. Distribution of defects over their non-dimensional coordinate *r/R*: *N* is the number of defects; *r* is the distance from a defect to the ingot axis; *R* is the ingot radius.

Genesis of Gas Containing Defects in Cast Titanium Parts 63

The distribution of defects over the height and radius of ingots was studied; it turned out to be strongly inhomogeneous. The defects were located predominately near the ingot base and the ingot axis. The number of defects in the unit volume of the metal in the lower part and near the axis of the ingots was almost 60 times greater than the average value in the ingots. The inhomogeneous distribution of defects in the ingots is due to different dissolution rate of defect sources during smelting in different parts of the ingot. In turn, this results from a low temperature of the liquid pool in the lower part of the ingots and a low

A large number of cases have been registered when not one but two, three or even more defects are formed simultaneously in the ingot. The probabilities of formation of one and several defects were comparable. Usually, several defects in the ingot were located very close to each other. These findings indicate that generally the formation of each defect

A method for identification of the sources of metallurgical defects in the smelting of titanium alloys has been developed. It is based on the diffusion mechanism for the modification of defect sources in smelting and the established dependences of the aluminum concentration at the center of a defect on its size. The identification scheme included data on microhardness and the average concentration of gaseous impurities in the zone of a defect.

This work was supported by the project No 11-2-06-AVI of the Basic Research Program of the Ural Division of Russian Academy of Sciences "Identification of stages which are responsible for formation of gas-containing metallurgical defects for the industrial VAR

Bellot, J.P. Mitchell, A. (1994). Hard-alpha particle behaviour in a titanium alloy liquid

Bellot, J.P. et al. (1997). Dissolution of Hard-Alpha Inclusions in Liquid Titanium Alloys,

Bewlay, B.P. & Gigliotti, M.F.X. (1997). Dissolution rate measurements of TiN in Ti6242,

David, D. et al (1979). Etude de la diffusion de'oxygene dans le titane α oxyde enter 700OC et

Fromm, E. and Gebhardt, E. (1976). *Gase und kohlenstoff in metallen*, Springer-Verlag, ISBN

Grala, E. M. (1968). *Characterization of Alpha Segregation Defects in Ti-6Al-4V Alloy*, Technical

Le Claire, A.D. & Neumann, G. (1990). In: *Diffusion in Solid Metals and Alloys*, Mehrer H.

(Ed.), Vol.III-26, pp. 85- 212, Landolt-Börnstein, Springer-Verlag, ISBN3-540-50886-

950OC, J.Less-Comm. Met., Vol.65, No. 1, pp. 51-69, ISSN 0022-5088

Report AFML-TR-68-304, AD0845805, TRW, Inc. Cleveland, OH

Acta Mater., Vol.45, No. 1, pp. 357-370, ISSN 1359-6454

pool. In: *Light Metals*, U. Mannweiler (Ed.), 1.187-1.193, The Minerals, Metals

Metallurgical and Materials Transactions B, Vol.28b, pp. 1001 -1010, ISSN 1073-

liquid flow velocity around solid particles near the ingot axis.

The efficiency of identification using this method was 80%.

Materials Society, ISBN 0-87339-264-7

3540072551, Berlin, Heidelberg, New York

cannot be considered as an independent event.

**5. Acknowledgment** 

**6. References** 

technology of titanium alloys".

5615 (print version)

4, Berlin

Fig. 14. Distribution of defects over their non-dimensional coordinate *r/R*: g is the number of defects in the unit of the ingot volume; *N* is the number of defects; *r* is the distance from a defect to the ingot axis; *R* is the ingot radius.

#### **4. Conclusion**

A database on the characteristics of metallurgical gas-containing defects in titanium alloys was obtained. This was achieved by the long term monitoring of a industrial process in which the vacuum arc remelting (VAR) technology was used. It is the most complete one that is available at present. It contains data on location of defects in ingots, the sizes of defects, their microhardness, and the concentrations of nitrogen, oxygen, and aluminum in them. Almost all the data are related to the Ti–6Al–4V alloy.

Almost all the revealed defects contained excess concentrations of oxygen and nitrogen in comparison with the matrix of titanium alloys; the excess ranged from several times to several tens of times larger. In almost of cases, the concentration of oxygen was slightly higher than that of nitrogen. At the registered concentration level of gaseous impurities, the alloying of inclusions by nitrogen is the predominant factor in the formation of defects. On the whole, the data on the concentrations of nitrogen and oxygen in the defects agree with a generally accepted concept, according to which the sources of defects are particles of titanium and its alloys with a high concentration of oxygen and nitrogen atoms.

It was shown that the approach that was used in practice for the identification of defect sources based on the measurement of the concentration of aluminum and vanadium in defects has no any substantial grounds.

The distribution of defects over the height and radius of ingots was studied; it turned out to be strongly inhomogeneous. The defects were located predominately near the ingot base and the ingot axis. The number of defects in the unit volume of the metal in the lower part and near the axis of the ingots was almost 60 times greater than the average value in the ingots. The inhomogeneous distribution of defects in the ingots is due to different dissolution rate of defect sources during smelting in different parts of the ingot. In turn, this results from a low temperature of the liquid pool in the lower part of the ingots and a low liquid flow velocity around solid particles near the ingot axis.

A large number of cases have been registered when not one but two, three or even more defects are formed simultaneously in the ingot. The probabilities of formation of one and several defects were comparable. Usually, several defects in the ingot were located very close to each other. These findings indicate that generally the formation of each defect cannot be considered as an independent event.

A method for identification of the sources of metallurgical defects in the smelting of titanium alloys has been developed. It is based on the diffusion mechanism for the modification of defect sources in smelting and the established dependences of the aluminum concentration at the center of a defect on its size. The identification scheme included data on microhardness and the average concentration of gaseous impurities in the zone of a defect. The efficiency of identification using this method was 80%.

#### **5. Acknowledgment**

62 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9

*r/R*

Fig. 14. Distribution of defects over their non-dimensional coordinate *r/R*: g is the number of defects in the unit of the ingot volume; *N* is the number of defects; *r* is the distance from a

A database on the characteristics of metallurgical gas-containing defects in titanium alloys was obtained. This was achieved by the long term monitoring of a industrial process in which the vacuum arc remelting (VAR) technology was used. It is the most complete one that is available at present. It contains data on location of defects in ingots, the sizes of defects, their microhardness, and the concentrations of nitrogen, oxygen, and aluminum in

Almost all the revealed defects contained excess concentrations of oxygen and nitrogen in comparison with the matrix of titanium alloys; the excess ranged from several times to several tens of times larger. In almost of cases, the concentration of oxygen was slightly higher than that of nitrogen. At the registered concentration level of gaseous impurities, the alloying of inclusions by nitrogen is the predominant factor in the formation of defects. On the whole, the data on the concentrations of nitrogen and oxygen in the defects agree with a generally accepted concept, according to which the sources of defects are particles of

It was shown that the approach that was used in practice for the identification of defect sources based on the measurement of the concentration of aluminum and vanadium in

titanium and its alloys with a high concentration of oxygen and nitrogen atoms.

0.0

**4. Conclusion** 

defect to the ingot axis; *R* is the ingot radius.

defects has no any substantial grounds.

them. Almost all the data are related to the Ti–6Al–4V alloy.

0.2

0.4

*g,* **arb.units**

0.6

0.8

1.0

This work was supported by the project No 11-2-06-AVI of the Basic Research Program of the Ural Division of Russian Academy of Sciences "Identification of stages which are responsible for formation of gas-containing metallurgical defects for the industrial VAR technology of titanium alloys".

#### **6. References**


**Part 2** 

**Properties of Titanium Alloys** 

**High Pressure Conditions** 

**Under High Temperature and Ultra** 


**Part 2** 

**Properties of Titanium Alloys Under High Temperature and Ultra High Pressure Conditions** 

64 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Le Claire, A.D. (1990). In: *Diffusion in Solid Metals and Alloys*, Mehrer H. (Ed.), Vol.III-26, p.

Mitchell, A. (1984). Final Report to General Electric Corp., pp. 10-29, University of British

Reddy, R.G. (1990). Kinetics of TiN Dissolution in Ti Alloys. In: *Electron Beam Melting and* 

Schwartz, F. (1993). Technical Report YKOG, No 3028/93, pp. 2-7, SNECMA, Paris, France Tarenkova, N.Yu. et al. (2011), Formation of Gas-Saturated Defects in Titanium Alloys

Vykhodets, V.B. et al. (1987). Oxygen diffusion in α-Ti. I. Anisotropy of oxygen diffusion in α-Ti, The Phys. Metal & Metallogr., Vol.64, pp. 127-133, ISSN 0031-918X Vykhodets, V.B. et al. (2006). Studying distribution of gaseous impurities and carbon in

Vykhodets, V.B. et al. (2007). Genesis of Metallurgical Defects in Titanium Alloys, Doklady

Vykhodets, V.B. et al (2011). Identification of the sources of gas containing metallurgical

Wood, F. W. (1969). *Elimination of Low-Density Inclusions in Titanium Alloy Ingots,* Final

Metallography, 2006, V.101, No.3, pp. 267-275, ISSN 0031-918X

Testing, Vol.47, No. 3, pp. 176–188, ISSN 1061 8309

technical report 1 Sep 67-30 Aug 68, AD0852028

*Refining State of the Art,* R. Bakish (ed.), 119-127, NV: Bakish Materials, ISBN-10

during Vacuum-Arc Remelting, Russian Metallurgy (Metally), Vol.2011, No. 2, pp.

titanium alloys using nuclear microanalysis, The Physics of Metals and

Physical Chemistry, Vol.416, part 2, pp. 285-288, ISSN PRINT 0012-5016, ISSN

defects during the smelting of titanium alloys, Russian Journal of Nondestructive

471, Landolt-Börnstein, Springer-Verlag, ISBN3-540-50886-4, Berlin

Columbia, Vancouver, Canada

127–132, ISSN 0036-0295

ONLINE 1608-3121

9992384719, ISBN-13 978-9992384718, Reno

**4** 

*1,3USA 2UK* 

**Titanium Alloys at** 

*1Los Alamos National Laboratory 2Atomic Weapons Establishment* 

*3Lawrence Livermore National Laboratory* 

**Extreme Pressure Conditions** 

Nenad Velisavljevic1, Simon MacLeod2 and Hyunchae Cynn3

The electronic structures of the early transition metals are characterised by the relationship that exists between the occupied narrow *d* bands and the broad *sp* bands. Under pressure, the *sp* bands rise faster in energy, causing electrons to be transferred to the *d* bands (Gupta et al., 2008). This process is known as the *s-d* transition and it governs the structural

At ambient conditions, pure Ti crystallizes in the 2-atom hcp, or phase crystal structure (space group P63/mmc) and has an axial ratio (*c*/*a*) ~ 1.58. Under pressure, the phase undergoes a martensitic transformation at room temperature (RT) into the 3-atom hexagonal, or phase structure (space group P6/mmm). The appearance of the ω phase at high pressure raises a number of scientific and engineering issues mainly because the phase appears to be fairly brittle compared with the phase, and this may significantly limit the use of Ti in high pressure applications. Furthermore, after pressure treatment the ω phase appears to be fully, or at least, partially recoverable at ambient conditions, thus raising questions as to which is the lowest thermodynamically stable crystallographic phase

This chapter deals with the behavior of Ti alloys under extreme pressure and temperature conditions. Our recent results are presented and comparison is made to data available in the open literature. We volume compressed Ti-6Al-4V (an alloy) and Ti-Beta-21S (a β alloy) in a series of RT diamond anvil cell (DAC) angle-dispersive X-ray diffraction (ADXD) experiments, to investigate the effects of alloying and the pressure environment on phase stability and the transformation pathway. However, before describing the results of these experiments in detail, we present a brief review of the current state of knowledge of Ti at

The phase relations of Ti have been studied extensively at high pressure. As indicated in the introduction, there is great interest in the properties of Ti at high pressure and temperature.

**1. Introduction** 

properties of the transition metals.

of Ti at RT and pressure.

**2. Ti at high pressure** 

high pressure.

### **Titanium Alloys at Extreme Pressure Conditions**

Nenad Velisavljevic1, Simon MacLeod2 and Hyunchae Cynn3

*1Los Alamos National Laboratory 2Atomic Weapons Establishment 3Lawrence Livermore National Laboratory 1,3USA 2UK* 

#### **1. Introduction**

The electronic structures of the early transition metals are characterised by the relationship that exists between the occupied narrow *d* bands and the broad *sp* bands. Under pressure, the *sp* bands rise faster in energy, causing electrons to be transferred to the *d* bands (Gupta et al., 2008). This process is known as the *s-d* transition and it governs the structural properties of the transition metals.

At ambient conditions, pure Ti crystallizes in the 2-atom hcp, or phase crystal structure (space group P63/mmc) and has an axial ratio (*c*/*a*) ~ 1.58. Under pressure, the phase undergoes a martensitic transformation at room temperature (RT) into the 3-atom hexagonal, or phase structure (space group P6/mmm). The appearance of the ω phase at high pressure raises a number of scientific and engineering issues mainly because the phase appears to be fairly brittle compared with the phase, and this may significantly limit the use of Ti in high pressure applications. Furthermore, after pressure treatment the ω phase appears to be fully, or at least, partially recoverable at ambient conditions, thus raising questions as to which is the lowest thermodynamically stable crystallographic phase of Ti at RT and pressure.

This chapter deals with the behavior of Ti alloys under extreme pressure and temperature conditions. Our recent results are presented and comparison is made to data available in the open literature. We volume compressed Ti-6Al-4V (an alloy) and Ti-Beta-21S (a β alloy) in a series of RT diamond anvil cell (DAC) angle-dispersive X-ray diffraction (ADXD) experiments, to investigate the effects of alloying and the pressure environment on phase stability and the transformation pathway. However, before describing the results of these experiments in detail, we present a brief review of the current state of knowledge of Ti at high pressure.

#### **2. Ti at high pressure**

The phase relations of Ti have been studied extensively at high pressure. As indicated in the introduction, there is great interest in the properties of Ti at high pressure and temperature.

Titanium Alloys at Extreme Pressure Conditions 69

Vohra et al. investigated the effects of the impurity levels of oxygen on the transition (Vohra et al., 1977). In this non-hydrostatic DAC study (no PTM), the oxygen content of the Ti samples was varied between 785 ppm and 3800 ppm (by weight) and the corresponding

Ti has been volume compressed at RT (no PTM) in a DAC to a pressure of 220 GPa and found to follow the transformation pathway δ (Vohra & Spencer, 2001; Akahama et al., 2001). The phase structure was observed to be stable to pressures greater than 100 GPa. The reported intermediate orthorhombic phase (transforming between ~ 116 GPa - 128 GPa) possessed a distorted hcp crystal structure (space group Cmcm), and the orthorhombic phase (transforming at ~ 145 GPa) was of a distorted bcc type structure (space group Cmcm) (Akahama et al., 2001). These observations for Ti are at variance with the other Titanium Group transition metals Zr and Hf, which have been observed to follow the β pathway at RT (Jayaraman et al., 1963; Xia et al., 1990a; Xia et al., 1990b). However, Ahuja et al. have observed the coexistence of a bcc-like structure (referred to by the authors as β') with the ω phase, during the compression of Ti at RT from 40 GPa to 80 GPa (Ahuja et al., 2004). In this study, Ti was embedded in a NaCl PTM. The subsequent ADXD patterns could only be fully analysed on the assumption that the additional reflections present in the patterns belonged to the β phase of Ti. Laser heating Ti to between 1200 K and 1300 K at 78-80 GPa resulted in the formation of a new orthorhombic -phase (space group Fmmm), which on decompression, at RT, transformed back into the β' phase below 40 GPa (Ahuja et al., 2004). On further decompression below 30 GPa, the β' phase reverted back into the phase, and this phase could be quenched to ambient conditions

Most recently, nano-Ti sample (with grain sizes ~ 100 nm) was compressed at RT in a DAC to 161 GPa and the transformation pathway was observed (Velisavljevic et al., n.d.). The slightly high transition pressure of 10 GPa observed (no PTM), may have been caused by the increase in interface to volume ratio, resulting from the reduction in grain size, and leading to increased resistance to shear deformation. Compared to coarse grained Ti, in nano-Ti there may also have been a larger concentration of interstitial impurities near the grain boundaries, which have been shown to help suppress the structural phase transition (see Hennig et al., 2005). The phase was observed to be stable up to ~ 120 GPa, and under compression to 127 GPa resulted in a phase transformation to the orthorhombic phase. A further compression resulted in a transition from to the phase at 140 GPa, in good agreement with a previous study (Akahama et al., 2001). The phase was stable up to 161 GPa. Figure 1 shows a stacked plot of nano-Ti ADXD patterns in the , and phases (Velisavljevic et al., n.d.). The metastable phase was recovered after pressure treatment, which is consistent with reports from other experiments on Ti, indicating that after pressure release, samples were either recovered in the phase or as a mixture of (see Errandonea et al., 2005 & Vohra et al., 2001). In the case of these nano-Ti experiments, the recovered sample was observed to consist of only the phase (Velisavljevic, n.d.). The high pressure behavior of nano-Ti, including the structural phase sequence (without the appearance of the β phase), the change in axial *c*/*a* ratio with pressure, recovery of phase, and change in volume with pressure and the EOS values (Velisavljevic, n.d.), is very similar and consistent with previous experimental results

on Ti (Akahama et al., 2001; Vohra et al., 2001; Errandonea et al., 2005;).

transition pressure was measured between 2.9 GPa and 6.0 GPa.

(Ahuja et al., 2004).

The RT phase transition has been observed to occur between 2 GPa and 12 GPa, depending on the experimental technique, the pressure environment, and the sample purity. Although most of the recent work has involved the use of diamond anvil cells (DACs) to volume compress Ti in the static regime (for example, Ming et al., 1981; Akahama et al., 2001; Vohra et al., 2001; Errandonea et al., 2005), many studies have also utilised large volume presses to statically compress Ti (Jamieson 1963; Jayaraman et al., 1963; Bundy 1963; Bundy 1965; Zhang et al., 2008), and to a lesser extent, shock techniques have been employed to compress Ti in the dynamic regime (see Gray et al., 1993; Trunin et al., 1999; Cerreta et al., 2006).

#### **2.1 Dynamic compression of Ti**

The effects oxygen and other interstitial impurities on phase formation in Ti has been studied using shock compression (Gray et al., 1993 & Cerreta et al., 2006). In these studies, samples with oxygen content 360 ppm (high purity Ti) and 3700 ppm (A-70) were shocked whilst simultaneously probed using a real-time velocity interferometer system for any reflector (VISAR) diagnostic and then analysed post-shock. A phase transition was reported at 10.4 GPa for the high purity sample, whereas the A-70 sample did not show any evidence of a phase change up to 35 GPa (Gray et al., 1993 & Cerreta et al., 2006). A post shock (11 GPa) recovered high purity sample retained 28% of the phase (Gray et al., 1993). The suppression of the stress transformation in Ti is likely caused by the presence of interstitial oxygen (Cerreta et al., 2006). Greef et al. performed an analysis of early Ti shock measurements, but were unable to allow for sample purity in their study, since the oxygen content was not measured in these early experiments, and as a consequence placed the phase transition at ~ 12 GPa (Greef et al., 2001).

#### **2.2 Static compression of Ti**

#### **2.2.1 Room temperature compression**

Using DACs and X-ray diffraction, the experimental technique RT phase transition in commercially available purity Ti has been observed to occur between 2 GPa and 12 GPa – in cases where data was collected during pressure release large transformation pressure hysteresis is observed and in some cases the phase can be recovered at ambient conditions (Vohra et al., 1977; Ming et al., 1981; Vohra et al., 2001; Akahama et al., 2001; Errandonea et al., 2005). The effect of the pressure environment (and therefore uniaxial stress) on the transformation in Ti was studied in a series of DAC and angle dispersive X-ray diffraction (ADXD) experiments (Errandonea et al., 2005). By embedding Ti samples in a variety of pressure-transmitting-media (PTM), Errandonea et al were able to demonstrate that the pressure at which the Ti transformed into the phase was found to increase with an increase in the hydrostaticity of the PTM. For an argon PTM (the most hydrostatic PTM used in the experiment), the transition occurred in the pressure range between 10.5 GPa and 14.9 GPa, and for the least hydrostatic environment (that is, no PTM), the transition occurred under pressure between 4.9 GPa and 12.4 GPa (Errandonea et al., 2005). A coexistence of the and phases over a largish pressure range was also observed, in agreement with earlier findings (Ming et al., 1981).

The RT phase transition has been observed to occur between 2 GPa and 12 GPa, depending on the experimental technique, the pressure environment, and the sample purity. Although most of the recent work has involved the use of diamond anvil cells (DACs) to volume compress Ti in the static regime (for example, Ming et al., 1981; Akahama et al., 2001; Vohra et al., 2001; Errandonea et al., 2005), many studies have also utilised large volume presses to statically compress Ti (Jamieson 1963; Jayaraman et al., 1963; Bundy 1963; Bundy 1965; Zhang et al., 2008), and to a lesser extent, shock techniques have been employed to compress Ti in the dynamic regime (see Gray et al., 1993; Trunin et al., 1999;

The effects oxygen and other interstitial impurities on phase formation in Ti has been studied using shock compression (Gray et al., 1993 & Cerreta et al., 2006). In these studies, samples with oxygen content 360 ppm (high purity Ti) and 3700 ppm (A-70) were shocked whilst simultaneously probed using a real-time velocity interferometer system for any reflector (VISAR) diagnostic and then analysed post-shock. A phase transition was reported at 10.4 GPa for the high purity sample, whereas the A-70 sample did not show any evidence of a phase change up to 35 GPa (Gray et al., 1993 & Cerreta et al., 2006). A post shock (11 GPa) recovered high purity sample retained 28% of the phase (Gray et al., 1993). The suppression of the stress transformation in Ti is likely caused by the presence of interstitial oxygen (Cerreta et al., 2006). Greef et al. performed an analysis of early Ti shock measurements, but were unable to allow for sample purity in their study, since the oxygen content was not measured in these early experiments, and as a consequence placed the

Using DACs and X-ray diffraction, the experimental technique RT phase transition in commercially available purity Ti has been observed to occur between 2 GPa and 12 GPa – in cases where data was collected during pressure release large transformation pressure hysteresis is observed and in some cases the phase can be recovered at ambient conditions (Vohra et al., 1977; Ming et al., 1981; Vohra et al., 2001; Akahama et al., 2001; Errandonea et al., 2005). The effect of the pressure environment (and therefore uniaxial stress) on the transformation in Ti was studied in a series of DAC and angle dispersive X-ray diffraction (ADXD) experiments (Errandonea et al., 2005). By embedding Ti samples in a variety of pressure-transmitting-media (PTM), Errandonea et al were able to demonstrate that the pressure at which the Ti transformed into the phase was found to increase with an increase in the hydrostaticity of the PTM. For an argon PTM (the most hydrostatic PTM used in the experiment), the transition occurred in the pressure range between 10.5 GPa and 14.9 GPa, and for the least hydrostatic environment (that is, no PTM), the transition occurred under pressure between 4.9 GPa and 12.4 GPa (Errandonea et al., 2005). A coexistence of the and phases over a largish pressure range was also observed, in

Cerreta et al., 2006).

**2.1 Dynamic compression of Ti** 

phase transition at ~ 12 GPa (Greef et al., 2001).

agreement with earlier findings (Ming et al., 1981).

**2.2 Static compression of Ti** 

**2.2.1 Room temperature compression** 

Vohra et al. investigated the effects of the impurity levels of oxygen on the transition (Vohra et al., 1977). In this non-hydrostatic DAC study (no PTM), the oxygen content of the Ti samples was varied between 785 ppm and 3800 ppm (by weight) and the corresponding transition pressure was measured between 2.9 GPa and 6.0 GPa.

Ti has been volume compressed at RT (no PTM) in a DAC to a pressure of 220 GPa and found to follow the transformation pathway δ (Vohra & Spencer, 2001; Akahama et al., 2001). The phase structure was observed to be stable to pressures greater than 100 GPa. The reported intermediate orthorhombic phase (transforming between ~ 116 GPa - 128 GPa) possessed a distorted hcp crystal structure (space group Cmcm), and the orthorhombic phase (transforming at ~ 145 GPa) was of a distorted bcc type structure (space group Cmcm) (Akahama et al., 2001). These observations for Ti are at variance with the other Titanium Group transition metals Zr and Hf, which have been observed to follow the β pathway at RT (Jayaraman et al., 1963; Xia et al., 1990a; Xia et al., 1990b). However, Ahuja et al. have observed the coexistence of a bcc-like structure (referred to by the authors as β') with the ω phase, during the compression of Ti at RT from 40 GPa to 80 GPa (Ahuja et al., 2004). In this study, Ti was embedded in a NaCl PTM. The subsequent ADXD patterns could only be fully analysed on the assumption that the additional reflections present in the patterns belonged to the β phase of Ti. Laser heating Ti to between 1200 K and 1300 K at 78-80 GPa resulted in the formation of a new orthorhombic -phase (space group Fmmm), which on decompression, at RT, transformed back into the β' phase below 40 GPa (Ahuja et al., 2004). On further decompression below 30 GPa, the β' phase reverted back into the phase, and this phase could be quenched to ambient conditions (Ahuja et al., 2004).

Most recently, nano-Ti sample (with grain sizes ~ 100 nm) was compressed at RT in a DAC to 161 GPa and the transformation pathway was observed (Velisavljevic et al., n.d.). The slightly high transition pressure of 10 GPa observed (no PTM), may have been caused by the increase in interface to volume ratio, resulting from the reduction in grain size, and leading to increased resistance to shear deformation. Compared to coarse grained Ti, in nano-Ti there may also have been a larger concentration of interstitial impurities near the grain boundaries, which have been shown to help suppress the structural phase transition (see Hennig et al., 2005). The phase was observed to be stable up to ~ 120 GPa, and under compression to 127 GPa resulted in a phase transformation to the orthorhombic phase. A further compression resulted in a transition from to the phase at 140 GPa, in good agreement with a previous study (Akahama et al., 2001). The phase was stable up to 161 GPa. Figure 1 shows a stacked plot of nano-Ti ADXD patterns in the , and phases (Velisavljevic et al., n.d.). The metastable phase was recovered after pressure treatment, which is consistent with reports from other experiments on Ti, indicating that after pressure release, samples were either recovered in the phase or as a mixture of (see Errandonea et al., 2005 & Vohra et al., 2001). In the case of these nano-Ti experiments, the recovered sample was observed to consist of only the phase (Velisavljevic, n.d.). The high pressure behavior of nano-Ti, including the structural phase sequence (without the appearance of the β phase), the change in axial *c*/*a* ratio with pressure, recovery of phase, and change in volume with pressure and the EOS values (Velisavljevic, n.d.), is very similar and consistent with previous experimental results on Ti (Akahama et al., 2001; Vohra et al., 2001; Errandonea et al., 2005;).

Titanium Alloys at Extreme Pressure Conditions 71

heating appeared to lower the transition pressure in pure T, as was determined in a follow up DAC study conducted by Errandonea at a synchrotron (Errandonea et al., 2005). At 5 GPa, Ti was transformed into the β phase by laser-heating to 1750 K and 2150 K (above melting), at which point the samples were quenched. A mixture of and phases were obtained at a pressure for which only the phase existed previously at RT, suggesting that thermal fluctuations may have a similar effect on the transformation as uniaxial

Although Ti has received much theoretical attention over the years, we will mention here only studies that are of direct relevance to this chapter. First-principles calculations have been used to generate the phase diagram of Ti (Trinkle et al., 2003; Pecker et al., 2005; Hennig et al., 2005; Trinkle et al., 2005; Verma et al., 2007; Hennig et al., 2008; Mei et al.,

A multiphase equation of state (EOS) of the three solid phases (, β and ), the liquid phase and gas phase, was calculated up to 100 GPa and predicted (based on experiment) the -βliquid triple point at around 45 GPa and 2200 K (Pecker et al., 2005). Verma et al. predicted the ω β pathway for Ti, and found the phase to be energetically unstable under hydrostatic conditions (Verma et al., 2007). The phase was found to be elastically stable between 102 GPa and 112 GPa. However, under non-hydrostatic conditions, the authors predicted the phase to exist over a larger pressure range (Verma et al., 2007). The influence of anisotropic stresses under very non-hydrostatic conditions may support the existence of the phase reported in DAC experiments (Akahama et al., 2001; Vohra et al., 2001;

The actual mechanism behind the martensitic transformation between the , β and phases in Ti has been explored in a series of molecular dynamics simulations (Trinkle et al., 2003; Trinkle et al., 2005; Hennig et al., 2005; Hennig et al., 2008). The authors propose the lowest energy pathway for the transition to be the TAO-1, ("titanium alpha to omega"), in which atoms in the phase transform through small shuffles and strains into the phase, without going through a metastable intermediate phase (Trinkle et al, 2003 & Trinkle et al., 2005). The presence of impurities in Ti such as O, N and C can affect the energy barrier to the phase and suppress the transformation (Hennig et al., 2005). Hennig et al. predicted the phase boundaries for Ti up to 15 GPa, and the -β- triple point at 8 GPa and

Mei et al. studied the thermodynamic properties and phase diagram of Ti and predicted the RP β transition at 1114 K and the β triple point at 11.1 GPa and 821 K (Mei et al., 2009). The transition was predicted at 1.8 GPa, slightly lower than experimental measurements. Hu et al. performed a detailed calculation to predict the phase diagram, thermal EOS and thermodynamic properties of Ti (Hu et al., 2010). The axial ratio of the phase was predicted to be almost invariant with pressure, in agreement with the anvil study of Zhang et al., and with calculation (Zhang et al., 2008 & Mei et al., 2009), but not with a DAC study in which the compression was found to be anisotropic (Errandonea et al., 2005). Hu et al. calculated the RT transition at 2.02 GPa and the triple point at 9.78 GPa and

stress.

2009; Hu et al., 2010).

Velisavljevic, n.d.).

1200 K. (Hennig et al., 2008).

**2.3 Theoretical treatments of Ti at high pressure** 

Fig. 1. A stack of ADXD patterns showing pressure induced structural phase transformation in nano-Ti (Velisavljevic et al., n.d.).

#### **2.2.2 High temperature compression**

Very few combined static high-pressure and high-temperature studies have been reported for Ti. Thermally treating Ti to temperatures exceeding 1155 K at room pressure (RP) transforms the phase into the more densely packed bcc or β phase (space group Im3m), without the intermediate or other crystallographic phase occurring.

Jayaraman et al. studied the β boundary up to ~ 6.5 GPa and 1100 K, and found that increasing the pressure lowered the β transition temperature (Jayaraman et al., 1963). A more extensive study conducted by Bundy, in which Ti was compressed to 16 GPa and heated to 1200 K revealed the location of the boundaries of the three solid phases, , β and (Bundy, 1963 & Bundy, 1965). More recently, Ti was compressed in a cubic anvil apparatus to 8.7 GPa and heated to 973 K (Zhang et al., 2008). The ADXD patterns collected confirmed the high-temperature phase diagram reported by Bundy. The --β triple point was estimated at 7.5 GPa and 913 K (Zhang et al., 2008), in agreement with the previous estimate of 8.0 0.7 GPa and 913 50 K (Bundy, 1965). Zhang et al. observed the -phase to undergo an isotropic compression between RP and 7.8 GPa, resulting in the axial ratio (*c*/*a*) being constant (1.587) over this pressure range (Zhang et al., 2008).

Errandonea et al. melted Ti in a DAC up to 80 GPa using single-sided laser-heating and the "speckle technique" to determine the onset of melting (Errandonea et al., 2001). No X-ray diffraction patterns were collected in this lab-based study, and so the authors were unable to discriminate between melting from either the β phase or the phase. Short term laser-

Fig. 1. A stack of ADXD patterns showing pressure induced structural phase transformation

Very few combined static high-pressure and high-temperature studies have been reported for Ti. Thermally treating Ti to temperatures exceeding 1155 K at room pressure (RP) transforms the phase into the more densely packed bcc or β phase (space group Im3m),

Jayaraman et al. studied the β boundary up to ~ 6.5 GPa and 1100 K, and found that increasing the pressure lowered the β transition temperature (Jayaraman et al., 1963). A more extensive study conducted by Bundy, in which Ti was compressed to 16 GPa and heated to 1200 K revealed the location of the boundaries of the three solid phases, , β and (Bundy, 1963 & Bundy, 1965). More recently, Ti was compressed in a cubic anvil apparatus to 8.7 GPa and heated to 973 K (Zhang et al., 2008). The ADXD patterns collected confirmed the high-temperature phase diagram reported by Bundy. The --β triple point was estimated at 7.5 GPa and 913 K (Zhang et al., 2008), in agreement with the previous estimate of 8.0 0.7 GPa and 913 50 K (Bundy, 1965). Zhang et al. observed the -phase to undergo an isotropic compression between RP and 7.8 GPa, resulting in the axial ratio (*c*/*a*) being

Errandonea et al. melted Ti in a DAC up to 80 GPa using single-sided laser-heating and the "speckle technique" to determine the onset of melting (Errandonea et al., 2001). No X-ray diffraction patterns were collected in this lab-based study, and so the authors were unable to discriminate between melting from either the β phase or the phase. Short term laser-

without the intermediate or other crystallographic phase occurring.

constant (1.587) over this pressure range (Zhang et al., 2008).

in nano-Ti (Velisavljevic et al., n.d.).

**2.2.2 High temperature compression** 

heating appeared to lower the transition pressure in pure T, as was determined in a follow up DAC study conducted by Errandonea at a synchrotron (Errandonea et al., 2005). At 5 GPa, Ti was transformed into the β phase by laser-heating to 1750 K and 2150 K (above melting), at which point the samples were quenched. A mixture of and phases were obtained at a pressure for which only the phase existed previously at RT, suggesting that thermal fluctuations may have a similar effect on the transformation as uniaxial stress.

#### **2.3 Theoretical treatments of Ti at high pressure**

Although Ti has received much theoretical attention over the years, we will mention here only studies that are of direct relevance to this chapter. First-principles calculations have been used to generate the phase diagram of Ti (Trinkle et al., 2003; Pecker et al., 2005; Hennig et al., 2005; Trinkle et al., 2005; Verma et al., 2007; Hennig et al., 2008; Mei et al., 2009; Hu et al., 2010).

A multiphase equation of state (EOS) of the three solid phases (, β and ), the liquid phase and gas phase, was calculated up to 100 GPa and predicted (based on experiment) the -βliquid triple point at around 45 GPa and 2200 K (Pecker et al., 2005). Verma et al. predicted the ω β pathway for Ti, and found the phase to be energetically unstable under hydrostatic conditions (Verma et al., 2007). The phase was found to be elastically stable between 102 GPa and 112 GPa. However, under non-hydrostatic conditions, the authors predicted the phase to exist over a larger pressure range (Verma et al., 2007). The influence of anisotropic stresses under very non-hydrostatic conditions may support the existence of the phase reported in DAC experiments (Akahama et al., 2001; Vohra et al., 2001; Velisavljevic, n.d.).

The actual mechanism behind the martensitic transformation between the , β and phases in Ti has been explored in a series of molecular dynamics simulations (Trinkle et al., 2003; Trinkle et al., 2005; Hennig et al., 2005; Hennig et al., 2008). The authors propose the lowest energy pathway for the transition to be the TAO-1, ("titanium alpha to omega"), in which atoms in the phase transform through small shuffles and strains into the phase, without going through a metastable intermediate phase (Trinkle et al, 2003 & Trinkle et al., 2005). The presence of impurities in Ti such as O, N and C can affect the energy barrier to the phase and suppress the transformation (Hennig et al., 2005). Hennig et al. predicted the phase boundaries for Ti up to 15 GPa, and the -β- triple point at 8 GPa and 1200 K. (Hennig et al., 2008).

Mei et al. studied the thermodynamic properties and phase diagram of Ti and predicted the RP β transition at 1114 K and the β triple point at 11.1 GPa and 821 K (Mei et al., 2009). The transition was predicted at 1.8 GPa, slightly lower than experimental measurements. Hu et al. performed a detailed calculation to predict the phase diagram, thermal EOS and thermodynamic properties of Ti (Hu et al., 2010). The axial ratio of the phase was predicted to be almost invariant with pressure, in agreement with the anvil study of Zhang et al., and with calculation (Zhang et al., 2008 & Mei et al., 2009), but not with a DAC study in which the compression was found to be anisotropic (Errandonea et al., 2005). Hu et al. calculated the RT transition at 2.02 GPa and the triple point at 9.78 GPa and

Titanium Alloys at Extreme Pressure Conditions 73

Ti alloys is governed by the presence of impurities, and for a ternary alloy such as Ti-6Al-4V, the substitutional impurities, Al (which is an -phase stabilizer) and V (a β-phase stabilizer), influence the onset of the phase transformation by changing the *d* electron concentration in the alloy (Vohra, 1979). The addition of Al reduces the *d* band concentration, whilst the addition of V increases it by one, thus resulting in an overall reduction in the *d* band concentration. Interstitial impurities such as O, N and C can retard the transformation. *Ab initio* calculations have shown that the presence of these impurities can affect the relative phase stability and the energy barrier of the phase transformation (Hennig et al., 2005). The presence of impurities in the commercial Ti alloys A-70 and Ti-6Al-4V, particularly O and Al, suppresses the onset of phase transformation by increasing the energy and energy barrier of relative to (Hennig et al., 2005). Thus, the stability range of the phase at RT is increased. Hennig et al. predicted the RT phase transformations in A-70 and Ti-6Al-4V to occur at 31 GPa and 63 GPa

As the most prevalent Ti alloy currently in commercial and industrial usage, it is perhaps not surprising that Ti-6Al-4V has received the most attention of all the Ti alloys at high pressure (Rosenberg & Meybar, 1981; Gray et al., 1993; Cerreta et al., 2006; Chesnut et al., 2008; Halevy et al., 2010; Tegner et al., 2011; Tegner et al., n.d). The alloying of a metal can have substantial effects on its properties, which of course is the desired effect, and so it is

In the dynamic regime, the Hugoniot curve of Ti-6Al-4V was generated up to 14 GPa using powder-gun driven shock waves and manganin stress gauges (Rosenberg & Meybar, 1981). A break in the stress-particle velocity curve near 10 GPa was indicative, the authors proposed, of a possible phase transformation, though they were unable to state unequivocally that the transformation was (Rosenberg & Meybar, 1981). As part of a study to examine the effects of alloy chemistry on phase formation in Ti alloys, Gray et al., shocked Ti-6Al-4V (oxygen content 0.18 by weight %) up to 25 GPa and used VISAR to measure wave profiles (Gray et al., 1993). No evidence of a phase transformation was detected using VISAR, and using neutron diffraction to analyse the recovered specimen did

The first study of Ti-6Al-4V in a DAC was reported by Chesnut et al., in which a sample was loaded into a 4:1 methanol: ethanol PTM and compressed to 37 GPa (Chesnut et al., 2008). The ambient conditions volume of Ti-6Al-4V (predominantly in the phase) was measured using ADXD and found to be *V0* = 17.208 Å3/atom. Chesnut et al. observed Ti-6Al-4V to undergo the phase transition at ~ 27.3 GPa (Chesnut et al., 2008). The phase was observed to be stable to 37 GPa (the pressure limit of the experiment). The volume change across the phase boundary, at around 1%, was considered too small to detect in shock experiments and may explain why this transformation has yet to be observed in

important that we understand how the alloy responds to extremes of pressure.

respectively (Hennig et al., 2005).

**3.1 Ti-6Al-4V at high pressure** 

**3.1.1 Dynamic compression of Ti-6Al-4V** 

not reveal the presence of -phase structure.

**3.1.2 Static compression of Ti-6Al-4V** 

931 K, which is in close agreement with experiment (Bundy 1965 & Zhang et al., 2008). The slope of the - boundary (*dT*/*dP* = 81 K/GPa) differs significantly to that measured by Zhang et al. in their cubic anvil study (*dT*/*dP* = 345 K/GPa) (Zhang et al., 2008). The slope of the -β boundary was calculated to be 2.4 K/GPa (Hu et al., 2010), in good agreement with earlier measurements (Bundy, 1965). The predicted RT ω transformation at 110 GPa concurred with the calculation of Verma et al. (Verma et al., 2007).

For reference, we show in figure 2 a representation of the Ti phase diagram as a function of pressure and temperature, based loosely on the experimental and theoretical studies discussed in this review.

Fig. 2. The *P-T* phase diagram of Ti (this is a representation based on published work).

#### **3. Ti alloys at high pressure**

Ti alloys are usually classified according to their phase stability. An alloy consists mainly of an -phase stabilizing element such as Al, O, N or C, which has the effect of extending the range of the more ductile -phase field to higher temperatures and higher pressures. Similarly, a β alloy contains a β-phase stabilizing element such as Mo, V or Ta, the presence of which will shift the β phase field to lower temperatures. Ti alloys containing a combination of both -phase and β-phase stabilizing elements are by far the most widely used alloys commercially. At ambient conditions, the +β alloys possess a β-phase fraction by volume that lies somewhere between 5 and 50% and so crystallizes predominantly in the phase. Of particular importance is the +β alloy Ti-6Al-4V (wt%), which has found many commercial applications as a result of its superior material properties. The phase stability in

931 K, which is in close agreement with experiment (Bundy 1965 & Zhang et al., 2008). The slope of the - boundary (*dT*/*dP* = 81 K/GPa) differs significantly to that measured by Zhang et al. in their cubic anvil study (*dT*/*dP* = 345 K/GPa) (Zhang et al., 2008). The slope of the -β boundary was calculated to be 2.4 K/GPa (Hu et al., 2010), in good agreement with earlier measurements (Bundy, 1965). The predicted RT ω transformation at 110 GPa

For reference, we show in figure 2 a representation of the Ti phase diagram as a function of pressure and temperature, based loosely on the experimental and theoretical studies

Fig. 2. The *P-T* phase diagram of Ti (this is a representation based on published work).

Ti alloys are usually classified according to their phase stability. An alloy consists mainly of an -phase stabilizing element such as Al, O, N or C, which has the effect of extending the range of the more ductile -phase field to higher temperatures and higher pressures. Similarly, a β alloy contains a β-phase stabilizing element such as Mo, V or Ta, the presence of which will shift the β phase field to lower temperatures. Ti alloys containing a combination of both -phase and β-phase stabilizing elements are by far the most widely used alloys commercially. At ambient conditions, the +β alloys possess a β-phase fraction by volume that lies somewhere between 5 and 50% and so crystallizes predominantly in the phase. Of particular importance is the +β alloy Ti-6Al-4V (wt%), which has found many commercial applications as a result of its superior material properties. The phase stability in

concurred with the calculation of Verma et al. (Verma et al., 2007).

discussed in this review.

**3. Ti alloys at high pressure** 

Ti alloys is governed by the presence of impurities, and for a ternary alloy such as Ti-6Al-4V, the substitutional impurities, Al (which is an -phase stabilizer) and V (a β-phase stabilizer), influence the onset of the phase transformation by changing the *d* electron concentration in the alloy (Vohra, 1979). The addition of Al reduces the *d* band concentration, whilst the addition of V increases it by one, thus resulting in an overall reduction in the *d* band concentration. Interstitial impurities such as O, N and C can retard the transformation. *Ab initio* calculations have shown that the presence of these impurities can affect the relative phase stability and the energy barrier of the phase transformation (Hennig et al., 2005). The presence of impurities in the commercial Ti alloys A-70 and Ti-6Al-4V, particularly O and Al, suppresses the onset of phase transformation by increasing the energy and energy barrier of relative to (Hennig et al., 2005). Thus, the stability range of the phase at RT is increased. Hennig et al. predicted the RT phase transformations in A-70 and Ti-6Al-4V to occur at 31 GPa and 63 GPa respectively (Hennig et al., 2005).

#### **3.1 Ti-6Al-4V at high pressure**

As the most prevalent Ti alloy currently in commercial and industrial usage, it is perhaps not surprising that Ti-6Al-4V has received the most attention of all the Ti alloys at high pressure (Rosenberg & Meybar, 1981; Gray et al., 1993; Cerreta et al., 2006; Chesnut et al., 2008; Halevy et al., 2010; Tegner et al., 2011; Tegner et al., n.d). The alloying of a metal can have substantial effects on its properties, which of course is the desired effect, and so it is important that we understand how the alloy responds to extremes of pressure.

#### **3.1.1 Dynamic compression of Ti-6Al-4V**

In the dynamic regime, the Hugoniot curve of Ti-6Al-4V was generated up to 14 GPa using powder-gun driven shock waves and manganin stress gauges (Rosenberg & Meybar, 1981). A break in the stress-particle velocity curve near 10 GPa was indicative, the authors proposed, of a possible phase transformation, though they were unable to state unequivocally that the transformation was (Rosenberg & Meybar, 1981). As part of a study to examine the effects of alloy chemistry on phase formation in Ti alloys, Gray et al., shocked Ti-6Al-4V (oxygen content 0.18 by weight %) up to 25 GPa and used VISAR to measure wave profiles (Gray et al., 1993). No evidence of a phase transformation was detected using VISAR, and using neutron diffraction to analyse the recovered specimen did not reveal the presence of -phase structure.

#### **3.1.2 Static compression of Ti-6Al-4V**

The first study of Ti-6Al-4V in a DAC was reported by Chesnut et al., in which a sample was loaded into a 4:1 methanol: ethanol PTM and compressed to 37 GPa (Chesnut et al., 2008). The ambient conditions volume of Ti-6Al-4V (predominantly in the phase) was measured using ADXD and found to be *V0* = 17.208 Å3/atom. Chesnut et al. observed Ti-6Al-4V to undergo the phase transition at ~ 27.3 GPa (Chesnut et al., 2008). The phase was observed to be stable to 37 GPa (the pressure limit of the experiment). The volume change across the phase boundary, at around 1%, was considered too small to detect in shock experiments and may explain why this transformation has yet to be observed in

Titanium Alloys at Extreme Pressure Conditions 75

Fig. 3. A stack of ADXD patterns showing structural change in Ti-6Al-4V with increasing

PTM (figure 3), we used a well known Cu shock study (Carter et al., 1971).

We determined the pressure in our experiments by analysing the reflections from the pressure markers (either Ta or Cu) and using a known EOS from previous shock measurements. In the case of the Cu marker used in the compression of Ti-6Al-4V in a neon

We now show in the *P-V* plot in figure 4 our measurements for Ti-6Al-4V embedded in a methanol: ethanol PTM, alongside previous DAC measurements (Chesnut et al., 2008 & Halevy et al., 2010). There is good agreement between the Chesnut et al. and our (Tegner et al., n.d.) measurements (both using a methanol: ethanol PTM), but not with the Halevy et al measurements (there was no mention of a PTM being used in their study). We measured the

 = 31.2 GPa, and an isothermal bulk modulus of *K* = 115 3 GPa and pressure derivative of *K'* = 3.22 0.22 after fitting a Vinet EOS (Vinet et al., 1987) to the data (Tegner et al., 2011). The volume change across the phase boundary was measured to be less than 1%, in agreement with the previous methanol: ethanol study (Chesnut et al., 2008).

The axial ratio (*c*/*a*) for the phase of Ti-6Al-4V (at ambient conditions) was measured to be 1.602, which is slightly higher than that reported for pure Ti (1.587), but is an expected result due to the presence of the stabiliser Al (possessing a smaller atomic radius than Ti) in the

For the phase of Ti-6Al-4V, in a methanol: ethanol PTM, we found the axial ratio to be almost constant between 1.600 and 1.602 up to 42 GPa, see figure 5. The measured *c*/*a* ratios for the phase are also effectively constant between 34 GPa and 74 GPa, varying between 0.616 and 0.617. We find similar results for a loading of Ti-6Al-4V in a neon PTM. These measurements are in broad agreement with those reported by Errandonea et al. for pure Ti

pressure (Tegnet et al., 2011).

*P→*

alloy.

dynamically driven Ti-6Al-4V (Rosenberg & Meybar, 1981 & Gray et al., 1993). Fitting a 3rd order Birch-Murnaghan EOS (Birch 1952) to the data generated an isothermal bulk modulus (which is a measure of the incompressibility of a material) for the phase of *K* = 125.24 GPa and the pressure derivative of the isothermal bulk modulus, *K'* = 2.409.

Halevy et al. compressed a sample of Ti-6Al-4V to 32.4 GPa in a DAC, and using energy dispersive X-ray diffraction (EDXD), did not observe a transformation to the phase (Halevy et al., 2010). A Vinet EOS (Vinet et al., 1987) fit to the experimental data returned values for the isothermal bulk modulus and the pressure derivative of *K* = 154 11 GPa and *K'* = 5.4 1.4. No mention was made of a PTM being used in this study.

#### **3.1.2.1 Effects of pressure media on the phase relations of Ti-6Al-4V**

In the most recent DAC study of Ti-6Al-4V, conducted by two of us (MacLeod and Cynn), powdered polycrystalline samples were embedded in a variety of PTMs to investigate the effects of the pressure environment on the RT phase transformation (Tegner et al., 2011; Tegner et al., n.d.). ADXD data were collected at the High Pressure Collaborative Access Team (HP-CAT) beamline 16-IDB at the Advanced Photon Source (APS) in Chicago, for Ti-6Al-4V samples embedded in neon, 4:1 methanol: ethanol and mineral oil. The oxygen content of the Ti-6Al-4V was 0.123 by weight %. The ambient conditions volume in the phase was measured to be *V0* = 17.252 Å3/atom. We observed the phase transformation to occur at 32.7 GPa for Ti-6Al-4V in the neon PTM, 31.2 GPa for the methanol: ethanol PTM and 26.2 GPa for the mineral oil PTM (in order of decreasing hydrostaticity in the PTM). At elevated pressures, ultimately all PTMs will become nonhydrostatic in nature (see Klotz et al., 2009, for a general discussion on the hydrostatic limits of various pressure media) and so it becomes more difficult to quantify the dependence of the phase transformation pressure in Ti-6Al-4V based on the hydrostaticity of the pressure environment, unlike in Ti where the transition is observed at a much lower pressure (see Errandonea et al., 2005).

A coexistence of the and phases was observed over a largish pressure range (of the order ~ 10 GPa or greater) for the different PTMs, similar to what was observed for Ti. A stacked plot of ADXD patterns, showing the structural response of Ti-6Al-4V to applied pressure, in a neon PTM, is shown in figure 3. Reflections from the Ti-6Al-4V sample in both and phases are present, together with reflections from neon (PTM) and copper (the pressure marker).

In figure 3, at 30.7 GPa, we observe the (100), (002), (101), (102) and (110) peaks that are characteristic of the phase. The dominant (110/101) reflections corresponding to the phase appear at ~ 32.7 GPa (between the phase (002) and (101) peaks) and then gradually grow in magnitude with pressure, whilst simultaneously the phase peaks diminish in intensity, until the pressure reaches ~ 44 GPa. By 44 GPa, the transformation is virtually completed. The phase (001), (201) and (210) reflections appear at a slightly higher pressure than the (110/101) peaks, at around 36 GPa to 39 GPa. In all, up to 10 phase peaks were indexed in this study. It is clear from figure 3 that both the and phases coexist over a large pressure range, of the order of 10 GPa (between ~ 32.7 GPa and ~ 44 GPa). We observed similar behaviour for Ti-6Al-4V embedded in methanol: ethanol and mineral oil, and also for an experiment with no PTM (Tegner et al., 2011). For Ti-6Al-4V embedded in methanol: ethanol, we decompressed our DAC from 75 GPa to ambient and observed the phase to gradually revert back to the phase.

dynamically driven Ti-6Al-4V (Rosenberg & Meybar, 1981 & Gray et al., 1993). Fitting a 3rd order Birch-Murnaghan EOS (Birch 1952) to the data generated an isothermal bulk modulus (which is a measure of the incompressibility of a material) for the phase of *K* = 125.24 GPa

Halevy et al. compressed a sample of Ti-6Al-4V to 32.4 GPa in a DAC, and using energy dispersive X-ray diffraction (EDXD), did not observe a transformation to the phase (Halevy et al., 2010). A Vinet EOS (Vinet et al., 1987) fit to the experimental data returned values for the isothermal bulk modulus and the pressure derivative of *K* = 154 11 GPa and

In the most recent DAC study of Ti-6Al-4V, conducted by two of us (MacLeod and Cynn), powdered polycrystalline samples were embedded in a variety of PTMs to investigate the effects of the pressure environment on the RT phase transformation (Tegner et al., 2011; Tegner et al., n.d.). ADXD data were collected at the High Pressure Collaborative Access Team (HP-CAT) beamline 16-IDB at the Advanced Photon Source (APS) in Chicago, for Ti-6Al-4V samples embedded in neon, 4:1 methanol: ethanol and mineral oil. The oxygen content of the Ti-6Al-4V was 0.123 by weight %. The ambient conditions volume in the phase was measured to be *V0* = 17.252 Å3/atom. We observed the phase transformation to occur at 32.7 GPa for Ti-6Al-4V in the neon PTM, 31.2 GPa for the methanol: ethanol PTM and 26.2 GPa for the mineral oil PTM (in order of decreasing hydrostaticity in the PTM). At elevated pressures, ultimately all PTMs will become nonhydrostatic in nature (see Klotz et al., 2009, for a general discussion on the hydrostatic limits of various pressure media) and so it becomes more difficult to quantify the dependence of the phase transformation pressure in Ti-6Al-4V based on the hydrostaticity of the pressure environment, unlike in Ti where the transition is observed at a much lower

A coexistence of the and phases was observed over a largish pressure range (of the order ~ 10 GPa or greater) for the different PTMs, similar to what was observed for Ti. A stacked plot of ADXD patterns, showing the structural response of Ti-6Al-4V to applied pressure, in a neon PTM, is shown in figure 3. Reflections from the Ti-6Al-4V sample in both and phases are

In figure 3, at 30.7 GPa, we observe the (100), (002), (101), (102) and (110) peaks that are characteristic of the phase. The dominant (110/101) reflections corresponding to the phase appear at ~ 32.7 GPa (between the phase (002) and (101) peaks) and then gradually grow in magnitude with pressure, whilst simultaneously the phase peaks diminish in intensity, until the pressure reaches ~ 44 GPa. By 44 GPa, the transformation is virtually completed. The phase (001), (201) and (210) reflections appear at a slightly higher pressure than the (110/101) peaks, at around 36 GPa to 39 GPa. In all, up to 10 phase peaks were indexed in this study. It is clear from figure 3 that both the and phases coexist over a large pressure range, of the order of 10 GPa (between ~ 32.7 GPa and ~ 44 GPa). We observed similar behaviour for Ti-6Al-4V embedded in methanol: ethanol and mineral oil, and also for an experiment with no PTM (Tegner et al., 2011). For Ti-6Al-4V embedded in methanol: ethanol, we decompressed our DAC from 75 GPa to ambient and

present, together with reflections from neon (PTM) and copper (the pressure marker).

observed the phase to gradually revert back to the phase.

and the pressure derivative of the isothermal bulk modulus, *K'* = 2.409.

*K'* = 5.4 1.4. No mention was made of a PTM being used in this study. **3.1.2.1 Effects of pressure media on the phase relations of Ti-6Al-4V** 

pressure (see Errandonea et al., 2005).

Fig. 3. A stack of ADXD patterns showing structural change in Ti-6Al-4V with increasing pressure (Tegnet et al., 2011).

We determined the pressure in our experiments by analysing the reflections from the pressure markers (either Ta or Cu) and using a known EOS from previous shock measurements. In the case of the Cu marker used in the compression of Ti-6Al-4V in a neon PTM (figure 3), we used a well known Cu shock study (Carter et al., 1971).

We now show in the *P-V* plot in figure 4 our measurements for Ti-6Al-4V embedded in a methanol: ethanol PTM, alongside previous DAC measurements (Chesnut et al., 2008 & Halevy et al., 2010). There is good agreement between the Chesnut et al. and our (Tegner et al., n.d.) measurements (both using a methanol: ethanol PTM), but not with the Halevy et al measurements (there was no mention of a PTM being used in their study). We measured the *P→* = 31.2 GPa, and an isothermal bulk modulus of *K* = 115 3 GPa and pressure derivative of *K'* = 3.22 0.22 after fitting a Vinet EOS (Vinet et al., 1987) to the data (Tegner et al., 2011). The volume change across the phase boundary was measured to be less than 1%, in agreement with the previous methanol: ethanol study (Chesnut et al., 2008).

The axial ratio (*c*/*a*) for the phase of Ti-6Al-4V (at ambient conditions) was measured to be 1.602, which is slightly higher than that reported for pure Ti (1.587), but is an expected result due to the presence of the stabiliser Al (possessing a smaller atomic radius than Ti) in the alloy.

For the phase of Ti-6Al-4V, in a methanol: ethanol PTM, we found the axial ratio to be almost constant between 1.600 and 1.602 up to 42 GPa, see figure 5. The measured *c*/*a* ratios for the phase are also effectively constant between 34 GPa and 74 GPa, varying between 0.616 and 0.617. We find similar results for a loading of Ti-6Al-4V in a neon PTM. These measurements are in broad agreement with those reported by Errandonea et al. for pure Ti

Titanium Alloys at Extreme Pressure Conditions 77

(Errandonea et al., 2005). We also include for reference in figure 5 our axial ratio results for nano-Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol: ethanol PTM (Velisavljevic & Chesnut, 2007) (see section 3.2). The *c*/*a* ratio for nano-Ti in the phase had a steady value of 0.612 initially, and as the pressure was increased above 20 GPa, this value increased slightly to 0.626 and then levelled off above 80 GPa. Under compression, the phase of Ti-6Al-4V in a methanol: ethanol PTM was observed to be stable to ~ 115 GPa (the

We observed, in further RT volume compression experiments of Ti-6Al-4V embedded in neon and mineral oil PTMs, a gradual transformation from the phase to the body-centredcubic β phase (space group Im3m) (Tegner et al., 2010 & Tegner et al., n.d.). For both neon and mineral oil PTMs, the transformation is completed between 115 GPa and 125 GPa. The β phase is formed by the splitting of the alternating (001) plane along the *c* axis of the phase into two (111) planes of the β phase (Xia et al., 1990a). All the β phase peaks are therefore contained in the diffraction pattern (that is, the peaks are coincident). With no detectable volume change from ω β, it was not possible to ascertain over what pressure range both phases coexisted. The disappearance of the phase (001), (210) and (002) peaks was the only indication that a solid-solid phase transformation had occurred. We observed the β phase of Ti-6Al-4V to be stable to at least 221 GPa (the pressure limit of the experiment) in a mineral oil PTM (Tegner et al., 2011) and stable to 130 GPa (the pressure limit of the experiment) in

The *P-V* plot for Ti-6Al-4V in a mineral oil PTM is shown in figure 6. No intermediate phases were observed in this experiment. For comparison, we include the *P*-*V* data for nano-Ti (Velisavljevic, n.d.), Ti-6Al-4V (Chesnut & Velisavljevic, 2008) and (see section 3) Ti-Beta-21S (Velisavljevic & Chesnut, 2007). Verma et al. proposed that the observations of the phase in Ti DAC experiments were likely caused by the presence of non-hydrostatic stresses and that the transition sequence ω β is thermodynamically preferable, with the intermediate phase existing over a range of at least 10 GPa (Verma et al., 2007). In our Ti-6Al-4V study, we collected ADXD patterns with pressure steps ~ 5 GPa and found no evidence for an intermediate phase in Ti-6Al-4V. Nor was there any evidence for an orthorhombic phase. The Ti-6Al-4V sample was embedded in a PTM (albeit the very nonhydrostatic mineral oil and neon PTMs above 100 GPa). Our own calculations predict the transformation pathway for Ti-6Al-4V to be ω β, with ω occurring at 24 GPa and ω β at ~ 105 GPa, in agreement with experiment (Tegner et al., 2010). Above 110 GPa we find that the orthorhombic and phases relax to the cubic β phase under hydrostatic conditions. As far as we know, all Ti DAC experiments compressed above 100 GPa have had samples compressed in the absence of a PTM (Vohra et al., 2001; Akahama et al., 2001;

In figure 7, an integrated ADXD pattern is shown of Ti-6Al-4V in a neon PTM, collected at 129 GPa (Tegner et al., n.d.). The peaks are indexed as the β-phase reflections (110), (200), (211) and (220) and face-centred cubic (fcc) neon (111), (002) and (220). The pressures were calculated using a previous static high pressure EOS study of neon (Dewaele et al., 2008). We performed a Rietveld analysis of our diffraction patterns to confirm the crystal structure

pressure limit in this experiment).

the neon PTM (Tegner et al., n.d.).

Velisavljevic et al., n.d.).

of Ti-6Al-4V to be the β phase.

Fig. 4. The *P-V* plot for Ti-6Al-4V in a methanol: ethanol PTM (Chesnut et al., 2008 & Tegner et al., 2011) and also for an unspecified PTM (Halevy et al., 2010).

Fig. 5. The axial ratios (*c/a*) for the and phases of Ti-6Al-4V in a methanol: ethanol PTM (Tegner et al., n.d.), Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol: ethanol PTM (Velisavljevic & Chesnut, 2007).

Fig. 4. The *P-V* plot for Ti-6Al-4V in a methanol: ethanol PTM (Chesnut et al., 2008 & Tegner

Fig. 5. The axial ratios (*c/a*) for the and phases of Ti-6Al-4V in a methanol: ethanol PTM (Tegner et al., n.d.), Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol:

et al., 2011) and also for an unspecified PTM (Halevy et al., 2010).

ethanol PTM (Velisavljevic & Chesnut, 2007).

(Errandonea et al., 2005). We also include for reference in figure 5 our axial ratio results for nano-Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol: ethanol PTM (Velisavljevic & Chesnut, 2007) (see section 3.2). The *c*/*a* ratio for nano-Ti in the phase had a steady value of 0.612 initially, and as the pressure was increased above 20 GPa, this value increased slightly to 0.626 and then levelled off above 80 GPa. Under compression, the phase of Ti-6Al-4V in a methanol: ethanol PTM was observed to be stable to ~ 115 GPa (the pressure limit in this experiment).

We observed, in further RT volume compression experiments of Ti-6Al-4V embedded in neon and mineral oil PTMs, a gradual transformation from the phase to the body-centredcubic β phase (space group Im3m) (Tegner et al., 2010 & Tegner et al., n.d.). For both neon and mineral oil PTMs, the transformation is completed between 115 GPa and 125 GPa. The β phase is formed by the splitting of the alternating (001) plane along the *c* axis of the phase into two (111) planes of the β phase (Xia et al., 1990a). All the β phase peaks are therefore contained in the diffraction pattern (that is, the peaks are coincident). With no detectable volume change from ω β, it was not possible to ascertain over what pressure range both phases coexisted. The disappearance of the phase (001), (210) and (002) peaks was the only indication that a solid-solid phase transformation had occurred. We observed the β phase of Ti-6Al-4V to be stable to at least 221 GPa (the pressure limit of the experiment) in a mineral oil PTM (Tegner et al., 2011) and stable to 130 GPa (the pressure limit of the experiment) in the neon PTM (Tegner et al., n.d.).

The *P-V* plot for Ti-6Al-4V in a mineral oil PTM is shown in figure 6. No intermediate phases were observed in this experiment. For comparison, we include the *P*-*V* data for nano-Ti (Velisavljevic, n.d.), Ti-6Al-4V (Chesnut & Velisavljevic, 2008) and (see section 3) Ti-Beta-21S (Velisavljevic & Chesnut, 2007). Verma et al. proposed that the observations of the phase in Ti DAC experiments were likely caused by the presence of non-hydrostatic stresses and that the transition sequence ω β is thermodynamically preferable, with the intermediate phase existing over a range of at least 10 GPa (Verma et al., 2007). In our Ti-6Al-4V study, we collected ADXD patterns with pressure steps ~ 5 GPa and found no evidence for an intermediate phase in Ti-6Al-4V. Nor was there any evidence for an orthorhombic phase. The Ti-6Al-4V sample was embedded in a PTM (albeit the very nonhydrostatic mineral oil and neon PTMs above 100 GPa). Our own calculations predict the transformation pathway for Ti-6Al-4V to be ω β, with ω occurring at 24 GPa and ω β at ~ 105 GPa, in agreement with experiment (Tegner et al., 2010). Above 110 GPa we find that the orthorhombic and phases relax to the cubic β phase under hydrostatic conditions. As far as we know, all Ti DAC experiments compressed above 100 GPa have had samples compressed in the absence of a PTM (Vohra et al., 2001; Akahama et al., 2001; Velisavljevic et al., n.d.).

In figure 7, an integrated ADXD pattern is shown of Ti-6Al-4V in a neon PTM, collected at 129 GPa (Tegner et al., n.d.). The peaks are indexed as the β-phase reflections (110), (200), (211) and (220) and face-centred cubic (fcc) neon (111), (002) and (220). The pressures were calculated using a previous static high pressure EOS study of neon (Dewaele et al., 2008). We performed a Rietveld analysis of our diffraction patterns to confirm the crystal structure of Ti-6Al-4V to be the β phase.

Titanium Alloys at Extreme Pressure Conditions 79

The observation of the ω β transformation pathway at RT for Ti-6Al-4V suggests that the slope of the -β phase boundary (see boundary in figure 2) is negative at high pressures, or that there are two separate areas of β phase separated by the phase (see for example,

In cases where a sufficient amount of alloying elements are introduced, known as Mo equivalent (Moeq), alloys with up to 50% β phase can be recovered after temperature treatment (Bania, 1993). The large β phase concentration can have a significant effect on the pseudo-elastic response (Zhou et al., 2004), while also effecting structural phase stability at high pressures. One example is the Mo rich Ti-Beta-21S (also known as TIMETAL21S®) alloy, which is a β-stabilized alloy with 15 wt.% Mo, 3 wt.% Al, 2.7 wt.% Nb, 0.2 wt% Si, and the remainder made up by Ti. The standard "solution-treated-and-aged" (STA) heat treatment, in which the samples were heated to 1098 K, held for 30 minutes, cooled to RT (air-cooled equivalent rate), subsequently heated to 828 K, held for 8 hours, and again cooled to RT, produced a sample resulting in a mixture of β phases, with 29% - 42% in the

An initial ADXD pattern collected at ambient conditions clearly shows the sample is composed of both the and β phases (Velisavljevic & Chesnut, 2007), as shown in figure 8. A close examination of the diffraction pattern does not indicate the existence of any peaks that could not be attributed to either the or β phase, and in particular, there is no evidence of phase in the sample. Cold compression of the sample in a methanol: ethanol PTM, in a DAC, shows very little change at low pressure. Up to ~ 11 GPa both and β phases appear to be stable, including the axial ratio of the phase, which remains fairly constant near the initial value of *c0*/*a0* = 1.601. Comparison of the axial ratio of the phase of Ti-Beta-21S and the *c0*/*a0* value of 1.602 reported for the phase of Ti-6Al-4V indicates that the inclusion of Mo, which has a larger atomic radius than Ti, does not have a significant effect on the initial axial ratio. However, the measured volume *V0* = 17.912 Å3/atom (*a0* = 2.956 ± 0.002 Å and *c0* = 4.734 ± 0.025 Å) for the phase of Ti-Beta-21S is larger than both the *V0* = 17.355 Å3/atom (*a0* = 3.262 ± 0.004 Å) measured for the β phase of Ti-Beta-21S and values reported for Ti-6Al-4V, which would suggest that a significant amount of alloying elements are still present in the phase portions of the sample. Above 11 GPa, anisotropic compression of the phase is observed, which leads to a steady decrease in the *c*/*a* ratio down to 1.568 at 36 GPa, as shown in figure 5. A sudden increase in the axial ratio up to 1.613 at 44 GPa, followed by a steady value up to 67 GPa is then observed. Over this same pressure range, besides a steady decrease in volume, no significant changes are observed for the β phase. The sudden change in the axial ratio in the 36-44 GPa region could denote a potential first order isostructural phase transition. Over the same pressure region of 36-44 GPa, no evidence of

**3.2 Absence of the phase in the Ti-beta-21S alloy at high pressure** 

phase (Velisavljevic & Chesnut, 2007; Honnell et al., 2007).

appearance of the phase or any other new phases could be detected.

With pressure increased to 58 GPa, the sample remains stable, as a mixture of β phases. However, above this pressure a relative change in the measured peak intensities between the two phases indicates the onset of a structural phase transition (Velisavljevic & Chesnut, 2007). A comparison of the intensity change of the (102) -phase peak and the

Xia, 1990a).

Fig. 6. The *P*-*V* plot for Ti-6Al-4V in a mineral oil PTM (Tegner et al., n.d.), Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol: ethanol PTM (Velisavljevic & Chesnut, 2007).

Fig. 7. An integrated ADXD pattern and Rietveld analysis for β (bcc) phase Ti-6Al-4V at 129 GPa and embedded in a neon PTM (Tegner et al., n.d.). The neon was used as both a PTM and pressure marker.

Fig. 6. The *P*-*V* plot for Ti-6Al-4V in a mineral oil PTM (Tegner et al., n.d.), Ti with no PTM (Velisavljevic et al., n.d.) and Ti-Beta-21S in a methanol: ethanol PTM (Velisavljevic &

Fig. 7. An integrated ADXD pattern and Rietveld analysis for β (bcc) phase Ti-6Al-4V at 129 GPa and embedded in a neon PTM (Tegner et al., n.d.). The neon was used as both a PTM

Chesnut, 2007).

and pressure marker.

The observation of the ω β transformation pathway at RT for Ti-6Al-4V suggests that the slope of the -β phase boundary (see boundary in figure 2) is negative at high pressures, or that there are two separate areas of β phase separated by the phase (see for example, Xia, 1990a).

#### **3.2 Absence of the phase in the Ti-beta-21S alloy at high pressure**

In cases where a sufficient amount of alloying elements are introduced, known as Mo equivalent (Moeq), alloys with up to 50% β phase can be recovered after temperature treatment (Bania, 1993). The large β phase concentration can have a significant effect on the pseudo-elastic response (Zhou et al., 2004), while also effecting structural phase stability at high pressures. One example is the Mo rich Ti-Beta-21S (also known as TIMETAL21S®) alloy, which is a β-stabilized alloy with 15 wt.% Mo, 3 wt.% Al, 2.7 wt.% Nb, 0.2 wt% Si, and the remainder made up by Ti. The standard "solution-treated-and-aged" (STA) heat treatment, in which the samples were heated to 1098 K, held for 30 minutes, cooled to RT (air-cooled equivalent rate), subsequently heated to 828 K, held for 8 hours, and again cooled to RT, produced a sample resulting in a mixture of β phases, with 29% - 42% in the phase (Velisavljevic & Chesnut, 2007; Honnell et al., 2007).

An initial ADXD pattern collected at ambient conditions clearly shows the sample is composed of both the and β phases (Velisavljevic & Chesnut, 2007), as shown in figure 8. A close examination of the diffraction pattern does not indicate the existence of any peaks that could not be attributed to either the or β phase, and in particular, there is no evidence of phase in the sample. Cold compression of the sample in a methanol: ethanol PTM, in a DAC, shows very little change at low pressure. Up to ~ 11 GPa both and β phases appear to be stable, including the axial ratio of the phase, which remains fairly constant near the initial value of *c0*/*a0* = 1.601. Comparison of the axial ratio of the phase of Ti-Beta-21S and the *c0*/*a0* value of 1.602 reported for the phase of Ti-6Al-4V indicates that the inclusion of Mo, which has a larger atomic radius than Ti, does not have a significant effect on the initial axial ratio. However, the measured volume *V0* = 17.912 Å3/atom (*a0* = 2.956 ± 0.002 Å and *c0* = 4.734 ± 0.025 Å) for the phase of Ti-Beta-21S is larger than both the *V0* = 17.355 Å3/atom (*a0* = 3.262 ± 0.004 Å) measured for the β phase of Ti-Beta-21S and values reported for Ti-6Al-4V, which would suggest that a significant amount of alloying elements are still present in the phase portions of the sample. Above 11 GPa, anisotropic compression of the phase is observed, which leads to a steady decrease in the *c*/*a* ratio down to 1.568 at 36 GPa, as shown in figure 5. A sudden increase in the axial ratio up to 1.613 at 44 GPa, followed by a steady value up to 67 GPa is then observed. Over this same pressure range, besides a steady decrease in volume, no significant changes are observed for the β phase. The sudden change in the axial ratio in the 36-44 GPa region could denote a potential first order isostructural phase transition. Over the same pressure region of 36-44 GPa, no evidence of appearance of the phase or any other new phases could be detected.

With pressure increased to 58 GPa, the sample remains stable, as a mixture of β phases. However, above this pressure a relative change in the measured peak intensities between the two phases indicates the onset of a structural phase transition (Velisavljevic & Chesnut, 2007). A comparison of the intensity change of the (102) -phase peak and the

Titanium Alloys at Extreme Pressure Conditions 81

Chesnut & Velisavljevic measured a 0.25% smaller value, V*0* = 17.208 Å3/atom (Chesnut &

A *P*-*T* phase diagram for Ti (Errandonea et al., 2001; Errandonea et al., 2005; Pecker et al., 2005; Zhang et al., 2008; Mei et al, 2009; Hu et al., 2010) and Ti-6Al-4V (Chesnut et al., 2008, Tegner et al., 2010) summarising the current state of knowledge of the phase relations of these systems up to 125 GPa is shown in figure 9. There is good agreement between experiment and theory for the location of the Ti phase boundary, and also the melt curve, but the location and slope of the β boundary is still in dispute and requires more study for clarification. Phase stability and the effects of anisotropic stresses on the transition in Ti is an issue that also requires more attention at high temperature. Very little is known about Ti alloys at high pressure, and even less at high pressure and high temperature. In figure 9, we suggest possible phase boundaries for Ti-6Al-4V as a dashed blue line. We are unsure about the exact location, or slope even, of the and β

Fig. 9. The combined *P-T* phase diagram of Ti (Errandonea et al., 2001; Errandonea et al., 2005; Pecker et al., 2005; Zhang et al., 2008; Mei et al, 2009; Hu et al., 2010) and Ti-6Al-4V (Chesnut et al., 2008, Tegner et al., 2010). The possible phase boundaries for the +β alloy Ti-

6Al-4V (wt%) are suggested by us as the dashed blue line.

Velisavljevic, 2007).

boundary for Ti-6Al-4V at high temperature.

**4. Discussion** 

(200) β-phase peak indicates a steady disappearance of the phase and a transition of the sample to 100% β phase above 67 GPa. The ADXD patterns collected at 18 GPa and 48 GPa in figure 8 show the mixture of the β phases. At 71 GPa in figure 8, only β phase peaks are now present. From the experimental data, a Vinet EOS fit returned values of *K* = 119 GPa and *K'* = 3.4 for the phase and *K* = 109 GPa and *K'* = 3.8 for the β phase (Honnell et al., 2007), and for a Birch-Murnaghan EOS fit, values of *K* = 117 GPa and *K'* = 3.4 for the phase and *K* = 110 GPa and *K'* = 3.7 for the β phase (Velisavljevic & Chesnut, 2007) were obtained. Although, as previously mentioned, there is a volume difference observed between the phase of Ti-Beta-21S and Ti-6Al-4V, the overall compressibility and EOS of Ti-Beta-21S are in good agreement with the various EOS values generated for Ti-6Al-4V, as shown in figure 6.

Fig. 8. A stack of ADXD patterns showing structural change in Ti-Beta-21S with increasing pressure. Initially sample is composed of mixture +β phase and with pressure increase sample transforms completely to β phase – ω or other intermediate phases were not observed at any point. Additional peaks in ADXD spectra belong to Cu, which was used as a pressure marker.

For data collection, and the values reported here for Ti-6Al-4V and Ti-beta-21S, synchrotron source monochromatic X-ray beams were used. The image plate detectors available at the synchrotrons had pixel sizes of 100 μm2 and so the diffraction patterns were generated with a resolution ~Δd/d = 10-3. As a consequence, the uncertainties in the measured volume data were of the order of ~0.3%. This is consistent with the different values reported for the volume of Ti-6Al-4V at ambient conditions by the various authors. For example, Tegner et al. reported *V0* = 17.252 Å3/atom (Tegner et al., 2010) whereas Chesnut & Velisavljevic measured a 0.25% smaller value, V*0* = 17.208 Å3/atom (Chesnut & Velisavljevic, 2007).

#### **4. Discussion**

80 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

(200) β-phase peak indicates a steady disappearance of the phase and a transition of the sample to 100% β phase above 67 GPa. The ADXD patterns collected at 18 GPa and 48 GPa in figure 8 show the mixture of the β phases. At 71 GPa in figure 8, only β phase peaks are now present. From the experimental data, a Vinet EOS fit returned values of *K* = 119 GPa and *K'* = 3.4 for the phase and *K* = 109 GPa and *K'* = 3.8 for the β phase (Honnell et al., 2007), and for a Birch-Murnaghan EOS fit, values of *K* = 117 GPa and *K'* = 3.4 for the phase and *K* = 110 GPa and *K'* = 3.7 for the β phase (Velisavljevic & Chesnut, 2007) were obtained. Although, as previously mentioned, there is a volume difference observed between the phase of Ti-Beta-21S and Ti-6Al-4V, the overall compressibility and EOS of Ti-Beta-21S are in good agreement with the various EOS values generated for

Fig. 8. A stack of ADXD patterns showing structural change in Ti-Beta-21S with increasing pressure. Initially sample is composed of mixture +β phase and with pressure increase sample transforms completely to β phase – ω or other intermediate phases were not observed at any point. Additional peaks in ADXD spectra belong to Cu, which was used as

For data collection, and the values reported here for Ti-6Al-4V and Ti-beta-21S, synchrotron source monochromatic X-ray beams were used. The image plate detectors available at the synchrotrons had pixel sizes of 100 μm2 and so the diffraction patterns were generated with a resolution ~Δd/d = 10-3. As a consequence, the uncertainties in the measured volume data were of the order of ~0.3%. This is consistent with the different values reported for the volume of Ti-6Al-4V at ambient conditions by the various authors. For example, Tegner et al. reported *V0* = 17.252 Å3/atom (Tegner et al., 2010) whereas

Ti-6Al-4V, as shown in figure 6.

a pressure marker.

A *P*-*T* phase diagram for Ti (Errandonea et al., 2001; Errandonea et al., 2005; Pecker et al., 2005; Zhang et al., 2008; Mei et al, 2009; Hu et al., 2010) and Ti-6Al-4V (Chesnut et al., 2008, Tegner et al., 2010) summarising the current state of knowledge of the phase relations of these systems up to 125 GPa is shown in figure 9. There is good agreement between experiment and theory for the location of the Ti phase boundary, and also the melt curve, but the location and slope of the β boundary is still in dispute and requires more study for clarification. Phase stability and the effects of anisotropic stresses on the transition in Ti is an issue that also requires more attention at high temperature. Very little is known about Ti alloys at high pressure, and even less at high pressure and high temperature. In figure 9, we suggest possible phase boundaries for Ti-6Al-4V as a dashed blue line. We are unsure about the exact location, or slope even, of the and β boundary for Ti-6Al-4V at high temperature.

Fig. 9. The combined *P-T* phase diagram of Ti (Errandonea et al., 2001; Errandonea et al., 2005; Pecker et al., 2005; Zhang et al., 2008; Mei et al, 2009; Hu et al., 2010) and Ti-6Al-4V (Chesnut et al., 2008, Tegner et al., 2010). The possible phase boundaries for the +β alloy Ti-6Al-4V (wt%) are suggested by us as the dashed blue line.

Titanium Alloys at Extreme Pressure Conditions 83

LANS, LLC for the DOE-NNSA – this work was, in part, supported by the US DOE under contract # DE-AC52-06NA25396. HP-CAT is supported by CIW, CDAC, UNLV and LLNL through funding from DOE-NNSA, DOE-BES and NSF. APS is supported by DOE-BES

Ahuja, R.; Dubrovinsky, L.; Dubrovinskaia, N.; Osorio Guillen, J.M.; Mattessini, M.;

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Dewaele, A.; Datchi, F.; Loubeyre, P. & Mezouar, M. (2008). High pressure-high temperature

Errandonea, D.; Schwager, B.; Ditz, R.; Gessmann, C.; Boehler, R. & Ross, M. (2001).

Errandonea, D.; Meng, Y.; Somayazulu, M. & Häusermann, D. (2005). Pressure-induced α

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oxygen content on the to phase transformation and shock hardening of

of Ti-6Al-4V. *Proceedings of the American Physical Society Topical Group on Shock Compression of Condensed Matter – 2007*, pp. 27-30, ISBN 978-0735404694, Kohala

equations of state of neon and diamond. *Physical Review B*, Vol.77, pp. 094106-1-

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under Contract No. DE-AC02-06CH11357.

Warrendale, USA, February 1993

*Research*, Vol.57, pp. 227-286

**6. References** 

184102-4

094106-9

132104-4

By comparing our high pressure Ti-Beta-21S (Velisavljevic & Chesnut, 2007), Ti-6Al-4V (Chesnut et al., 2008; Tegner et al., 2010; Tegner et al., n.d.) and nano-Ti (Velisavljevic et al., n.d.) data, we observe that the addition of alloying elements can have a significant influence on the structural phase transition sequence in these metals at high pressure and temperature. For example, the presence of alloys and interstitial impurities in Ti-6Al-4V suppresses the onset of the phase transformation, thus ensuring the predominance of the -phase alloy over a much larger pressure range than exists for pure Ti, which is desirable in industrial and commercial applications. The main effect observed in Ti-Beta-21S is the complete suppression of the brittle phase at high pressures.

However, it is also important to recognize that changes in the electronic configuration cause changes in crystal structures. These changes can be induced either by an increase in pressure or an increase in the occupancy of the *d* bands. The structural trend exhibited by the 3*d*, 4*d*, and 5*d* transition metals is well known. The variation in the electronic configuration affects crystal structures and mechanical properties such as microstructures and dislocations. For industrial applications, it is the machineability and superior mechanical properties that are of paramount interest. Pressure allows us to measure the differences in crystal structures induced by changes in the electronic configuration.

Among the Group IV elements, similar structural changes occur. Group IV elements and their alloys apparently favor a transformation pathway β at high pressure, based on recent DAC experiments and theoretical calculations. The intermediate structures, δ and γ, appear to be metastable and are shear driven. Based on the experimental results one can conclude that high pressure structural phase transitions in Ti and Ti alloys are highly susceptible to loading conditions and stress distribution, as shown from experiments using various PMTs. However, stability of various phases can be controlled by other variables, such as alloying, which can change electronic structure by increasing/decreasing *d* band occupancy, inclusion of interstitial impurities, which help reduce shear deformation, and in some cases it appears that shear driven structural phase transitions can be controlled by varying sample grain size as well. Although these factors play a significant role in controlling structural phase transitions they appear to have only a slight effect on the initial compressibility (i.e. EOS) – measured phase EOS parameters for Ti-Beta-21S, Ti-6Al-4V, Ti, and nano-Ti are all relatively close with values of *K* = 115- 125 GPa and *K'* = 2.4-3.9. Overall it appears that with stress conditions, grain size, and presence of impurities, there is a systematic shift of the transition pressure *P* , as the transition pressure increases with improved hydrostaticity of the pressure environment and by grain size reduction.

#### **5. Acknowledgements**

SM would like to acknowledge the support of Professor Malcolm McMahon and Dr John Proctor of the Centre for Science at Extreme Conditions (CSEC), Edinburgh University, in collecting the Ti-6Al-4V data. This work was performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory in part under contract W-7405-Eng-48 and in part under Contract DE-AC52-07NA27344. LANL is operated by LANS, LLC for the DOE-NNSA – this work was, in part, supported by the US DOE under contract # DE-AC52-06NA25396. HP-CAT is supported by CIW, CDAC, UNLV and LLNL through funding from DOE-NNSA, DOE-BES and NSF. APS is supported by DOE-BES under Contract No. DE-AC02-06CH11357.

#### **6. References**

82 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

By comparing our high pressure Ti-Beta-21S (Velisavljevic & Chesnut, 2007), Ti-6Al-4V (Chesnut et al., 2008; Tegner et al., 2010; Tegner et al., n.d.) and nano-Ti (Velisavljevic et al., n.d.) data, we observe that the addition of alloying elements can have a significant influence on the structural phase transition sequence in these metals at high pressure and temperature. For example, the presence of alloys and interstitial impurities in Ti-6Al-4V suppresses the onset of the phase transformation, thus ensuring the predominance of the -phase alloy over a much larger pressure range than exists for pure Ti, which is desirable in industrial and commercial applications. The main effect observed in Ti-Beta-21S

However, it is also important to recognize that changes in the electronic configuration cause changes in crystal structures. These changes can be induced either by an increase in pressure or an increase in the occupancy of the *d* bands. The structural trend exhibited by the 3*d*, 4*d*, and 5*d* transition metals is well known. The variation in the electronic configuration affects crystal structures and mechanical properties such as microstructures and dislocations. For industrial applications, it is the machineability and superior mechanical properties that are of paramount interest. Pressure allows us to measure the differences in crystal structures induced by changes in the electronic

Among the Group IV elements, similar structural changes occur. Group IV elements and their alloys apparently favor a transformation pathway β at high pressure, based on recent DAC experiments and theoretical calculations. The intermediate structures, δ and γ, appear to be metastable and are shear driven. Based on the experimental results one can conclude that high pressure structural phase transitions in Ti and Ti alloys are highly susceptible to loading conditions and stress distribution, as shown from experiments using various PMTs. However, stability of various phases can be controlled by other variables, such as alloying, which can change electronic structure by increasing/decreasing *d* band occupancy, inclusion of interstitial impurities, which help reduce shear deformation, and in some cases it appears that shear driven structural phase transitions can be controlled by varying sample grain size as well. Although these factors play a significant role in controlling structural phase transitions they appear to have only a slight effect on the initial compressibility (i.e. EOS) – measured phase EOS parameters for Ti-Beta-21S, Ti-6Al-4V, Ti, and nano-Ti are all relatively close with values of *K* = 115- 125 GPa and *K'* = 2.4-3.9. Overall it appears that with stress conditions, grain size, and presence of impurities, there is a systematic shift of the transition pressure *P* , as the transition pressure increases with improved hydrostaticity of the pressure environment

SM would like to acknowledge the support of Professor Malcolm McMahon and Dr John Proctor of the Centre for Science at Extreme Conditions (CSEC), Edinburgh University, in collecting the Ti-6Al-4V data. This work was performed under the auspices of the U.S. Department of Energy by Lawrence Livermore National Laboratory in part under contract W-7405-Eng-48 and in part under Contract DE-AC52-07NA27344. LANL is operated by

is the complete suppression of the brittle phase at high pressures.

configuration.

and by grain size reduction.

**5. Acknowledgements** 


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**5** 

*Poland* 

**Hot Plasticity of Alpha Beta Alloys** 

Two phase titanium alloys most often are hot deformed, mainly by open die or close-die forging. Desired mechanical properties can be achieved in these alloys by development of proper microstructure in plastic working and heat treatment processes. Irreversible microstructural changes caused by deformation at the temperature in α+β↔β phase transformation range quite often cannot be eliminated or reduced by heat treatment and therefore required properties of products cannot be achieved (Bylica & Sieniawski, 1985; Lütjering, 1998; Zwicker, 1974). Some of the properties of titanium alloys, such as: high chemical affinity to oxygen, low thermal conductivity, high heat capacity and significant dependence of plastic flow resistance on strain rate, make it very difficult to obtain finished products having desired microstructure and properties by hot working. Differences in temperature across the material volume, which result from various deformation conditions (local strain and strain rate) lead to formation of zones having various phase composition (equilibrium α and β phases, martensitic phases α'(α")), morphology (equiaxial, lamellar, bimodal) and dispersion (fine- or coarse-grained) and therefore various mechanical properties

Obtaining desired microstructure of Ti-6Al-4V titanium alloy using plastic deformation in the α+β↔β phase transformation range is related to proper conditions selection taking into account plastic deformation, phase transformation, dynamic recovery and recrystallization effects (Ding et al., 2002; Kubiak, 2004; Kubiak & Sieniawski, 1998). Grain refinement can be achieved by including preliminary heat treatment in thermomechanical process. Final heat treatment operations are usually used for stabilization of microstructure (they restrict grain

Titanium alloys together with aluminium alloys belong to the largest group of superplastic materials used in industrial SPF. Their main advantages are good superplasticity combined with relatively high susceptibility to diffusion bonding. Among them two-phase α+β Ti-6Al-4V alloy has been the most popular for many years as it exhibits superplasticity even after application of conventional plastic working methods

**1. Introduction** 

(Kubiak & Sieniawski, 1998).

growth) (Motyka & Sieniawski, 2010).

(Sieniawski & Motyka, 2007).

Maciej Motyka, Krzysztof Kubiak, Jan Sieniawski and Waldemar Ziaja

> *Department of Materials Science, Rzeszow University of Technology*

Zhou, T.; Aindow, M; Alpay, S.P.; Blackburn, M.J.; Wu, M.H. (2004). Pseudo-elastic deformation behavior in a Ti/Mo-based alloy. *Scripta Materialia,* Vol.50, pp. 343– 348

### **Hot Plasticity of Alpha Beta Alloys**

Maciej Motyka, Krzysztof Kubiak, Jan Sieniawski and Waldemar Ziaja *Department of Materials Science, Rzeszow University of Technology Poland* 

#### **1. Introduction**

86 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Zhou, T.; Aindow, M; Alpay, S.P.; Blackburn, M.J.; Wu, M.H. (2004). Pseudo-elastic

348

deformation behavior in a Ti/Mo-based alloy. *Scripta Materialia,* Vol.50, pp. 343–

Two phase titanium alloys most often are hot deformed, mainly by open die or close-die forging. Desired mechanical properties can be achieved in these alloys by development of proper microstructure in plastic working and heat treatment processes. Irreversible microstructural changes caused by deformation at the temperature in α+β↔β phase transformation range quite often cannot be eliminated or reduced by heat treatment and therefore required properties of products cannot be achieved (Bylica & Sieniawski, 1985; Lütjering, 1998; Zwicker, 1974). Some of the properties of titanium alloys, such as: high chemical affinity to oxygen, low thermal conductivity, high heat capacity and significant dependence of plastic flow resistance on strain rate, make it very difficult to obtain finished products having desired microstructure and properties by hot working. Differences in temperature across the material volume, which result from various deformation conditions (local strain and strain rate) lead to formation of zones having various phase composition (equilibrium α and β phases, martensitic phases α'(α")), morphology (equiaxial, lamellar, bimodal) and dispersion (fine- or coarse-grained) and therefore various mechanical properties (Kubiak & Sieniawski, 1998).

Obtaining desired microstructure of Ti-6Al-4V titanium alloy using plastic deformation in the α+β↔β phase transformation range is related to proper conditions selection taking into account plastic deformation, phase transformation, dynamic recovery and recrystallization effects (Ding et al., 2002; Kubiak, 2004; Kubiak & Sieniawski, 1998). Grain refinement can be achieved by including preliminary heat treatment in thermomechanical process. Final heat treatment operations are usually used for stabilization of microstructure (they restrict grain growth) (Motyka & Sieniawski, 2010).

Titanium alloys together with aluminium alloys belong to the largest group of superplastic materials used in industrial SPF. Their main advantages are good superplasticity combined with relatively high susceptibility to diffusion bonding. Among them two-phase α+β Ti-6Al-4V alloy has been the most popular for many years as it exhibits superplasticity even after application of conventional plastic working methods (Sieniawski & Motyka, 2007).

Hot Plasticity of Alpha Beta Alloys 89

a) b) Fig. 2. Critical resolved shear stress as a function of: a) interstitial impurities content and temperature (Zwicker, 1974) b) temperature and slip plane in titanium with 0.1% interstitial

The number of twinning planes is higher in titanium than in other metals (except for zirconium) having hcp crystal structure. Twinning in single crystal titanium occurs in { 1102 }, { 1121 }, { 1122 }, { 1123 }, { 1124 } planes, while in polycrystalline titanium in { 1102 }, { 1121 }, { 1122 } planes (Fig. 3). The smallest value of critical stress for twinning occurs in { 1102 } plane. It was found that slip and twinning interact with each other. Twinning in { 1122 } plane hinders slip in that plane. Twinning is intensified by increase in deformation,

metal purity, grain size and by decrease in temperature (Churchman, 1955).

Fig. 3. Twinning planes in titanium (Balasubramanian & Anand, 2002)

impurities content (Levine, 1966)

#### **2. Deformation behaviour of titanium and its alloys**

Titanium has two allotropic forms: Ti with hexagonal close packed (hcp) crystal structure (up to 882.5°C) and Ti with body centered cubic (bcc) crystal structure (between 882.5 and 1662°C). Each of the allotropes exhibit different plasticity resulting from its crystal structure and different number of the slip systems (A. D. Mc Quillan & M. K. Mc Quillan, 1956; Zwicker, 1974).

Deformation of titanium, both at room and elevated temperature, occurs by slip and twinning (Tab. 1). Primary slip systems in titanium are: { 1100 } 1120 and { 1101 } 1120 (Fig. 1). Critical resolved shear stress value (crss) depends on deformation temperature, impurities content and slip system (Fig. 2). If the relative value of crss stress for { 1100 } 1120 slip system is set to 1, for { 1101 } 1120 system it equals to 1.75, and for {0001} 1120 system – 1.92. Coarse grained and single crystal titanium deforms in {0001} 1120 system, because stacking fault energy and atom packing density are larger in { 1100 } planes than in {0001} planes (Kajbyszew & Krajuchijn, 1967; Zwicker, 1974).


Table 1. Slip and twinning planes and critical resolved shear stress in metals with hcp crystal structure (Kajbyszew & Krajuchijn, 1967)

Fig. 1. Slip systems in titanium (Balasubramanian & Anand, 2002)

Titanium has two allotropic forms: Ti with hexagonal close packed (hcp) crystal structure (up to 882.5°C) and Ti with body centered cubic (bcc) crystal structure (between 882.5 and 1662°C). Each of the allotropes exhibit different plasticity resulting from its crystal structure and different number of the slip systems (A. D. Mc Quillan & M. K. Mc Quillan, 1956;

Deformation of titanium, both at room and elevated temperature, occurs by slip and twinning (Tab. 1). Primary slip systems in titanium are: { 1100 } 1120 and { 1101 } 1120 (Fig. 1). Critical resolved shear stress value (crss) depends on deformation temperature, impurities content and slip system (Fig. 2). If the relative value of crss stress for { 1100 } 1120 slip system is set to 1, for { 1101 } 1120 system it equals to 1.75, and for {0001} 1120 system – 1.92. Coarse grained and single crystal titanium deforms in {0001} 1120 system, because stacking fault energy and atom packing density are larger in { 1100 } planes than in

{ 1100 }

{0001} { 1100 }

Fig. 1. Slip systems in titanium (Balasubramanian & Anand, 2002)

Table 1. Slip and twinning planes and critical resolved shear stress in metals with hcp crystal

Critical resolved shear stress crss, MPa

0.34

4.5 5.1

{ <sup>1100</sup> } 39.2 { <sup>1103</sup> }

{0001} 62.1 { <sup>1122</sup> }

{ <sup>1100</sup> } 13.83 { <sup>1122</sup> }

10-15 { <sup>1102</sup> }

Twinning plane

{ 1102 } { 1101 }

{ 1121 }

{ 1121 }

{ 1102 }

**2. Deformation behaviour of titanium and its alloys** 

{0001} planes (Kajbyszew & Krajuchijn, 1967; Zwicker, 1974).

Metal c/a ratio Slip plane

Zinc 1.856 {0001}

Manganese 1.624

Titanium 1.587

structure (Kajbyszew & Krajuchijn, 1967)

Zwicker, 1974).

Fig. 2. Critical resolved shear stress as a function of: a) interstitial impurities content and temperature (Zwicker, 1974) b) temperature and slip plane in titanium with 0.1% interstitial impurities content (Levine, 1966)

The number of twinning planes is higher in titanium than in other metals (except for zirconium) having hcp crystal structure. Twinning in single crystal titanium occurs in { 1102 }, { 1121 }, { 1122 }, { 1123 }, { 1124 } planes, while in polycrystalline titanium in { 1102 }, { 1121 }, { 1122 } planes (Fig. 3). The smallest value of critical stress for twinning occurs in { 1102 } plane. It was found that slip and twinning interact with each other. Twinning in { 1122 } plane hinders slip in that plane. Twinning is intensified by increase in deformation, metal purity, grain size and by decrease in temperature (Churchman, 1955).

Fig. 3. Twinning planes in titanium (Balasubramanian & Anand, 2002)

Hot Plasticity of Alpha Beta Alloys 91

Plasticity of commercial pure (CP) titanium and single-phase alloys depends on type and content of impurities, alloying elements, temperature of deformation and strain rate. Increase in impurities and alloying elements content reduces plasticity due to solid solution

> Ti-0.3Fe-0.07C-0.04N-0.2O2-0.01H Ti-0.2Fe-0.05C-0.04N-0.1O2-0.008H

Ti-5Al-2.5Sn

Fig. 5. The influence of temperature on plasticity of CP titanium and single-phase alloys


Increase in temperature of deformation reduces plastic flow stress of the CP titanium and

Critical strain of CP titanium depends on temperature of deformation and strain rate

Ti-5Al-2.5Sn

Ti-0.3Fe-0.07C-0.04N-0.2O2-0.01H

Fig. 6. Dependence of plastic flow stress of CP titanium and alloys on deformation

Roughly 90% of about 70 grades of titanium alloys that are manufactured by conventional methods are two-phase martensitic or transition alloys. They exhibit high relative strength (UTS/), good creep resistance (up to 450°C), corrosion resistance in many environments, good weldability and formability. The most widely used representative of this group of alloys is Ti-6Al-4V showing good balance of mechanical and technological properties.

700 800 900 1000 1100 Deformation temperature, °C

strengthening (Fig. 5) (Glazunov & Mojsiejew, 1974).

0

100

200

300

Plastic flow stress p, MPa

400

500

Elongation

A5, %

(Glazunov & Mojsiejew, 1974)

single-phase alloys (Fig. 6).

temperature (Hadasik, 1979)

(Fig. 7).

Plastic deformation in titanium takes place by mechanisms characteristic of metals with bcc structure. In titanium following slip systems operate: {110} 111, {112} 111, {123} 111 along with twinning system {112} 111 (Balasubramanian & Anand, 2002; Wassermann & Grewen, 1962). Alloying elements affect plasticity of titanium to different degree. Cold strain of titanium (with high content of Mo, Nb or Ta) can exceed 90%. However addition of ruthenium or rhodium results in very large decrease in plasticity, rendering cold deformation of titanium practically impossible (Raub & Röschelb, 1963).

Titanium alloys can be classified according to their microstructure in particular state (e.g. after normalizing). It must be emphasized that this classification is questionable, because phase transformations in alloys with transition elements proceed so slowly, that very often the microstructure consistent with phase equilibrium diagram cannot be obtained at room temperature. According to widely accepted classification of the alloys in normalized state following types of titanium alloys can be distinguished (Fig. 4) (Glazunov & Kolachev, 1980):

	- not heat strengthened
	- heat strengthened due to metastable phases decomposition
	- strengthened by quenching
	- with increased plasticity after quenching
	- with mechanically unstable MN phase (decomposing under the stress)
	- with mechanically stable MS phase (not decomposing under the stress)
	- with thermodynamically stable TS phase.

Two transition types of alloys can also be distinguished:


Fig. 4. Classification of titanium alloys on the basis of stabilizers content: 1 – alloys, 2 – near alloys, 3 – martensitic + alloys, 4 – transition + alloys, 5 – near alloys, 6 – alloys (Glazunov & Kolachev, 1980)

Plastic deformation in titanium takes place by mechanisms characteristic of metals with bcc structure. In titanium following slip systems operate: {110} 111, {112} 111, {123} 111 along with twinning system {112} 111 (Balasubramanian & Anand, 2002; Wassermann & Grewen, 1962). Alloying elements affect plasticity of titanium to different degree. Cold strain of titanium (with high content of Mo, Nb or Ta) can exceed 90%. However addition of ruthenium or rhodium results in very large decrease in plasticity, rendering cold deformation of titanium practically impossible (Raub & Röschelb, 1963).

Titanium alloys can be classified according to their microstructure in particular state (e.g. after normalizing). It must be emphasized that this classification is questionable, because phase transformations in alloys with transition elements proceed so slowly, that very often the microstructure consistent with phase equilibrium diagram cannot be obtained at room temperature. According to widely accepted classification of the alloys in normalized state following types of titanium alloys can be distinguished (Fig. 4) (Glazunov & Kolachev, 1980):



1. near alloys – which contain up to 5% of stabilizers (microstructure is composed of

2. near alloys – their properties match the properties of + alloys with high volume fraction of phase (the microstructure after solutionizing is composed of metastable <sup>M</sup>

Fig. 4. Classification of titanium alloys on the basis of stabilizers content: 1 – alloys, 2 – near alloys, 3 – martensitic + alloys, 4 – transition + alloys, 5 – near alloys,

1. alloys

2. + alloys

3. alloys

phase).



6 – alloys (Glazunov & Kolachev, 1980)



phase and small amount, about 3-5%, of phase),

Plasticity of commercial pure (CP) titanium and single-phase alloys depends on type and content of impurities, alloying elements, temperature of deformation and strain rate. Increase in impurities and alloying elements content reduces plasticity due to solid solution strengthening (Fig. 5) (Glazunov & Mojsiejew, 1974).

Fig. 5. The influence of temperature on plasticity of CP titanium and single-phase alloys (Glazunov & Mojsiejew, 1974)

Increase in temperature of deformation reduces plastic flow stress of the CP titanium and single-phase alloys (Fig. 6).

Critical strain of CP titanium depends on temperature of deformation and strain rate (Fig. 7).

Fig. 6. Dependence of plastic flow stress of CP titanium and alloys on deformation temperature (Hadasik, 1979)

Roughly 90% of about 70 grades of titanium alloys that are manufactured by conventional methods are two-phase martensitic or transition alloys. They exhibit high relative strength (UTS/), good creep resistance (up to 450°C), corrosion resistance in many environments, good weldability and formability. The most widely used representative of this group of alloys is Ti-6Al-4V showing good balance of mechanical and technological properties.

Hot Plasticity of Alpha Beta Alloys 93

Fig. 9. The schematic diagram of plastic work and heat treatment processes of two phase

The processes of hot working and heat treatment of two-phase titanium alloys, e.g. Ti-6Al-4V, allow to obtain various types of microstructure (Fig.10) (Ezugwu & Wang, 1997; Kubiak

a) b) c)

d) e)

d) lamellar, e) bi-modal (Kubiak & Sieniawski, 1998)

Fig. 10. Microstructure of two-phase Ti-6Al-4V alloy: a) martensitic, b) globular, c) necklace,

titanium alloys (Peters et al., 1983)




& Sieniawski, 1998):

Fig. 7. The effect of temperature of deformation and strain rate on the critical strain of CP titanium (Hadasik, 1979)

The basic technological processes enabling final product manufacturing and development of mechanical properties of titanium alloys are hot working and heat treatment. Application of cold working is limited to the operations of bending (sheets, flat bars, tubes and bars) and shallow drawing of sheets allowing to obtain large elements. Bulk cold forming is not used due to high resistance of titanium and its alloys to plastic flow. (Fig. 8) (Lee & Lin, 1997).

Fig. 8. The influence of strain rate on plastic flow stress of cold worked Ti-6Al-4V alloy (Lee & Lin, 1997)

Plasticity, microstructure and mechanical properties of titanium alloys depend on hot working conditions (Peters et al., 1983; Brooks, 1996):


Application of all or selected operations of hot working and heat treatment (Fig. 9) allows to vary to a large extent the microstructure e.g. morphology and dispersion of phases in + titanium alloys (Fig. 10).

Fig. 7. The effect of temperature of deformation and strain rate on the critical strain of CP

Deformation temperature, °C 700 800 900

0.1

1

10

Critical strain, <sup>f</sup>

100

Fig. 8. The influence of strain rate on plastic flow stress of cold worked Ti-6Al-4V alloy


working conditions (Peters et al., 1983; Brooks, 1996):



Plasticity, microstructure and mechanical properties of titanium alloys depend on hot

Application of all or selected operations of hot working and heat treatment (Fig. 9) allows to vary to a large extent the microstructure e.g. morphology and dispersion of phases in +

The basic technological processes enabling final product manufacturing and development of mechanical properties of titanium alloys are hot working and heat treatment. Application of cold working is limited to the operations of bending (sheets, flat bars, tubes and bars) and shallow drawing of sheets allowing to obtain large elements. Bulk cold forming is not used due to high resistance of titanium and its alloys to plastic flow. (Fig. 8) (Lee & Lin, 1997).

0.1 1 10 100 Strain rate • , s-1

1000 1100

titanium (Hadasik, 1979)

(Lee & Lin, 1997)


titanium alloys (Fig. 10).

Fig. 9. The schematic diagram of plastic work and heat treatment processes of two phase titanium alloys (Peters et al., 1983)

The processes of hot working and heat treatment of two-phase titanium alloys, e.g. Ti-6Al-4V, allow to obtain various types of microstructure (Fig.10) (Ezugwu & Wang, 1997; Kubiak & Sieniawski, 1998):


Fig. 10. Microstructure of two-phase Ti-6Al-4V alloy: a) martensitic, b) globular, c) necklace, d) lamellar, e) bi-modal (Kubiak & Sieniawski, 1998)

Hot Plasticity of Alpha Beta Alloys 95

High sensitivity to strain rate is a characteristic feature of two-phase titanium alloys. The coefficient of strain rate sensitivity depends on temperature of deformation and strain rate (Fig. 13), grain size (Fig. 14), strain magnitude (Fig. 15) and morphology of phase

Fig. 13. The dependence of strain rate sensitivity factor *m* on temperature (volume fraction

Fig. 14. The dependence of strain rate sensitivity factor *m* on grain size of phase for

Ti-6Al-4V alloy (Semiatin et al., 1998)

of phase) and strain rate for Ti-6Al-4V alloy (Semiatin et al., 1998)

constituents (Fig. 16) (Semiatin et al., 1998).

#### **3. The influence of deformation conditions and morphology of phases on the plasticity of + titanium alloys**

Hot deformation behaviour of two-phase titanium alloys depends on chemical and phase composition, stereological parameters of microstructure and process conditions (deformation temperature, strain rate, stress and strain distribution). They exhibit analogous dependences of the plastic flow stress on temperature and strain rate like other metals alloys. Increase in strain rate raises plastic flow stress while increase in deformation temperature reduces it (Sakai, 1995).

Increase in temperature of deformation of two-phase titanium alloys reduces plastic flow stress (Fig. 11) more effectively in + field than in phase field, as a result of change in volume fraction of and phases (Fig. 12) (Ding et al., 2002; Sheppard & Norley, 1988).

Fig. 11. The effect of temperature of deformation and strain rate on plastic flow stress for Ti-6Al-4V alloy: a) torsion test (Sheppard & Norley, 1988), b) compression test (Ding et. al., 2002)

Fig. 12. The effect of temperature on volume fraction of and phases in Ti-6Al-4V alloy during torsion test (Sheppard & Norley, 1988)

**3. The influence of deformation conditions and morphology of phases on the** 

Hot deformation behaviour of two-phase titanium alloys depends on chemical and phase composition, stereological parameters of microstructure and process conditions (deformation temperature, strain rate, stress and strain distribution). They exhibit analogous dependences of the plastic flow stress on temperature and strain rate like other metals alloys. Increase in strain rate raises plastic flow stress while increase in deformation

Increase in temperature of deformation of two-phase titanium alloys reduces plastic flow stress (Fig. 11) more effectively in + field than in phase field, as a result of change in volume fraction of and phases (Fig. 12) (Ding et al., 2002; Sheppard & Norley, 1988).

a) b) Fig. 11. The effect of temperature of deformation and strain rate on plastic flow stress for Ti-6Al-4V alloy: a) torsion test (Sheppard & Norley, 1988), b) compression test (Ding et. al., 2002)

Fig. 12. The effect of temperature on volume fraction of and phases in Ti-6Al-4V alloy

during torsion test (Sheppard & Norley, 1988)

**plasticity of + titanium alloys** 

temperature reduces it (Sakai, 1995).

High sensitivity to strain rate is a characteristic feature of two-phase titanium alloys. The coefficient of strain rate sensitivity depends on temperature of deformation and strain rate (Fig. 13), grain size (Fig. 14), strain magnitude (Fig. 15) and morphology of phase constituents (Fig. 16) (Semiatin et al., 1998).

Fig. 13. The dependence of strain rate sensitivity factor *m* on temperature (volume fraction of phase) and strain rate for Ti-6Al-4V alloy (Semiatin et al., 1998)

Fig. 14. The dependence of strain rate sensitivity factor *m* on grain size of phase for Ti-6Al-4V alloy (Semiatin et al., 1998)

Hot Plasticity of Alpha Beta Alloys 97

These processes occur simultaneously during deformation and have an impact both on

Fig. 18. The texture diagram – pole figure (0002) – of two-phase titanium alloys for various

After deformation in the + temperature range and air cooling following microstructure of




Increase in degree or rate of deformation may lead to local increase in temperature even above finish temperature of +↔ transformation and development of martensitic '(") phases or colonies of lamellar and phases with stereological parameters depending on

Upon deformation of + titanium alloys above the finish temperature of +↔ transformation dynamic recrystallization occurs. New grains nucleate at primary grain boundaries and then grow until complete ring arranged along these grain boundaries is formed. New nuclei form also inside formerly recrystallized grains in the ring and the process repeats until new grains completely fill up the 'old' grain (Fig. 20) (Ding et al.,

temperature at the cooling rate slightly lower than critical (Fig. 19-1),

texture (Fig. 18) and morphology of phases (Lütjering, 1998).

modes and temperature of deformation (Lütjering, 1998)

phase can be obtained (Fig. 19) (Kubiak & Sieniawski, 1998):

temperature of +↔ transformation (Fig. 19-2),

of +↔ transformation (Fig. 19-6).

cooling rate.

2002).

Fig. 15. The dependence of strain rate sensitivity factor *m* on the strain magnitude for Ti-6Al-4V alloy (Ghosh & Hamilton, 1979)

Fig. 16. The dependence of strain rate sensitivity factor *m* on microstructure morphology (type of heat treatment) for hot-rolled sheets of Ti-6Al-4V alloy (Semiatin et al., 1998)

Volume fraction of and phases affects the value of the coefficient of strain rate sensitivity *m*. For two-phase titanium alloys it reaches maximum value when Vv=Vv 50%, at the temperature close to start temperature of + phase transformation and strain rate in the range of = 4·10-5 – 1·10-3 s-1. Grain refinement leads to increase in *m* value and growth of strain magnitude reduces it (Ghosh & Hamilton, 1979).

Deformation of titanium alloys in two-phase range leads to distortion of and grains, fragmentation of grains and their globularization. As a result of these processes (Fig. 17) elongated phase grains develop with particular orientation which are arranged along direction of maximum deformation (Seshacharyulu et al., 2002).

Fig. 17. The deformation process in + field – fragmentation and globularization of phase lamellae (Seshacharyulu et al., 2002)

Fig. 15. The dependence of strain rate sensitivity factor *m* on the strain magnitude for Ti-6Al-

Fig. 16. The dependence of strain rate sensitivity factor *m* on microstructure morphology (type of heat treatment) for hot-rolled sheets of Ti-6Al-4V alloy (Semiatin et al., 1998)

Volume fraction of and phases affects the value of the coefficient of strain rate sensitivity *m*. For two-phase titanium alloys it reaches maximum value when Vv=Vv 50%, at the temperature close to start temperature of + phase transformation and strain rate in the

Deformation of titanium alloys in two-phase range leads to distortion of and grains, fragmentation of grains and their globularization. As a result of these processes (Fig. 17) elongated phase grains develop with particular orientation which are arranged along

Fig. 17. The deformation process in + field – fragmentation and globularization of phase

= 4·10-5 – 1·10-3 s-1. Grain refinement leads to increase in *m* value and growth of

4V alloy (Ghosh & Hamilton, 1979)

range of

lamellae (Seshacharyulu et al., 2002)

strain magnitude reduces it (Ghosh & Hamilton, 1979).

direction of maximum deformation (Seshacharyulu et al., 2002).

These processes occur simultaneously during deformation and have an impact both on texture (Fig. 18) and morphology of phases (Lütjering, 1998).

Fig. 18. The texture diagram – pole figure (0002) – of two-phase titanium alloys for various modes and temperature of deformation (Lütjering, 1998)

After deformation in the + temperature range and air cooling following microstructure of phase can be obtained (Fig. 19) (Kubiak & Sieniawski, 1998):


Increase in degree or rate of deformation may lead to local increase in temperature even above finish temperature of +↔ transformation and development of martensitic '(") phases or colonies of lamellar and phases with stereological parameters depending on cooling rate.

Upon deformation of + titanium alloys above the finish temperature of +↔ transformation dynamic recrystallization occurs. New grains nucleate at primary grain boundaries and then grow until complete ring arranged along these grain boundaries is formed. New nuclei form also inside formerly recrystallized grains in the ring and the process repeats until new grains completely fill up the 'old' grain (Fig. 20) (Ding et al., 2002).

Hot Plasticity of Alpha Beta Alloys 99

Increase in deformation rate decreases the size of recrystallized grains (Fig. 21). The

Fig. 21. The effect of strain rate on average diameter of primary phase grains in Ti 6Al-4V

However presented model of dynamic recrystallization (Fig. 20) does not take into account the change of shape of primary grains and deformation of recrystallized grains. Thus it can be stated that it describes process of metadynamic recrystallization. Upon cooling after dynamic or metadynamic recrystallization following phases can be formed within grains (depending on cooling rate): colonies of and lamellae, equiaxed and grains,

The character of flow curves obtained during sequential deformation of two phase titanium alloys at the temperature in field confirms occurrence of dynamic processes of microstructure recovery. Increase in time of metadynamic recrystallization reduces strengthening. The character of the curves describing volume fraction of recrystallized phase confirms the influence of chemical composition on recrystallization kinetics (Fig. 22). Increasing content of alloying elements leads to decrease in rate of recrystallization. Time for recrystallization of 50% of phase is equal t0,5 3.5 s for Ti-6Al-5Mo-5V-2Cr-1Fe alloy and

alloy with globular initial microstructure (Seshacharyulu et al., 1999)

martensitic '(") phases or the mixture of them.

t0,5 1.5 s for Ti-6Al-4V alloy.

, induced by deformation, can be described for Ti-6Al-4V alloy by following equations,

and average thickness of lamellae –

= 1954.3 *Z*-0,172 (m) (Seshacharyulu et al., 2002)

= 1406.4 *Z*-0,139 (m) (Seshacharyulu et al., 2002)

) = 3.22 0.16 log (*Z*) (Seshacharyulu et al., 1999)

change in average diameter of primary grains – *d <sup>p</sup>*

depending on morphology of primary and phases:



where: *Z* – Zener-Holomon parameter.


*g <sup>p</sup>*

Fig. 19. The microstructure of Ti-6Al-4V alloy after die forging at 950°C (Kubiak & Sieniawski, 1998)

Fig. 20. Simulation of microstructure development for Ti-6Al-4V alloy deformed at 1050°C at the strain rate = 1.0 s-1 for various strains: a) = 0.7; b) = 5.0; c) = 15, d) = 45; *d <sup>R</sup>* – average diameter of recrystallized grains (Ding et al., 2002)

1 2 3

4 5 6

a) b)

c) d) Fig. 20. Simulation of microstructure development for Ti-6Al-4V alloy deformed at 1050°C at

= 1.0 s-1 for various strains: a) = 0.7; b) = 5.0; c) = 15, d) = 45; *d*

*<sup>R</sup>* –

Fig. 19. The microstructure of Ti-6Al-4V alloy after die forging at 950°C (Kubiak &

Sieniawski, 1998)

the strain rate

average diameter of recrystallized grains (Ding et al., 2002)

Increase in deformation rate decreases the size of recrystallized grains (Fig. 21). The change in average diameter of primary grains – *d <sup>p</sup>* and average thickness of lamellae – *g <sup>p</sup>* , induced by deformation, can be described for Ti-6Al-4V alloy by following equations, depending on morphology of primary and phases:


) = 3.22 0.16 log (*Z*) (Seshacharyulu et al., 1999)

where: *Z* – Zener-Holomon parameter.

Fig. 21. The effect of strain rate on average diameter of primary phase grains in Ti 6Al-4V alloy with globular initial microstructure (Seshacharyulu et al., 1999)

However presented model of dynamic recrystallization (Fig. 20) does not take into account the change of shape of primary grains and deformation of recrystallized grains. Thus it can be stated that it describes process of metadynamic recrystallization. Upon cooling after dynamic or metadynamic recrystallization following phases can be formed within grains (depending on cooling rate): colonies of and lamellae, equiaxed and grains, martensitic '(") phases or the mixture of them.

The character of flow curves obtained during sequential deformation of two phase titanium alloys at the temperature in field confirms occurrence of dynamic processes of microstructure recovery. Increase in time of metadynamic recrystallization reduces strengthening. The character of the curves describing volume fraction of recrystallized phase confirms the influence of chemical composition on recrystallization kinetics (Fig. 22). Increasing content of alloying elements leads to decrease in rate of recrystallization. Time for recrystallization of 50% of phase is equal t0,5 3.5 s for Ti-6Al-5Mo-5V-2Cr-1Fe alloy and t0,5 1.5 s for Ti-6Al-4V alloy.

Hot Plasticity of Alpha Beta Alloys 101

In the 1980s and 1990s deformation maps for CP titanium were calculated on the basis of material constants, describing possible deformation mechanisms operating during processing. In the late 1990s the map of microstructure changes depending on deformation

conditions was developed for Ti-6Al-4V alloy (Fig. 25) (Seshacharyulu et al., 1999).

Fig. 25. The map of microstructure changes and phenomena occurring during hot

Technological hot plasticity of + titanium alloys, characterized by plastic flow stress and critical strain, depends on morphology and stereological parameters of the phases in the

The flow curves = f() of + titanium alloys have similar character. Three stages of flow stress changes can be distinguished, what is characteristic for materials in which dynamic

Fig. 26. The effect of strain rate on plastic flow stress at 900°C for Ti-6Al-4V alloy – coarse-

1 2 3 4 5 6 7 8 9 10 11 12 Strain

f

0.04 0.4 4.0

Strain rate , s-1

deformation of Ti-6Al-4V alloy (Seshacharyulu et al., 1999)

recrystallization occurs (Fig. 26): - increase up to pm value, - decrease down to ps value, - stabilization at ps value.

grained lamellar microstructure (Kubiak, 2004)

p

Plastic flow stressp, MPa

pp

ps

alloy microstructure and deformation conditions (Kubiak, 2004).

Fig. 22. The dependence of volume fraction of recrystallized phase on dwell time during sequential deformation of Ti-6Al-4V and Ti-6Al-5V-5Mo-2Cr-1Fe alloys (Kubiak, 2004)

Microstructure of + titanium alloys after sequential deformation at the temperature in field and air cooling is composed of globular and heavily deformed grains (Fig. 23) in the matrix of the lamellar and phases (Fig. 24a, b). Increase in dwell time during sequential deformation leads to reduction of the strengthening effect and moreover to reduction of dispersion of the phases – growth of both globular and deformed grains. New, recrystallized grains nucleate at the primary grain boundaries, forming chains (Figs 23b, 24b). The dislocation density in globular and lamellar phase and phase is low (Fig. 24c, d).

Fig. 23. The microstructure of Ti-6Al-4V alloy after sequential deformation with various dwell time: a) 1s – transverse section, b) 1s – longitudinal section, c) 100s – transverse section, b) 100s – longitudinal section (Kubiak, 2004)

Fig. 24. The microstructure of Ti-6Al-4V alloy after sequential deformation with various dwell time: a) 1s – transverse section, b) 1s – longitudinal section, c) 100s – transverse section, b) 100s – longitudinal section (Kubiak, 2004)

Fig. 22. The dependence of volume fraction of recrystallized phase on dwell time during sequential deformation of Ti-6Al-4V and Ti-6Al-5V-5Mo-2Cr-1Fe alloys (Kubiak, 2004)

1 10 100 Time, s

Ti-6Al-5Mo-5V-2Cr-1Fe

Ti-6Al-4V

Microstructure of + titanium alloys after sequential deformation at the temperature in field and air cooling is composed of globular and heavily deformed grains (Fig. 23) in the matrix of the lamellar and phases (Fig. 24a, b). Increase in dwell time during sequential deformation leads to reduction of the strengthening effect and moreover to reduction of dispersion of the phases – growth of both globular and deformed grains. New, recrystallized grains nucleate at the primary grain boundaries, forming chains (Figs 23b, 24b). The

a) b) c) d) Fig. 23. The microstructure of Ti-6Al-4V alloy after sequential deformation with various dwell time: a) 1s – transverse section, b) 1s – longitudinal section, c) 100s – transverse

a) b) c) d) Fig. 24. The microstructure of Ti-6Al-4V alloy after sequential deformation with various dwell time: a) 1s – transverse section, b) 1s – longitudinal section, c) 100s – transverse

section, b) 100s – longitudinal section (Kubiak, 2004)

0.2

0.4

0.6

Softening factor for phase

0.8

1

section, b) 100s – longitudinal section (Kubiak, 2004)

dislocation density in globular and lamellar phase and phase is low (Fig. 24c, d).

In the 1980s and 1990s deformation maps for CP titanium were calculated on the basis of material constants, describing possible deformation mechanisms operating during processing. In the late 1990s the map of microstructure changes depending on deformation conditions was developed for Ti-6Al-4V alloy (Fig. 25) (Seshacharyulu et al., 1999).

Fig. 25. The map of microstructure changes and phenomena occurring during hot deformation of Ti-6Al-4V alloy (Seshacharyulu et al., 1999)

Technological hot plasticity of + titanium alloys, characterized by plastic flow stress and critical strain, depends on morphology and stereological parameters of the phases in the alloy microstructure and deformation conditions (Kubiak, 2004).

The flow curves = f() of + titanium alloys have similar character. Three stages of flow stress changes can be distinguished, what is characteristic for materials in which dynamic recrystallization occurs (Fig. 26):


Fig. 26. The effect of strain rate on plastic flow stress at 900°C for Ti-6Al-4V alloy – coarsegrained lamellar microstructure (Kubiak, 2004)

Hot Plasticity of Alpha Beta Alloys 103

The influence of microstructure morphology and deformation condition on maximum flow stress pm for Ti-6Al-5Mo-5V-2Cr-1Fe alloy is similar to that found for Ti-6Al-4V alloy.

The effect of conditions of heat treatment and degree of plastic deformation in thermomechanical process on development of microstructure and plasticity of Ti-6Al-4V and Ti-6Al-2Mo-2Cr titanium alloys in hot tensile test was also investigated (Motyka & Sieniawski, 2010). On the basis of dilatometric results and previous findings conditions of heat treatment and plastic deformation were defined and two schemes of thermomechanical

Fig. 29. Schemes of thermomechanical processing of Ti-6Al-4V alloy with forging reduction

Initial microstructure of Ti-6Al-4V alloy was composed of globular, fine α grains and β phase in the form of thin layers separating α grains (Fig. 30a). Quenching of Ti-6Al-4V alloy from the β phase temperature range led to formation of microstructure composed solely of martensitic α'(α'') phase (Fig. 30b). Microstructure after following plastic deformation in the α+β↔β range with forging reduction of about 20% (TMP-I) and 50% (TMP-II) comprised elongated and deformed grains of primary α phase in the matrix of β transformed phase containing fine globular grains of α secondary phase (Fig. 30c,d). Higher degree of initial deformation led to obtaining finer microstructure containing more elongated α grains – *f*α = 16 for ε = 20% and 21.1 for ε = 50%. The larger volume fraction of α phase was also found

a) b) c) d) Fig. 30. Microstructure (DIC) of Ti-6Al-4V alloy before thermomechanical processing (a), after quenching from the β phase range (b) and after deformation in the α+β↔β range with

forging reduction of 20% (c) and 50% (d) (Motyka & Sieniawski, 2010)

ε ≈ 20% (a) and ε ≈ 50% (b) (WQ - water quenching) (Motyka & Sieniawski, 2010)

(Tab. 2).

processing were worked out, denoted TMP-I and TMP-II respectively (Fig. 29).

Deformation of Ti-6Al-4V alloy at 1050°C – the range of phase stability – results in reduction of maximum plastic flow stress pm regardless of the initial phase morphology and dispersion. Coarse lamellar microstructure shows the maximum plastic flow stress. Reduction of the size of colonies and lamellae of and phases leads to significant decrease in flow stress in comparison with bi-modal and globular microstructure. It was also found that deformation rate have a pronounced effect on the flow stress pm and m strain for Ti-6Al-4V alloy. Increase in strain rate results in higher pm and m values (Figs 27 and 28).

Fig. 27. The dependence of m strain on strain rate, temperature of deformation and microstructure of Ti-6Al-4V alloy: a) bi-modal, b) globular, c) coarse-grained lamellar, d) fine-grained lamellar microstructure (Kubiak, 2004)

Fig. 28. The dependence of maximum plastic flow stress pm on strain rate, temperature of deformation and microstructure of Ti-6Al-4V alloy: a) bi-modal, b) globular, c) coarsegrained lamellar, d) fine-grained lamellar microstructure (Kubiak, 2004)

Deformation of Ti-6Al-4V alloy at 1050°C – the range of phase stability – results in reduction of maximum plastic flow stress pm regardless of the initial phase morphology and dispersion. Coarse lamellar microstructure shows the maximum plastic flow stress. Reduction of the size of colonies and lamellae of and phases leads to significant decrease in flow stress in comparison with bi-modal and globular microstructure. It was also found that deformation rate have a pronounced effect on the flow stress pm and m strain for Ti-6Al-4V alloy. Increase in strain rate results in higher pm and m values (Figs 27 and 28).

Fig. 27. The dependence of m strain on strain rate, temperature of deformation and microstructure of Ti-6Al-4V alloy: a) bi-modal, b) globular, c) coarse-grained lamellar,

a) b)

Strain rate , s-1

0.2 0.4 0.6 0.8 1

0.2 0.4 0.6 0.8 1

d)

0.1 <sup>1</sup> <sup>0</sup>

0.1 <sup>1</sup> <sup>0</sup>

Deformation temperature, °C

> 900 950 1050

Fig. 28. The dependence of maximum plastic flow stress pm on strain rate, temperature of deformation and microstructure of Ti-6Al-4V alloy: a) bi-modal, b) globular, c) coarse-

Strain rate , s-1

d)

0.1 <sup>1</sup> <sup>20</sup>

0.1 <sup>1</sup> <sup>20</sup>

1050 950 900

Deformation temperature, °C

grained lamellar, d) fine-grained lamellar microstructure (Kubiak, 2004)

0.1 1

d) fine-grained lamellar microstructure (Kubiak, 2004)

0.1 <sup>1</sup> <sup>20</sup>

0 0.2 0.4 0.6 0.8

20 40 60 80 100 120 140 Maximum plastic flow stress pm, MPa

c)

Strain

m

0.2 0.4 0.6 0.8 1

1

c)

0.1 1

0.1 <sup>1</sup> <sup>0</sup>

a) b)

The influence of microstructure morphology and deformation condition on maximum flow stress pm for Ti-6Al-5Mo-5V-2Cr-1Fe alloy is similar to that found for Ti-6Al-4V alloy.

The effect of conditions of heat treatment and degree of plastic deformation in thermomechanical process on development of microstructure and plasticity of Ti-6Al-4V and Ti-6Al-2Mo-2Cr titanium alloys in hot tensile test was also investigated (Motyka & Sieniawski, 2010). On the basis of dilatometric results and previous findings conditions of heat treatment and plastic deformation were defined and two schemes of thermomechanical processing were worked out, denoted TMP-I and TMP-II respectively (Fig. 29).

Fig. 29. Schemes of thermomechanical processing of Ti-6Al-4V alloy with forging reduction ε ≈ 20% (a) and ε ≈ 50% (b) (WQ - water quenching) (Motyka & Sieniawski, 2010)

Initial microstructure of Ti-6Al-4V alloy was composed of globular, fine α grains and β phase in the form of thin layers separating α grains (Fig. 30a). Quenching of Ti-6Al-4V alloy from the β phase temperature range led to formation of microstructure composed solely of martensitic α'(α'') phase (Fig. 30b). Microstructure after following plastic deformation in the α+β↔β range with forging reduction of about 20% (TMP-I) and 50% (TMP-II) comprised elongated and deformed grains of primary α phase in the matrix of β transformed phase containing fine globular grains of α secondary phase (Fig. 30c,d). Higher degree of initial deformation led to obtaining finer microstructure containing more elongated α grains – *f*α = 16 for ε = 20% and 21.1 for ε = 50%. The larger volume fraction of α phase was also found (Tab. 2).

Fig. 30. Microstructure (DIC) of Ti-6Al-4V alloy before thermomechanical processing (a), after quenching from the β phase range (b) and after deformation in the α+β↔β range with forging reduction of 20% (c) and 50% (d) (Motyka & Sieniawski, 2010)

Hot Plasticity of Alpha Beta Alloys 105

a) b)

In Ti-6Al-2Mo-2Cr alloy after TMP-I processing dislocations were observed mainly near grain boundaries (Fig. 34a). It was found that the secondary α phase in β transformed matrix occurs in lamellar form (Fig. 34b). Higher degree of deformation in TMP-II process led to higher dislocation density in α phase grains (Fig. 35a) and fragmentation of elongated α

In Ti-6Al-2Mo-2Cr alloy higher volume fraction of β phase (Tab. 3) was found than in Ti-6Al-4V alloy which can be explained by higher value of coefficient of β phase stabilisation *Kβ*.

a) b)

a) b) c) Fig. 34. Microstructure (TEM) of Ti-6Al-2Mo-2Cr alloy after TMP-II process: dislocations in α grains (a), lamellae precipitations of α grains in β transformed phase (b) with indexing of

Fig. 33. Microstructure (TEM) of Ti-6Al-4V alloy after TMP II process: high dislocation

density in α grains (Motyka & Sieniawski, 2010)

diffraction (c) (Motyka & Sieniawski, 2010)

Fig. 32. Microstructure (TEM) of Ti-6Al-4V alloy after TMP-I process: fragmentation of

α phase (a), globular secondary α phase grain (b) (Motyka & Sieniawski, 2010)

grains (Fig. 35b).

Fig. 31. Microstructure (DIC) of Ti-6Al-2Mo-2Cr alloy before thermomechanical processing (a), after quenching from the β phase range (b) and after deformation in the α+β↔β range with forging reduction of 20% (c) and 50% (d) (Motyka & Sieniawski, 2010)

Initial microstructure of Ti-6Al-2Mo-2Cr alloy was composed of colonies of parallel α-lamellae enclosed in primary β phase grains (Fig. 31a). Solution heat treatment led to formation of microstructure composed of martensitic α'(α'') phase, similarly to Ti-6Al-4V alloy (Fig. 31b). Microstructure after thermomechanical processes (TMP-I and TMP-II) comprised fine, elongated grains of α phase in the matrix of β transformed phase (Figs 31c,d). In contrary to Ti-6Al-4V alloy primary β phase grain boundaries were observed. Higher degree of initial deformation in thermomechanical process led to obtaining finer microstructure and larger volume fraction of α phase (Table 2).


Table 2. Stereological parameters of microstructure of as-received and thermomechanically processed Ti-6Al-4V and Ti-6Al-2Mo-2Cr alloys; where: *V* - volume fraction of α phase, *a* and *b* - length of sides of rectangular circumscribed on α grain section, *f* - elongation factor of α phase grains, *a prim* and *b prim* - length of sides of rectangular circumscribed on primary grain section, *f prim* - elongation factor of primary β phase grains, *R* - size of the colony of parallel α lamellae, *g* - thickness of α-lamellae (Motyka & Sieniawski, 2010)

TEM examination of Ti-6Al-4V alloy revealed fragmentation of elongated α grains (Fig. 32a) and presence of globular secondary α grains in the β transformed matrix (Fig. 32b) after TMP-I thermomechanical processing. Higher dislocation density in elongated α grains was observed after TMP-II processing (larger forging reduction) (Figs 33a,b).

a) b) c) d) Fig. 31. Microstructure (DIC) of Ti-6Al-2Mo-2Cr alloy before thermomechanical processing (a), after quenching from the β phase range (b) and after deformation in the α+β↔β range

Initial microstructure of Ti-6Al-2Mo-2Cr alloy was composed of colonies of parallel α-lamellae enclosed in primary β phase grains (Fig. 31a). Solution heat treatment led to formation of microstructure composed of martensitic α'(α'') phase, similarly to Ti-6Al-4V alloy (Fig. 31b). Microstructure after thermomechanical processes (TMP-I and TMP-II) comprised fine, elongated grains of α phase in the matrix of β transformed phase (Figs 31c,d). In contrary to Ti-6Al-4V alloy primary β phase grain boundaries were observed. Higher degree of initial deformation in thermomechanical process led to obtaining finer

> *Va*

> > *a prim b*

As received 82 4.1 5.3 0.77 TMP-I processed 59 51.3 3.2 16 TMP-II processed 79 23.2 1.1 21.1

As received 76 137 42 3.26 12 1 TMP-I processed 34 - - - - 4 TMP-II processed 40 - - - - 1

Table 2. Stereological parameters of microstructure of as-received and thermomechanically

TEM examination of Ti-6Al-4V alloy revealed fragmentation of elongated α grains (Fig. 32a) and presence of globular secondary α grains in the β transformed matrix (Fig. 32b) after TMP-I thermomechanical processing. Higher dislocation density in elongated α grains was


colony of parallel α lamellae, *g* - thickness of α-lamellae (Motyka & Sieniawski, 2010)

Stereological parameters


*f*

*R g* 


*b*

% m

*prim fprim*

*prim* - length of sides of rectangular circumscribed on

% m m

*prim* - elongation factor of primary β phase grains, *R* - size of the

with forging reduction of 20% (c) and 50% (d) (Motyka & Sieniawski, 2010)

microstructure and larger volume fraction of α phase (Table 2).

Condition of Ti-6Al-2Mo-2Cr alloy *<sup>V</sup>*

processed Ti-6Al-4V and Ti-6Al-2Mo-2Cr alloys; where: *V*

*prim* and *b*

observed after TMP-II processing (larger forging reduction) (Figs 33a,b).

Condition of Ti-6Al-4V alloy

and *b* 

factor of α phase grains, *a*

primary grain section, *f*

Fig. 32. Microstructure (TEM) of Ti-6Al-4V alloy after TMP-I process: fragmentation of α phase (a), globular secondary α phase grain (b) (Motyka & Sieniawski, 2010)

In Ti-6Al-2Mo-2Cr alloy after TMP-I processing dislocations were observed mainly near grain boundaries (Fig. 34a). It was found that the secondary α phase in β transformed matrix occurs in lamellar form (Fig. 34b). Higher degree of deformation in TMP-II process led to higher dislocation density in α phase grains (Fig. 35a) and fragmentation of elongated α grains (Fig. 35b).

In Ti-6Al-2Mo-2Cr alloy higher volume fraction of β phase (Tab. 3) was found than in Ti-6Al-4V alloy which can be explained by higher value of coefficient of β phase stabilisation *Kβ*.

Fig. 33. Microstructure (TEM) of Ti-6Al-4V alloy after TMP II process: high dislocation density in α grains (Motyka & Sieniawski, 2010)

Fig. 34. Microstructure (TEM) of Ti-6Al-2Mo-2Cr alloy after TMP-II process: dislocations in α grains (a), lamellae precipitations of α grains in β transformed phase (b) with indexing of diffraction (c) (Motyka & Sieniawski, 2010)

Hot Plasticity of Alpha Beta Alloys 107

Dilatometric examination revealed the influence of forging reduction on critical temperatures of α+β↔β phase transformation. The TMP-II thermomechanical processing with highest strain applied (ε≈50%) caused significant increase in finish temperature of α+β↔β phase transformation in both examined titanium alloys (Tab. 3 and Fig. 36). The temperature range of phase transformation was considerably wider in Ti-6Al-2Mo-2Cr alloy

On the basis of tensile tests at 850°C and 925°C on thermomechanically processed Ti-6Al-4V and Ti-6Al-2Mo-2Cr alloys it was found that the maximum flow stress σpm decreased with growing temperature of deformation but increased with strain rate (Fig. 37). It was found that the maximum flow stress σpm determined in tensile test is higher at lower test temperature 850°C for the strain rate range applied (Fig. 37). There is no significant effect of degree of initial deformation (forging) of two investigated alloys on σpm value for both

The relative elongation *A* of hot deformed Ti-6Al-4V and Ti-6Al-2Mo-2Cr titanium alloys

where the maximum *A* value was achieved for both alloys deformed at 850°C. After thermomechanical processing TMP-II (ε ≈ 50%) alloys exhibit maximum elongations, typical for superplastic deformation (Fig. 38). It seems that higher grain refinement obtained in thermomechanical process enhanced hot plasticity of two-phase titanium alloys deformed with low strain rates. Similar behaviour was observed in previous works on superplasticity of thermomechanically processed Ti-6Al-4V alloy (Motyka, 2007; Motyka & Sieniawski, 2004). It was found that fragmentation and globularization of elongated α phase grains during initial stage of hot deformation restricted grain growth and resulted in higher values

a) b) 1) Results obtained in tensile tests in fine-grained superplasticity region for Ti-6Al-4V alloy after TMP-II

2Mo-2Cr (b) alloys after processing TMP-I and TMP-II (Motyka & Sieniawski, 2010)

dependence (on the basis of tensile test) for Ti-6Al-4V (a) and Ti-6Al-

0

50

100

Stress

pm , MPa

150

200

. 0,01 0,1 <sup>1</sup>

*T* = 850o C, = 20%

*T* = 850o C, = 50%

*T* = 925o C, = 20%

*T* = 925o C, = 50%

Strain rate , s-1

.

above 0.1 the influence of forging reduction ε in thermomechanical processing and

in the whole range applied (Fig. 38). For strain

= 1·10-2 s-1

tensile test temperature is very slight. Considerable differences are visible for

(Tab. 3).

rate 

850°C and 925°C test temperature (Fig. 37).

decreased with the increasing strain rate

of total elongation in tensile test.

*T* = 875o C, = 50% **1**)

*T* = 850o C, = 20%

*T* = 850o C, = 50%

*T* = 925o C, = 20%

*T* = 925o C, = 50%

processing (Motyka & Sieniawski, 2004)

1E-3 0,01 0,1 1

Strain rate , s-1

Fig. 37. The σpm -

0

50

100

Stress

pm , MPa

150

200

Fig. 35. Microstructure (TEM) of Ti-6Al-2Mo-2Cr alloy after TMP-II process: dislocations in α grains (a), precipitations of α grains in β transformed phase (b) (Motyka & Sieniawski, 2010)


Table 3. Critical temperatures of α+β↔β phase transformation of as received and thermomechanically processed two-phase alloys (Motyka & Sieniawski, 2010)

Fig. 36. Dilatometric curves of Ti-6Al-4V (a) and Ti-6Al-2Mo-2Cr (b) alloys in as-received state and after thermomechanical processing (Motyka & Sieniawski, 2010)

a) b) Fig. 35. Microstructure (TEM) of Ti-6Al-2Mo-2Cr alloy after TMP-II process: dislocations in α grains (a), precipitations of α grains in β transformed phase (b) (Motyka & Sieniawski,

> Start of α+β↔β 894 882 912 Finish of α+β↔β 979 976 1009

> Start of α+β↔β 803 800 809 Finish of α+β↔β 991 992 1011

a) b) Fig. 36. Dilatometric curves of Ti-6Al-4V (a) and Ti-6Al-2Mo-2Cr (b) alloys in as-received

state and after thermomechanical processing (Motyka & Sieniawski, 2010)

Table 3. Critical temperatures of α+β↔β phase transformation of as received and thermomechanically processed two-phase alloys (Motyka & Sieniawski, 2010)

Condition of Ti-6Al-4V alloy

processing

Condition of Ti-6Al-2Mo-2Cr alloy

processing

TMP-II processing

TMP-II processing

As-received TMP-I

As-received TMP-I

2010)

Critical temperatures of α+β↔β phase transformation [ºC]

Dilatometric examination revealed the influence of forging reduction on critical temperatures of α+β↔β phase transformation. The TMP-II thermomechanical processing with highest strain applied (ε≈50%) caused significant increase in finish temperature of α+β↔β phase transformation in both examined titanium alloys (Tab. 3 and Fig. 36). The temperature range of phase transformation was considerably wider in Ti-6Al-2Mo-2Cr alloy (Tab. 3).

On the basis of tensile tests at 850°C and 925°C on thermomechanically processed Ti-6Al-4V and Ti-6Al-2Mo-2Cr alloys it was found that the maximum flow stress σpm decreased with growing temperature of deformation but increased with strain rate (Fig. 37). It was found that the maximum flow stress σpm determined in tensile test is higher at lower test temperature 850°C for the strain rate range applied (Fig. 37). There is no significant effect of degree of initial deformation (forging) of two investigated alloys on σpm value for both 850°C and 925°C test temperature (Fig. 37).

The relative elongation *A* of hot deformed Ti-6Al-4V and Ti-6Al-2Mo-2Cr titanium alloys decreased with the increasing strain rate in the whole range applied (Fig. 38). For strain rate above 0.1 the influence of forging reduction ε in thermomechanical processing and tensile test temperature is very slight. Considerable differences are visible for = 1·10-2 s-1 where the maximum *A* value was achieved for both alloys deformed at 850°C. After thermomechanical processing TMP-II (ε ≈ 50%) alloys exhibit maximum elongations, typical for superplastic deformation (Fig. 38). It seems that higher grain refinement obtained in thermomechanical process enhanced hot plasticity of two-phase titanium alloys deformed with low strain rates. Similar behaviour was observed in previous works on superplasticity of thermomechanically processed Ti-6Al-4V alloy (Motyka, 2007; Motyka & Sieniawski, 2004). It was found that fragmentation and globularization of elongated α phase grains during initial stage of hot deformation restricted grain growth and resulted in higher values of total elongation in tensile test.

1) Results obtained in tensile tests in fine-grained superplasticity region for Ti-6Al-4V alloy after TMP-II processing (Motyka & Sieniawski, 2004)

Fig. 37. The σpm - dependence (on the basis of tensile test) for Ti-6Al-4V (a) and Ti-6Al-2Mo-2Cr (b) alloys after processing TMP-I and TMP-II (Motyka & Sieniawski, 2010)

Hot Plasticity of Alpha Beta Alloys 109

Elongation 

Grain size *d* [m]

Temperature *T*  [ºC]

Strain rate [s-1]

8.310-3

[%]

Two-phase + alloys

(*IMI550*) + 2000 4 885 510-4

(*SP-700*) + 2500 2-3 750 10-3

Ti-6Al-4V + 2100 2 850 10-2 Ti-6Al-2Sn-4Zr-2Mo + 2700 1-2 900 10-2 Ti-6Al-2Sn-4Zr-6Mo + 2200 1-2 750 10-2 Ti-6Al-7Nb (*IMI367*) + 300 6 900 310-4 Ti-6.5Al-3.7Mo-1,5Zr + 640 6-7 600 10-4

0,15Si + 2000 4 885 510-4

Ti-10Co-4Al + Ti2Co 1000 0.5 700 510-2 Titanium matrix composites Ti-6Al-4V + 10%TiC + TiC 270 5 870 1.710-4 Ti-6Al-4V + 10%TiN + TiN 410 5 920 1.710-4 Table 4. Superplastic deformation conditions of selected titanium alloys and titanium matrix

Relatively new group of superplastic titanium alloys are TiAl or Ti3Al intermetallics based alloys (Tab. 2). It is well known that intermetallics based alloys have a high relative strength, and good high-temperature creep resistance. Widespread usage of those materials is limited mainly by their low plasticity precluding forming of structural components using conventional plastic working methods. In this case pursuit to obtain fine-grained microstructure enabling superplastic deformation seems to be very promising [Hofmann et

al., 1995; Imayev et al., 1999; Kobayashi et al., 1994; Nieh et al., 1997).

Intermetallics based alloys Ti-24Al-11Nb 2 (Ti3Al) + 1280 4 970 10-3

2 (Ti3Al) 380 2-5 1050 10-3

2 (Ti3Al) 350 0.3 800 8.310-4

2 (Ti3Al) 250 <5 900-1050 210-4-

*-CEZ*) + 1100 2-3 72 210-4

composition

Alloy Phase

Ti-4Al-4Mo-2Sn-0.5Si

Ti-4.5Al-3V-2Mo-2Fe

1Fe (

Ti-5Al-2Sn-4Zr-4Mo-2Cr-

Ti-6Al-2Sn-2Zr-2Mo-2Cr-

Ti-46Al-1Cr-0.2Si (TiAl) +

Ti-48Al-2Nb-2Cr (TiAl) +

Ti-50Al (TiAl) +

composites (Sieniawski & Motyka, 2007)

1) Results obtained in tensile tests in fine-grained superplasticity region for Ti-6Al-4V alloy after TMP-II processing (Motyka & Sieniawski, 2004)

Fig. 38. The A - dependence (on the basis of tensile test) for Ti-6Al-4V (a) and Ti-6Al-2Mo-2Cr (b) alloys after processing TMP-I and TMP-II (Motyka & Sieniawski, 2010)

#### **4. Superplasticity of titanium alloys**

Superplasticity is the ability of polycrystalline materials to exhibit very high value of strain (tensile elongation can be even more than 2000%), appearing in high homologous temperature under exceptionally low stress which is strongly dependent on strain rate. Generally two types of superplasticity are distinguished: fine-structure superplasticity (*FSS*) – considered as an internal structural feature of material and internal-stress superplasticity (*ISS*) caused by special external conditions (e.g. thermal or pressure cycling) generating internal structural transformations that produce high internal stresses independent on external stresses.

*FSS* phenomenon is observed in isotropic fine-grained metallic materials under special conditions: limited range of low strain rates and temperature above 0.4 *T*m. Main features of superplastic deformation are: high value of strain rate sensitivity parameter (*m* > 0.3), lack of strain hardening, equiaxial shape of grains not undergoing changes, conversion of texture during deformation, low activity of lattice dislocations in grains and occurrence of intensive grain boundary sliding (GBS) with associated accommodation mechanisms (Grabski, 1973; Nieh et al., 1997).

One of the titanium alloys which has been extensively studied in aspect of superplasticity is widely used Ti-6Al-4V alloy. Results concerning research on this alloy published in world scientific literature indicate meaningful progress in evaluation and applications of superplasticity in last 30 years. In the beginning of 70's maximum superplastic tensile elongation of Ti-6Al-4V alloy was about 1000% at the strain rate of 10-4 s-1 (Grabski, 1973), whereas in few last years special thermomechanical methods were developed that enabled doubling the tensile elongation and increasing strain rate by the factor of 100 (Inagaki, 1996) (Table 4).

a) b) 1) Results obtained in tensile tests in fine-grained superplasticity region for Ti-6Al-4V alloy

. 0,01 0,1

200

400

Elongation *A*, %

600

800

Superplasticity is the ability of polycrystalline materials to exhibit very high value of strain (tensile elongation can be even more than 2000%), appearing in high homologous temperature under exceptionally low stress which is strongly dependent on strain rate. Generally two types of superplasticity are distinguished: fine-structure superplasticity (*FSS*) – considered as an internal structural feature of material and internal-stress superplasticity (*ISS*) caused by special external conditions (e.g. thermal or pressure cycling) generating internal structural transformations that produce high internal stresses independent on

*FSS* phenomenon is observed in isotropic fine-grained metallic materials under special conditions: limited range of low strain rates and temperature above 0.4 *T*m. Main features of superplastic deformation are: high value of strain rate sensitivity parameter (*m* > 0.3), lack of strain hardening, equiaxial shape of grains not undergoing changes, conversion of texture during deformation, low activity of lattice dislocations in grains and occurrence of intensive grain boundary sliding (GBS) with associated accommodation mechanisms (Grabski, 1973;

One of the titanium alloys which has been extensively studied in aspect of superplasticity is widely used Ti-6Al-4V alloy. Results concerning research on this alloy published in world scientific literature indicate meaningful progress in evaluation and applications of superplasticity in last 30 years. In the beginning of 70's maximum superplastic tensile elongation of Ti-6Al-4V alloy was about 1000% at the strain rate of 10-4 s-1 (Grabski, 1973), whereas in few last years special thermomechanical methods were developed that enabled doubling the tensile elongation and increasing strain rate by the factor of 100 (Inagaki, 1996)

2Cr (b) alloys after processing TMP-I and TMP-II (Motyka & Sieniawski, 2010)

dependence (on the basis of tensile test) for Ti-6Al-4V (a) and Ti-6Al-2Mo-

Strain rate , s-1

.

*T* = 850o C, = 20%

*T* = 850o C, = 50%

*T* = 925o C, = 20%

*T* = 925o C, = 50%

after TMP-II processing (Motyka & Sieniawski, 2004)

Strain rate , s-1

*T* = 875o C, = 50% **1**)

*T* = 850o C, = 20%

*T* = 850o C, = 50%

*T* = 925o C, = 20%

*T* = 925o C, = 50%

Fig. 38. The A -

Elongation *A*, %

external stresses.

Nieh et al., 1997).

(Table 4).

**4. Superplasticity of titanium alloys** 

1E-3 0,01 0,1


Table 4. Superplastic deformation conditions of selected titanium alloys and titanium matrix composites (Sieniawski & Motyka, 2007)

Relatively new group of superplastic titanium alloys are TiAl or Ti3Al intermetallics based alloys (Tab. 2). It is well known that intermetallics based alloys have a high relative strength, and good high-temperature creep resistance. Widespread usage of those materials is limited mainly by their low plasticity precluding forming of structural components using conventional plastic working methods. In this case pursuit to obtain fine-grained microstructure enabling superplastic deformation seems to be very promising [Hofmann et al., 1995; Imayev et al., 1999; Kobayashi et al., 1994; Nieh et al., 1997).

Hot Plasticity of Alpha Beta Alloys 111

a) b) Fig. 40. Microstructure of Ti-6Al-4V alloy before (a) and after superplastic deformation (b) -

Along with grain size and shape, volume fraction of particular phases in the alloy also affects its superplasticity. Properties of phases in two-phase α+β titanium alloys differ considerably. α phase (hcp) has less slip systems and two order of magnitude lower selfdiffusion coefficient than β phase (bcc). These features suggest that in the superplasticity conditions α phase has a higher plasticity than β phase. It was confirmed by results obtained from experiments on superplasticity in Ti-6Al-4V alloy where deformation in α grains was observed. Density of dislocations was found to be very low in β grains (Bylica & Sieniawski,

It was established that increase in volume fraction of β phase in alloy causes decrease of the effect of α grain size (Meier et al., 1991). Maximum values of elongation and strain rate sensitivity factor *m* as a function of β volume fraction is shown in Figure 41. Increase in relative volume of β phase causes improvement of superplasticity of titanium alloys. The best superplastic properties of two-phase α+β titanium alloys are achieved for 40-50% volume fraction of β phase (Nieh et al., 1997). Whereas similar properties of intermetallics based alloys are possible for about (20-30)% volume fraction of β phase (Kim et al., 1999; Lee

Superplasticity of titanium alloys depends on relationship between grain growth control and plasticity. β grains are characterized by high diffusivity therefore they grow extremely rapidly at the superplastic deformation temperature which does not favour superplastic flow (Meier et al., 1992). Particular volume fraction of α phase considerably limits β grains growth because in this case long distance diffusion of alloying elements is necessary (e.g. vanadium in β phase). The second phase, besides stabilization of microstructure, influences the rate of grain boundary (α/α, β/β) and phase boundary (α/β) sliding (Inagaki, 1996; Jain et al., 1991, Meier et al., 1991; Nieh et al., 1997). Increase of volume fraction of β phase causes decrease of α/α grain boundary areas and consequently their contribution to deformation by GBS. It is thought that improvement of superplasticity of α+β titanium alloys caused by increase of volume of β phase should be considered in following aspects (Inagaki, 1996): α/β phase boundary sliding, β/β GBS

temperature 850ºC and strain rate of 10-3 s-1 (Sieniawski & Motyka, 2007)

1985; Inagaki, 1996; Jain et al., 1991, Meier et al., 1991; Nieh et al., 1997).

and contribution of other deformation mechanisms.

et al., 1995).

Main criterion for superplastic materials is possibility of obtaining fine-grained and equiaxial microstructure. Desired microstructure is most often obtained by conventional plastic working methods coupled with suitable heat treatment and severe plastic deformation methods (i.e. equal-channel angular pressing – ECAP). Superplastic forming (SPF) of titanium alloys is limited by relatively long time and high deformation temperature. It was established that grain refinement causes increase of strain rate and decrease of superplastic deformation temperature (Fig. 39) (Sieniawski & Motyka, 2007).

Fig. 39. Effect of grain size on strain rate of superplastic deformation for Ti-6Al-4V and Ti-5Al-2.5Sn alloys (a) and on superplastic deformation temperature for Ti-6.5Al-3.3Mo-1.8Zr-0.26Si alloy (b) (Sieniawski & Motyka, 2007)

Taking into account the mechanism of superplastic deformation equiaxed microstructure favours proceeding of GBS. It was found that in fine grained polycrystalline materials with grains elongated crosswise deformation direction GBS is limited. The main reason is difficulty of deformation accommodation in triple points. Transverse deformation is also related to cavities formation along grain boundaries and precludes superplastic deformation (Nieh et al., 1997). It is emphasised that superplastic deformation does not cause shape changes of equiaxed grains. However, gradual transformation of texture is observed what indicates that GBS plays a crucial role in superplastic deformation (Grabski, 1973; Zelin, 1996).

On the basis of the results of research works conducted at the Department of Materials Science of Rzeszow University of Technology it was found that initial microstructure of superplastic titanium alloy can be different from equiaxed one. High superplasticity was observed in Ti-6Al-4V alloy with microstructure composed of strongly elongated and deformed α grains (Fig. 40a) (Motyka & Sieniawski, 2004). It was established that during heating and first stage of superplastic deformation significant changes of the morphology of phases occur (Fig. 40b) (Motyka, 2007).

Main criterion for superplastic materials is possibility of obtaining fine-grained and equiaxial microstructure. Desired microstructure is most often obtained by conventional plastic working methods coupled with suitable heat treatment and severe plastic deformation methods (i.e. equal-channel angular pressing – ECAP). Superplastic forming (SPF) of titanium alloys is limited by relatively long time and high deformation temperature. It was established that grain refinement causes increase of strain rate and decrease of

a) b)

Taking into account the mechanism of superplastic deformation equiaxed microstructure favours proceeding of GBS. It was found that in fine grained polycrystalline materials with grains elongated crosswise deformation direction GBS is limited. The main reason is difficulty of deformation accommodation in triple points. Transverse deformation is also related to cavities formation along grain boundaries and precludes superplastic deformation (Nieh et al., 1997). It is emphasised that superplastic deformation does not cause shape changes of equiaxed grains. However, gradual transformation of texture is observed what indicates that GBS plays a crucial role in superplastic deformation (Grabski, 1973; Zelin,

On the basis of the results of research works conducted at the Department of Materials Science of Rzeszow University of Technology it was found that initial microstructure of superplastic titanium alloy can be different from equiaxed one. High superplasticity was observed in Ti-6Al-4V alloy with microstructure composed of strongly elongated and deformed α grains (Fig. 40a) (Motyka & Sieniawski, 2004). It was established that during heating and first stage of superplastic deformation significant changes of the morphology of

Fig. 39. Effect of grain size on strain rate of superplastic deformation for Ti-6Al-4V and Ti-5Al-2.5Sn alloys (a) and on superplastic deformation temperature for Ti-6.5Al-3.3Mo-

1.8Zr-0.26Si alloy (b) (Sieniawski & Motyka, 2007)

phases occur (Fig. 40b) (Motyka, 2007).

1996).

superplastic deformation temperature (Fig. 39) (Sieniawski & Motyka, 2007).

Fig. 40. Microstructure of Ti-6Al-4V alloy before (a) and after superplastic deformation (b) temperature 850ºC and strain rate of 10-3 s-1 (Sieniawski & Motyka, 2007)

Along with grain size and shape, volume fraction of particular phases in the alloy also affects its superplasticity. Properties of phases in two-phase α+β titanium alloys differ considerably. α phase (hcp) has less slip systems and two order of magnitude lower selfdiffusion coefficient than β phase (bcc). These features suggest that in the superplasticity conditions α phase has a higher plasticity than β phase. It was confirmed by results obtained from experiments on superplasticity in Ti-6Al-4V alloy where deformation in α grains was observed. Density of dislocations was found to be very low in β grains (Bylica & Sieniawski, 1985; Inagaki, 1996; Jain et al., 1991, Meier et al., 1991; Nieh et al., 1997).

It was established that increase in volume fraction of β phase in alloy causes decrease of the effect of α grain size (Meier et al., 1991). Maximum values of elongation and strain rate sensitivity factor *m* as a function of β volume fraction is shown in Figure 41. Increase in relative volume of β phase causes improvement of superplasticity of titanium alloys. The best superplastic properties of two-phase α+β titanium alloys are achieved for 40-50% volume fraction of β phase (Nieh et al., 1997). Whereas similar properties of intermetallics based alloys are possible for about (20-30)% volume fraction of β phase (Kim et al., 1999; Lee et al., 1995).

Superplasticity of titanium alloys depends on relationship between grain growth control and plasticity. β grains are characterized by high diffusivity therefore they grow extremely rapidly at the superplastic deformation temperature which does not favour superplastic flow (Meier et al., 1992). Particular volume fraction of α phase considerably limits β grains growth because in this case long distance diffusion of alloying elements is necessary (e.g. vanadium in β phase). The second phase, besides stabilization of microstructure, influences the rate of grain boundary (α/α, β/β) and phase boundary (α/β) sliding (Inagaki, 1996; Jain et al., 1991, Meier et al., 1991; Nieh et al., 1997). Increase of volume fraction of β phase causes decrease of α/α grain boundary areas and consequently their contribution to deformation by GBS. It is thought that improvement of superplasticity of α+β titanium alloys caused by increase of volume of β phase should be considered in following aspects (Inagaki, 1996): α/β phase boundary sliding, β/β GBS and contribution of other deformation mechanisms.

Hot Plasticity of Alpha Beta Alloys 113

technological plasticity criterion. The value of critical strain depends on microstructure morphology and deformation conditions. The character of dependence of plastic flow stress as well as results of microstructure examination support conclusion that dynamic processes of microstructure recovery take place above the temperature range of α+β↔β

Thermomechanical processing enables microstructure and hot plasticity development of two-phase α+β titanium alloys. Increase in degree of initial deformation (forging) in proposed thermomechanical processing leads to formation of more elongated and refined α grains in tested α+β alloys. The most significant effect of degree of initial deformation occurs for the lowest strain rate and lower tensile test temperature used, resulting in

High superplasticity of the Ti-6Al-4V alloy does not necessarily require equiaxial microstructure. Changes of the morphology of phases during heating and first stage of superplastic deformation enables superplastic behaviour of the alloy with initial

Balasubramanian, S. & Anand, L. (2002). Plasticity of Initially Textured Hexagonal

Brooks, J.W. (1996). Processing Wrought Nickel and Titanium Superalloys, *Proceedings of Int.* 

Bylica, A. & Sieniawski, J. (1985). *Titanium and Its Alloys* (in Polish), PWN, ISBN 83-0105-

Churchman, A.T. (1955). The Formation and Removal of Twins in Ti at Low and High

Ding, R., Guo, Z.X. & Wilson A. (2002). Microstructural Evolution of a Ti-6Al-4V Alloy

Ezugwu, E.O. & Wang, Z.M. (1997). Titanium Alloys and their Machinability -

Ghosh, A.K. & Hamilton, C. H. (1979). Mechanical Behaviour and Hardening Characteristics

Glazunov, S. & Kolachev, B.A. (1980). *Physical Metallurgy of Titanium Alloys* (in Russian),

Glazunov, S. & Mojsiejew, W.N. (1974). *Structural Titanium Alloys* (in Russian),

Hadasik, E. (1979). *Analysis of plastic working processes of titanium alloys based on plastometer examinations* (in Polish), PhD Thesis (not published), Katowice, Poland

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microstructure composed of strongly elongated and deformed α grains.

*Materialia,* Vol. 50, pp. 133-148, ISSN 1359-6454

916-305-4213, Stockholm, Sweden, September 1996

phase transformation.

**6. References** 

considerable rise of elongation *A*.

888-9, Warsaw, Poland

ISSN 0924-0136

706, ISSN 0026-086X

A327, pp. 233-245, ISSN 0921-5093

Mietallurgiya, Moscow, Russia

Mietallurgiya, Moscow, Russia

Fig. 41. Effect of volume fraction of β phase on elongation (a) and strain rate sensitivity *m* (b) in selected titanium alloys (Nieh et al., 1997)

Most often microstructure of α+β superplastic titanium alloys is composed of α and β grains which have similar size and shape. Interesting results was obtained for Ti-6Al-4V alloy where α grains were separated by thin films of β phase. Superplastic elongation in this case was more than 2000%. Further investigations indicated that during superplastic deformation thin films of β phase coagulated in triple points into larger particles having irregular forms. Thanks to that α/α grain boundaries free of β thin films were formed. It can be expected that sliding along these grain boundaries proceeds easily. However it was revealed that at this stage superplastic deformation is almost completed and deformation within grains becomes dominant deformation mechanism. It seems that α/α grain boundary sliding is not dominant superplastic deformation mechanism. In this case the effect of β phase thin film can be comparable to role of grain boundaries in single phase materials. Slip and shearing in β phase thin film is caused by movement and rotation of neighbouring α grains. Mentioned processes enable accommodation of grain boundary and phase boundary sliding (Inagaki, 1996). Other investigations also indicate accommodative role of β phase, in which substantially higher dislocations density is observed than in α phase grains. It was noticed simultaneously that dislocations density in α phase increases together with decrease in temperature and increase in strain rate of superplastic deformation (Kim et al., 1999; Meier et al., 1991). Superplasticity of titanium alloys with intermetallic phases like Ti-12Co-5Al and Ti-6Co-6Ni-5Al is observed for grain size about 0.5 m. Particles of Ti2Co and Ti2Ni phases (about 27% of volume) advantageously influence the grain refinement and limit grain growth during superplastic deformation (Nieh et al., 1997).

#### **5. Conclusion**

Hot plasticity of two-phase titanium alloys strongly depends on values of stereological parameters of microstructure. It is possible to develop appropriate microstructure of these alloys yielding optimum plastic flow stress and critical strain values based on

a) b) Fig. 41. Effect of volume fraction of β phase on elongation (a) and strain rate sensitivity *m*

Most often microstructure of α+β superplastic titanium alloys is composed of α and β grains which have similar size and shape. Interesting results was obtained for Ti-6Al-4V alloy where α grains were separated by thin films of β phase. Superplastic elongation in this case was more than 2000%. Further investigations indicated that during superplastic deformation thin films of β phase coagulated in triple points into larger particles having irregular forms. Thanks to that α/α grain boundaries free of β thin films were formed. It can be expected that sliding along these grain boundaries proceeds easily. However it was revealed that at this stage superplastic deformation is almost completed and deformation within grains becomes dominant deformation mechanism. It seems that α/α grain boundary sliding is not dominant superplastic deformation mechanism. In this case the effect of β phase thin film can be comparable to role of grain boundaries in single phase materials. Slip and shearing in β phase thin film is caused by movement and rotation of neighbouring α grains. Mentioned processes enable accommodation of grain boundary and phase boundary sliding (Inagaki, 1996). Other investigations also indicate accommodative role of β phase, in which substantially higher dislocations density is observed than in α phase grains. It was noticed simultaneously that dislocations density in α phase increases together with decrease in temperature and increase in strain rate of superplastic deformation (Kim et al., 1999; Meier et al., 1991). Superplasticity of titanium alloys with intermetallic phases like Ti-12Co-5Al and Ti-6Co-6Ni-5Al is observed for grain size about 0.5 m. Particles of Ti2Co and Ti2Ni phases (about 27% of volume) advantageously influence the grain refinement and limit grain

Hot plasticity of two-phase titanium alloys strongly depends on values of stereological parameters of microstructure. It is possible to develop appropriate microstructure of these alloys yielding optimum plastic flow stress and critical strain values based on

(b) in selected titanium alloys (Nieh et al., 1997)

growth during superplastic deformation (Nieh et al., 1997).

**5. Conclusion** 

technological plasticity criterion. The value of critical strain depends on microstructure morphology and deformation conditions. The character of dependence of plastic flow stress as well as results of microstructure examination support conclusion that dynamic processes of microstructure recovery take place above the temperature range of α+β↔β phase transformation.

Thermomechanical processing enables microstructure and hot plasticity development of two-phase α+β titanium alloys. Increase in degree of initial deformation (forging) in proposed thermomechanical processing leads to formation of more elongated and refined α grains in tested α+β alloys. The most significant effect of degree of initial deformation occurs for the lowest strain rate and lower tensile test temperature used, resulting in considerable rise of elongation *A*.

High superplasticity of the Ti-6Al-4V alloy does not necessarily require equiaxial microstructure. Changes of the morphology of phases during heating and first stage of superplastic deformation enables superplastic behaviour of the alloy with initial microstructure composed of strongly elongated and deformed α grains.

#### **6. References**


Hot Plasticity of Alpha Beta Alloys 115

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Raub, E. & Röschel, E. (1963). Ruthenium with Titanium and Zirconium Alloys (in German),

Sakai, T. (1995). Dynamic Recrystallization Microstructure Under Hot Working Conditions.

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Seshacharyulu, T., Medeiros, S.C., Frazier W.G. & Prasad Y.V.R.K. (2002). Microstructural

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**6** 

*1,2Malaysia 3Japan* 

*1Universiti Teknologi Malaysia,* 

*2Universiti Tun Hussein Onn Malaysia,* 

*3Tokyo University of Agriculture and Technology,* 

**Machinability of Titanium Alloys in Drilling** 

Hole making is an essential process in the structural frames of an aircraft and contributes to 40 to 60% of the total material removal operations (Brinksmeier, 1990). This process is commonly divided into short hole or deep hole drilling. Short hole drilling typically covers holes with a small depth to diameter ratio having diameter up to 30 mm and a depth of not more than 5 times the diameter. Meanwhile deep hole drilling caters for holes greater than 30 mm in diameter and the depths are usually greater than 2.5 times the hole diameter. Drilling deeper hole with conventional drills requires pecking method to enable easy flow of the chips out of the hole. Deep hole drilling is more difficult especially when hole straightness is the main concern. Therefore, a usual method is to make a circular cut using a hollow-core cutting tool. This technique allows larger hole diameter to be drilled with lesser power. In addition, holes can be produced in many forms which include through holes or blind holes (Fig. 1). Through hole is one which is drilled completely through the workpiece

A twist drill is fabricated with 3 major parts as shown in Fig. 2. The most important features from the analytical point of view are rake angle, point angle, web thickness, nominal clearance angle, drill diameter, inclination angle and chisel edge angle. The rake angle is usually specified as helix angle at the periphery. The direction of the chip flow is attributed to the point angle. The torque decreases with increasing point angle due to the increase of orthogonal rake angle at each point on the main cutting edges. Furthermore, the thrust force

Fig. 3 shows the phases involved in a drilling operation, first is the start and centering phase, second is the full drilling phase and finally the break through phase (Tonshoff et al., 1994). To ensure good surface quality and accuracy of the holes are achieved, the first phase is very important (Fig. 3 (a)) in order to avoid the occurrence of premature wear and breakage of the drill. In this phase, the torque and force on the tool constantly increase. The full drilling phase starts once the main cutting edges are fully engaged (Fig. 3 (b)). The break through phase begins when the drill point breaks through the underside of the work piece and the process is stopped when the drill body passed through the work piece (Fig. 3 (c)).

**1. Introduction** 

**1.1 Drilling technology** 

while a blind hole is drilled only to a certain depth.

always increases with increasing point angle.

Safian Sharif1, Erween Abd Rahim2 and Hiroyuki Sasahara3


### **Machinability of Titanium Alloys in Drilling**

Safian Sharif1, Erween Abd Rahim2 and Hiroyuki Sasahara3

*1Universiti Teknologi Malaysia, 2Universiti Tun Hussein Onn Malaysia, 3Tokyo University of Agriculture and Technology, 1,2Malaysia 3Japan* 

#### **1. Introduction**

116 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Wassermann, G. & Grewen, J. (1962). *Texturen Metallischer Werkstoffe,* Springer Verlag, ISBN

Zelin, M.G. (1996). Processes of Microstructural Evolution During Superplastic Deformation,

Zwicker, U. (1974). *Titanium and Titanium Alloys* (in German), Springer Verlag, ISBN 978-

*Materials Characterization,* Vol. 37, pp. 311-329, ISSN 1044-5803

354-002-9214, Berlin, Germany

354-0052-33-3, Berlin-Heidelberg-New York

#### **1.1 Drilling technology**

Hole making is an essential process in the structural frames of an aircraft and contributes to 40 to 60% of the total material removal operations (Brinksmeier, 1990). This process is commonly divided into short hole or deep hole drilling. Short hole drilling typically covers holes with a small depth to diameter ratio having diameter up to 30 mm and a depth of not more than 5 times the diameter. Meanwhile deep hole drilling caters for holes greater than 30 mm in diameter and the depths are usually greater than 2.5 times the hole diameter. Drilling deeper hole with conventional drills requires pecking method to enable easy flow of the chips out of the hole. Deep hole drilling is more difficult especially when hole straightness is the main concern. Therefore, a usual method is to make a circular cut using a hollow-core cutting tool. This technique allows larger hole diameter to be drilled with lesser power. In addition, holes can be produced in many forms which include through holes or blind holes (Fig. 1). Through hole is one which is drilled completely through the workpiece while a blind hole is drilled only to a certain depth.

A twist drill is fabricated with 3 major parts as shown in Fig. 2. The most important features from the analytical point of view are rake angle, point angle, web thickness, nominal clearance angle, drill diameter, inclination angle and chisel edge angle. The rake angle is usually specified as helix angle at the periphery. The direction of the chip flow is attributed to the point angle. The torque decreases with increasing point angle due to the increase of orthogonal rake angle at each point on the main cutting edges. Furthermore, the thrust force always increases with increasing point angle.

Fig. 3 shows the phases involved in a drilling operation, first is the start and centering phase, second is the full drilling phase and finally the break through phase (Tonshoff et al., 1994). To ensure good surface quality and accuracy of the holes are achieved, the first phase is very important (Fig. 3 (a)) in order to avoid the occurrence of premature wear and breakage of the drill. In this phase, the torque and force on the tool constantly increase. The full drilling phase starts once the main cutting edges are fully engaged (Fig. 3 (b)). The break through phase begins when the drill point breaks through the underside of the work piece and the process is stopped when the drill body passed through the work piece (Fig. 3 (c)).

Machinability of Titanium Alloys in Drilling 119

Heat generation, pressure, friction and stress distribution are the main contributors of drill wear. The drill wear can be classified into (Kanai et al., 1978): outer corner (*w*), flank wear (*Vb*), margin wear (*Mw*), crater wear (*KM*), along with two types of chisel edge wear (*CT* and *CM*) and chipping at the cutting lips (*PT* and *PM*). Fig. 4 shows the aforementioned types of wear. Wear starts at the sharp corners of the cutting edges and distributed along the cutting edges until the chisel and drill margin (Schnieder, 2001). Flank wear is considered as one of the criterion to measure the performance of a drill. It occurs due to the friction between the workpiece and the contact area on the clearance surface. However, Kanai et al. (1978) suggested that outer corner wear should be used as the main criteria of tool performance because of the relative ease of measurement and the close relationship between this type of

a) Outer corner wear b) Flank Wear

c) Margin Wear d) Crater Wear

e) Chisel Edge Wear f) Chipping

Crater wear was also observed on the rake face of the drill and can be found clearly around the outer corners of the cutting edges (Choudhury & Raju, 2000; Kaldor & Lenz, 1980). According to Dolinsek et al., (2001), wear land behind the cutting edges is less significant as

Fig. 4. Types of drill wear (Kanai et al., 1978)

**2. Tool wear in drilling process** 

wear and the drill life.

Fig. 1. Type of hole, (a) Through hole and (b) blind hole

Fig. 2. Drill geometry (Lindberg, 1990)

Fig. 3. Drilling phases, (a) centering phase, (b) full drilling phase and (c) break through phase

(a) (b)

Fig. 3. Drilling phases, (a) centering phase, (b) full drilling phase and (c) break through phase

(a) (b) (c)

Fig. 1. Type of hole, (a) Through hole and (b) blind hole

Fig. 2. Drill geometry (Lindberg, 1990)

### **2. Tool wear in drilling process**

Heat generation, pressure, friction and stress distribution are the main contributors of drill wear. The drill wear can be classified into (Kanai et al., 1978): outer corner (*w*), flank wear (*Vb*), margin wear (*Mw*), crater wear (*KM*), along with two types of chisel edge wear (*CT* and *CM*) and chipping at the cutting lips (*PT* and *PM*). Fig. 4 shows the aforementioned types of wear. Wear starts at the sharp corners of the cutting edges and distributed along the cutting edges until the chisel and drill margin (Schnieder, 2001). Flank wear is considered as one of the criterion to measure the performance of a drill. It occurs due to the friction between the workpiece and the contact area on the clearance surface. However, Kanai et al. (1978) suggested that outer corner wear should be used as the main criteria of tool performance because of the relative ease of measurement and the close relationship between this type of wear and the drill life.

e) Chisel Edge Wear f) Chipping

Fig. 4. Types of drill wear (Kanai et al., 1978)

Crater wear was also observed on the rake face of the drill and can be found clearly around the outer corners of the cutting edges (Choudhury & Raju, 2000; Kaldor & Lenz, 1980). According to Dolinsek et al., (2001), wear land behind the cutting edges is less significant as

Machinability of Titanium Alloys in Drilling 121

Lightweight materials such as titanium alloys are now being constituted in modern aircraft structure especially in jet engine components that are subjected to temperatures up to 1000° C. Titanium alloys possess the best combination of physical and metallurgical properties and have established to be quite attractive as engineering materials due to their high strength-to-weight ratio, low density, excellent corrosion resistance, excellent erosion

Titanium alloys are classified into groups based on the alloying elements and the resultant predominant room temperature constituent phases. These groups include α alloy, α- β alloy and β alloy. The α alloys can be divided into two types, commercially pure grades of titanium and those with additions of α- stabilizers such as Al and Sn. α alloys are non-heat treatable and are generally very weldable. They have low to medium strength, good notch toughness, reasonably good ductility and possess excellent mechanical properties which offer optimum high temperature creep strength and oxidation resistance (Boyer, 1996; Ezugwu and Wang, 1997). These include alloys such as Ti-3Al-2.5V, Ti-5Al-2.5Sn, Ti-8Al-1Mo-1V and Ti-6Al-2Sn-4Zr-2Mo. A wide variety of application for α alloys includes gas turbine engine casings, air

Most of the titanium alloys used in the industry contain α- and β- stabilizers. These alloys include Ti-6Al-4V, Ti-6Al-6V-2Sn and Ti-6Al-2Sn-4Zr-6Mo. They are heat treatable and most are weldable especially with the lower β- stabilizer. Their strength levels are medium to high. These alloys possess excellent combination of strength, toughness and corrosion resistance. Typical applications include blades and discs for jet engine turbines and compressors, structural aircraft components and landing gear, chemical process equipment, marine components and surgical implants. Meanwhile, β alloys contain small amounts of αstabilizing elements as strengtheners and generally weldable, high corrosion resistance and good creep resistance to intermediate temperatures. Additions of vanadium, iron and chromium as stabilizing elements, provide superior hot working characteristics. Ti-10V-2Fe-3Al, Ti-15V-3Cr-3Al-3Sn, Ti-15Mo-2.7Nb-3Al-0.2Si and Ti-3Al-8V-6Cr-4Mo-4Zr are examples of these alloys. Typical applications include airframe components, fasteners,

Research works on the machinability of titanium alloys have been conducted extensively and reviewed comprehensively by several researchers. The increasing demands of titanium alloys with excellent high temperature, mechanical and chemical properties make them more difficult to machine. According to Ezugwu et al. (2003), machinability can be phrased as the difficulty to machine a particular material under a given set of the machining parameters such as cutting speed, feed rate and depth of cut. It can be rated in terms of tool life, surface quality, the reaction of cutting forces and also machining cost per part. Basically, work hardening, low thermal conductivity, abrasiveness, high strength level and high heat generated were the dominant reasons for the difficulty in machining titanium alloys. Heat is the most important factor that needs to be aware of when machining titanium alloys. Excessive heat could damage the cutting tool rapidly. The main sources of heat during machining are from the shear zone, from the tool-chip interface friction and from the tool-

resistance and low modulus of elasticity (Brewer et al., 1998)

frame skin and structural components and jet engine compressor blades.

springs, pipe and commercial and consumer products.

**4. Machinability of titanium alloys** 

**3. Titanium alloys** 

an indicator of tool wear because it depends on the relief angle. They suggested that the drill will be considered damaged once the corner of the drill has been rounded off as shown in Fig. 5. However, Fujise and Ohtani (1998) and Harris et al. (2003) considered the outer corner wear as their tool rejection criteria (Fig. 6). The tools were rejected when the outer corner wear reached 75% of the total margin width. Kaldor and Lenz (1980) also employed the corner wear as the tool life criterion in drilling because of the similar wear behavior of other cutting tools.

Fig. 5. Location of flank wear land on the drill (Dolinsek et al., 2001)

Fig. 6. A method to measure outer corner wear from a fixed reference point (Harris et al., 2003)

Tetsutaro & Zhao (1989) considered that the tool is rejected when the maximum flank wear width, *Vb,max* reached 0.7 mm when drilling plain steel. Wen & Xiao (2000) used to measure the wear width developed on the flank surfaces when they drilled stainless steel. Lin (2002) rejected the tool based on the tool rejection criteria when maximum flank wear land exceeded 0.8 mm, surface roughness value exceeded 5.0 μm, excessive outer corner tearing and chipping of the helix flutes. Choudhury & Raju (2000) have studied the influence of feed and speed on crater wear at different points along the cutting lip in drilling. Ezugwu & Lai (1995) rejected the drill bit when maximum flank wear in excess of 0.38 mm on any of the drill lips, a squeaking noise occurring during machining and fracture or catastrophic failure of the drill. These criteria were used when they investigated the drilling of Inconel 901 using HSS drills.

#### **3. Titanium alloys**

120 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

an indicator of tool wear because it depends on the relief angle. They suggested that the drill will be considered damaged once the corner of the drill has been rounded off as shown in Fig. 5. However, Fujise and Ohtani (1998) and Harris et al. (2003) considered the outer corner wear as their tool rejection criteria (Fig. 6). The tools were rejected when the outer corner wear reached 75% of the total margin width. Kaldor and Lenz (1980) also employed the corner wear as the tool life criterion in drilling because of the similar wear behavior of

Fig. 5. Location of flank wear land on the drill (Dolinsek et al., 2001)

Margin width

Fig. 6. A method to measure outer corner wear from a fixed reference point (Harris et al.,

Fixed reference point

Tetsutaro & Zhao (1989) considered that the tool is rejected when the maximum flank wear width, *Vb,max* reached 0.7 mm when drilling plain steel. Wen & Xiao (2000) used to measure the wear width developed on the flank surfaces when they drilled stainless steel. Lin (2002) rejected the tool based on the tool rejection criteria when maximum flank wear land exceeded 0.8 mm, surface roughness value exceeded 5.0 μm, excessive outer corner tearing and chipping of the helix flutes. Choudhury & Raju (2000) have studied the influence of feed and speed on crater wear at different points along the cutting lip in drilling. Ezugwu & Lai (1995) rejected the drill bit when maximum flank wear in excess of 0.38 mm on any of the drill lips, a squeaking noise occurring during machining and fracture or catastrophic failure of the drill. These criteria were used when they investigated the drilling of Inconel

Unknown Margin

Outer corner flank wear land

other cutting tools.

2003)

901 using HSS drills.

Lightweight materials such as titanium alloys are now being constituted in modern aircraft structure especially in jet engine components that are subjected to temperatures up to 1000° C. Titanium alloys possess the best combination of physical and metallurgical properties and have established to be quite attractive as engineering materials due to their high strength-to-weight ratio, low density, excellent corrosion resistance, excellent erosion resistance and low modulus of elasticity (Brewer et al., 1998)

Titanium alloys are classified into groups based on the alloying elements and the resultant predominant room temperature constituent phases. These groups include α alloy, α- β alloy and β alloy. The α alloys can be divided into two types, commercially pure grades of titanium and those with additions of α- stabilizers such as Al and Sn. α alloys are non-heat treatable and are generally very weldable. They have low to medium strength, good notch toughness, reasonably good ductility and possess excellent mechanical properties which offer optimum high temperature creep strength and oxidation resistance (Boyer, 1996; Ezugwu and Wang, 1997). These include alloys such as Ti-3Al-2.5V, Ti-5Al-2.5Sn, Ti-8Al-1Mo-1V and Ti-6Al-2Sn-4Zr-2Mo. A wide variety of application for α alloys includes gas turbine engine casings, air frame skin and structural components and jet engine compressor blades.

Most of the titanium alloys used in the industry contain α- and β- stabilizers. These alloys include Ti-6Al-4V, Ti-6Al-6V-2Sn and Ti-6Al-2Sn-4Zr-6Mo. They are heat treatable and most are weldable especially with the lower β- stabilizer. Their strength levels are medium to high. These alloys possess excellent combination of strength, toughness and corrosion resistance. Typical applications include blades and discs for jet engine turbines and compressors, structural aircraft components and landing gear, chemical process equipment, marine components and surgical implants. Meanwhile, β alloys contain small amounts of αstabilizing elements as strengtheners and generally weldable, high corrosion resistance and good creep resistance to intermediate temperatures. Additions of vanadium, iron and chromium as stabilizing elements, provide superior hot working characteristics. Ti-10V-2Fe-3Al, Ti-15V-3Cr-3Al-3Sn, Ti-15Mo-2.7Nb-3Al-0.2Si and Ti-3Al-8V-6Cr-4Mo-4Zr are examples of these alloys. Typical applications include airframe components, fasteners, springs, pipe and commercial and consumer products.

#### **4. Machinability of titanium alloys**

Research works on the machinability of titanium alloys have been conducted extensively and reviewed comprehensively by several researchers. The increasing demands of titanium alloys with excellent high temperature, mechanical and chemical properties make them more difficult to machine. According to Ezugwu et al. (2003), machinability can be phrased as the difficulty to machine a particular material under a given set of the machining parameters such as cutting speed, feed rate and depth of cut. It can be rated in terms of tool life, surface quality, the reaction of cutting forces and also machining cost per part. Basically, work hardening, low thermal conductivity, abrasiveness, high strength level and high heat generated were the dominant reasons for the difficulty in machining titanium alloys. Heat is the most important factor that needs to be aware of when machining titanium alloys. Excessive heat could damage the cutting tool rapidly. The main sources of heat during machining are from the shear zone, from the tool-chip interface friction and from the tool-

Machinability of Titanium Alloys in Drilling 123

the low thermal conductivity of titanium alloys, the temperature on the rake face can be

Various types of wear can be observed when drilling titanium alloys, namely non-uniform flank wear, excessive chipping and micro-cracking. These types of wear are the dominant tool failure modes when drilling Ti-6Al-4V. The wear occurs along the drill's cutting edges or the flank faces. An increase in cutting speed led to a proportional increase in the flank wear width. The increase of flank wear rate may encourage adherence of workpiece material

During drilling of Ti-48Al-2Mn-2Nb, Mantle and co-workers (Mantle et al. (1995)) found that the workpiece material adhered to the chisel and the cutting edges. The adherence of Ti-48Al-2Mn-2Nb was thinner than Ti-6Al-4V and after verification under the SEM, they concluded that the adhered material was the main contributor to the tool failure. Titanium is highly chemically reactive with the tendency of welding onto the cutting tool during machining. In the beginning, the adhered material may protect the cutting edges from wear as shown in Fig. 8. In this figure, the adhesion occurred mainly at the cutting edge, near the periphery and on the chisel edge. However with prolonged drilling, the adhered material becomes unstable and breaks away from the tool carrying along small amount of tool

(a)

(b) Fig. 8. Adherence of workpiece materials observed at: (a) chisel edge and (b) cutting edge of

Adhered material

after drilling Ti-6Al-4V (Rahim, 2005)

and may lead to attrition wear and eventually ended up in severe chipping.

particles. This situation may lead to severe chipping on the cutting edge.

Adhered material

above 900º C even at moderate cutting speed.

workpiece interface friction. However, too much heat is not the only reason associated with tool failures. The lack of rigidity in holding the tool holder with cutting tool and workpiece can also shorten the tool life. Non-rigid setups with vibration or inconsistent cutting pressure and interrupted cuts often cause tool chipping or fracturing. Prolong machining also causes severe chipping and fracture of the tool edge.

### **5. Performance evaluation in drilling of titanium alloys**

Among the various machining processes, drilling can be considerably as the most difficult process in comparison to milling and turning. Many researchers have studied the machinability of titanium alloys in the past, especially in turning and milling operations. Although extensive investigation reports have been published, no considerable progress is being made and reported on the drilling of these alloys.

#### **5.1 Tool wear**

Tool wear of cutting tools in metal cutting accounts for a significant portion of the production costs of a component. Tool wear occurs due to the physical and chemical interaction between the cutting tool and workpiece as a result of the removal of small particles of the tool material from the edge of the cutting tool.Tool wear takes place in three stages as shown in Fig. 7 (Vaughn, 1966). Tool wear developed rapidly in the initial stage and then grew uniformly until it reached its limiting value. In the third stage, the tool wear developed rapidly and caused tool failure. Machining beyond this limit will cause catastrophic failures on the tool and usually this should be avoided.

Fig. 7. Typical stages of tool wear in machining (Vaughn, 1966)

The main problem in drilling titanium and its alloys is the rapid wear of the cutting tool. Permissible rates of metal removal are low, in spite of the low cutting forces. The inhibitor in machining titanium alloys are the high temperature generated and the unfavorable temperature distribution in the cutting tool (Ezugwu & Wang, 1997; Vaughn, 1966). Due to

workpiece interface friction. However, too much heat is not the only reason associated with tool failures. The lack of rigidity in holding the tool holder with cutting tool and workpiece can also shorten the tool life. Non-rigid setups with vibration or inconsistent cutting pressure and interrupted cuts often cause tool chipping or fracturing. Prolong machining

Among the various machining processes, drilling can be considerably as the most difficult process in comparison to milling and turning. Many researchers have studied the machinability of titanium alloys in the past, especially in turning and milling operations. Although extensive investigation reports have been published, no considerable progress is

Tool wear of cutting tools in metal cutting accounts for a significant portion of the production costs of a component. Tool wear occurs due to the physical and chemical interaction between the cutting tool and workpiece as a result of the removal of small particles of the tool material from the edge of the cutting tool.Tool wear takes place in three stages as shown in Fig. 7 (Vaughn, 1966). Tool wear developed rapidly in the initial stage and then grew uniformly until it reached its limiting value. In the third stage, the tool wear developed rapidly and caused tool failure. Machining beyond this limit will cause

Steady-state wear region

Uniform wear rate

Failure region

Accelerating wear rate

The main problem in drilling titanium and its alloys is the rapid wear of the cutting tool. Permissible rates of metal removal are low, in spite of the low cutting forces. The inhibitor in machining titanium alloys are the high temperature generated and the unfavorable temperature distribution in the cutting tool (Ezugwu & Wang, 1997; Vaughn, 1966). Due to

Cutting time

also causes severe chipping and fracture of the tool edge.

being made and reported on the drilling of these alloys.

**5.1 Tool wear** 

Tool flank wear

**5. Performance evaluation in drilling of titanium alloys** 

catastrophic failures on the tool and usually this should be avoided.

Break-in period

Fig. 7. Typical stages of tool wear in machining (Vaughn, 1966)

Rapid initial wear

the low thermal conductivity of titanium alloys, the temperature on the rake face can be above 900º C even at moderate cutting speed.

Various types of wear can be observed when drilling titanium alloys, namely non-uniform flank wear, excessive chipping and micro-cracking. These types of wear are the dominant tool failure modes when drilling Ti-6Al-4V. The wear occurs along the drill's cutting edges or the flank faces. An increase in cutting speed led to a proportional increase in the flank wear width. The increase of flank wear rate may encourage adherence of workpiece material and may lead to attrition wear and eventually ended up in severe chipping.

During drilling of Ti-48Al-2Mn-2Nb, Mantle and co-workers (Mantle et al. (1995)) found that the workpiece material adhered to the chisel and the cutting edges. The adherence of Ti-48Al-2Mn-2Nb was thinner than Ti-6Al-4V and after verification under the SEM, they concluded that the adhered material was the main contributor to the tool failure. Titanium is highly chemically reactive with the tendency of welding onto the cutting tool during machining. In the beginning, the adhered material may protect the cutting edges from wear as shown in Fig. 8. In this figure, the adhesion occurred mainly at the cutting edge, near the periphery and on the chisel edge. However with prolonged drilling, the adhered material becomes unstable and breaks away from the tool carrying along small amount of tool particles. This situation may lead to severe chipping on the cutting edge.

Fig. 8. Adherence of workpiece materials observed at: (a) chisel edge and (b) cutting edge of after drilling Ti-6Al-4V (Rahim, 2005)

Machinability of Titanium Alloys in Drilling 125

workpiece during drilling Ti-6Al-4V (Rahim & Sharif, 2009). Diffusion wear is significant at the tool-workpiece interface, especially at high cutting temperature. Due to high chemical reactivity of titanium alloys, carbon reacts readily with titanium. Therefore, the formation of

Fig. 10. EDAX Section of worn tool, showing adherent workpiece material on the cutting edge after drilling Ti-6Al-4V for 2 minutes at 45 m/min and 0.06 mm/rev (Rahim & Sharif, 2009)

Piezoelectric dynamometer is commonly used to measure the cutting force in most machining processes. The various types of the dynamometer depend on the machining process such as turning, milling, drilling or grinding. A three components dynamometer is able to measure the cutting force and feed force, especially in milling and turning operations. Meanwhile, two components dynamometer is normally used in drilling process

Comparison of thrust force against the coolant-lubricant conditions at cutting speed of 60 m/min and feed rate of 0.1 mm/rev is presented in Fig. 11 (Rahim & Sasahara, 2011). It was found that the air blow condition produced the highest thrust force in comparison to the other coolant-lubricant conditions. In contrast, the MQLPO (palm oil using MQL condition) and flood conditions exhibited comparable and the lowest thrust force among the other conditions tested. As expected, the flood condition demonstrated the lowest torque among the other conditions tested as shown in Fig.12. Through the comparison, it was found that air blow did not reduce the drilling torque as much as the other coolant-lubricant conditions. They concluded that the highest value of thrust force and torque for the air blow condition could be attributed to higher amount of friction between tool-chip interface, hence, more heat is generated during the drilling process. Furthermore, Lopez and co-

**5.2 Thrust force and torque** 

to measure the thrust force and torque.

titanium carbide occurred at the interface between the tool and work material.

Spot area

Fig. 9 shows a thermal crack on the flank face of the drill (Rahim, 2005). It can be seen that the crack line propagated perpendicularly to the cutting edge. Cracks on the cutting tool and fracture of the entire cutting edge were mainly observed when machining titanium alloys at higher cutting conditions (Ezugwu et al., 2000; Jawaid et al., 2000). Cracks usually originate from the chipped area and gradually propagate along the worn flank face. Chipping at cutting edges is attributes mainly by the generation of cyclic surface stresses during drilling, which may lead to the stress cycling results in the formation of cracks parallel on the cutting edge. The propagation of cracks with prolonged machining, leads to chipping along the cutting edge. Chipping can also occur without the presence of crack formation, especially at the initial stages of the wear progress. If cracks become very numerous, they may join and cause small fragments of the cutting edge to break away.

Fig. 9. Crack on the flank face after drilling Ti-6Al-4V for 1 minute at 55 m/min and 0.06 mm/rev (Rahim, 2005)

Cantero et al. (2005) reported on the approach in drilling Ti-6Al-4V under dry condition. Using a 6 mm diameter with TiN coated carbide drill, they recommended that speed and feed rate for drilling of Ti-6Al-4V were 50 m/min and 0.07 mm/rev respectively. Attrition and diffusion were the dominant tool wear mechanisms, especially in the helical flute of drill. With prolonged drilling, these tool wear mechanisms lead to the catastrophic failure of the drill. Attrition wear is a removal of grains or agglomerates of tool material by the adherent chip or workpiece (Dearnley and Grearson, 1986). This could be due to intermittent adhesion between the tool and the workpiece as a result of the irregular chip flow and the breaking of a partially stable built-up edge. When seizure between the tool and the workpiece is broken, small fragments of the tool can be plucked out due to weakening of the binder and transported material via the underside of the chip or by the workpiece. The presence of fatigue during machining operation can initiate cracks and also encourage cracks propagation on the tool.

Furthermore, diffusion wear is associated with the chemical affinity between the tool and workpiece materials under high temperature and pressure during machining of titanium alloys (Hartung and Kramer, 1982; Kramer, 1987). An intimate contact between the toolworkpiece interface at temperature above 800º C provides an ideal environment for diffusion of tool material across the tool-workpiece interface. The EDAX analysis (Fig. 10) confirmed that tool elements (C, Co and W) had diffused into the interface between tool-

Fig. 9 shows a thermal crack on the flank face of the drill (Rahim, 2005). It can be seen that the crack line propagated perpendicularly to the cutting edge. Cracks on the cutting tool and fracture of the entire cutting edge were mainly observed when machining titanium alloys at higher cutting conditions (Ezugwu et al., 2000; Jawaid et al., 2000). Cracks usually originate from the chipped area and gradually propagate along the worn flank face. Chipping at cutting edges is attributes mainly by the generation of cyclic surface stresses during drilling, which may lead to the stress cycling results in the formation of cracks parallel on the cutting edge. The propagation of cracks with prolonged machining, leads to chipping along the cutting edge. Chipping can also occur without the presence of crack formation, especially at the initial stages of the wear progress. If cracks become very numerous, they may join and

Fig. 9. Crack on the flank face after drilling Ti-6Al-4V for 1 minute at 55 m/min and 0.06

Cantero et al. (2005) reported on the approach in drilling Ti-6Al-4V under dry condition. Using a 6 mm diameter with TiN coated carbide drill, they recommended that speed and feed rate for drilling of Ti-6Al-4V were 50 m/min and 0.07 mm/rev respectively. Attrition and diffusion were the dominant tool wear mechanisms, especially in the helical flute of drill. With prolonged drilling, these tool wear mechanisms lead to the catastrophic failure of the drill. Attrition wear is a removal of grains or agglomerates of tool material by the adherent chip or workpiece (Dearnley and Grearson, 1986). This could be due to intermittent adhesion between the tool and the workpiece as a result of the irregular chip flow and the breaking of a partially stable built-up edge. When seizure between the tool and the workpiece is broken, small fragments of the tool can be plucked out due to weakening of the binder and transported material via the underside of the chip or by the workpiece. The presence of fatigue during machining operation can initiate cracks and also encourage

Furthermore, diffusion wear is associated with the chemical affinity between the tool and workpiece materials under high temperature and pressure during machining of titanium alloys (Hartung and Kramer, 1982; Kramer, 1987). An intimate contact between the toolworkpiece interface at temperature above 800º C provides an ideal environment for diffusion of tool material across the tool-workpiece interface. The EDAX analysis (Fig. 10) confirmed that tool elements (C, Co and W) had diffused into the interface between tool-

cause small fragments of the cutting edge to break away.

Crack

mm/rev (Rahim, 2005)

cracks propagation on the tool.

workpiece during drilling Ti-6Al-4V (Rahim & Sharif, 2009). Diffusion wear is significant at the tool-workpiece interface, especially at high cutting temperature. Due to high chemical reactivity of titanium alloys, carbon reacts readily with titanium. Therefore, the formation of titanium carbide occurred at the interface between the tool and work material.

Fig. 10. EDAX Section of worn tool, showing adherent workpiece material on the cutting edge after drilling Ti-6Al-4V for 2 minutes at 45 m/min and 0.06 mm/rev (Rahim & Sharif, 2009)

#### **5.2 Thrust force and torque**

Piezoelectric dynamometer is commonly used to measure the cutting force in most machining processes. The various types of the dynamometer depend on the machining process such as turning, milling, drilling or grinding. A three components dynamometer is able to measure the cutting force and feed force, especially in milling and turning operations. Meanwhile, two components dynamometer is normally used in drilling process to measure the thrust force and torque.

Comparison of thrust force against the coolant-lubricant conditions at cutting speed of 60 m/min and feed rate of 0.1 mm/rev is presented in Fig. 11 (Rahim & Sasahara, 2011). It was found that the air blow condition produced the highest thrust force in comparison to the other coolant-lubricant conditions. In contrast, the MQLPO (palm oil using MQL condition) and flood conditions exhibited comparable and the lowest thrust force among the other conditions tested. As expected, the flood condition demonstrated the lowest torque among the other conditions tested as shown in Fig.12. Through the comparison, it was found that air blow did not reduce the drilling torque as much as the other coolant-lubricant conditions. They concluded that the highest value of thrust force and torque for the air blow condition could be attributed to higher amount of friction between tool-chip interface, hence, more heat is generated during the drilling process. Furthermore, Lopez and co-

Machinability of Titanium Alloys in Drilling 127

Fig. 13. Comparison of thrust force when drilling Ti-6Al-4V using cemented carbide tool

Fig. 14. Comparison of torque when drilling Ti-6Al-4V using cemented carbide tool under

Some researchers have tested several techniques in drilling of titanium alloys. A step feed drilling or intermittently decelerated feed drilling and vibratory drilling were conducted by Sakurai co-wrokers (Sakurai et al., 1992; Sakurai et al., 1996) to examine the cutting force and cutting characteristic of TiN coated cobalt HSS and oxide treatment nitridized cobalt HSS when drilling Ti-6Al-4V. Results of their study showed that step feed drilling contributed a lower thrust force and torque as compared to continuous conventional drilling. In addition, the thrust force and torque on TiN drills are lower than oxide treatment nitridized drills in both conventional and step feed drilling. As reported by Okamura and co-workers (Okamura et al., 2006), the non-vibration drilling shows a tremendous reduction on thrust force. However, the value tends to decrease once the vibration exceeded 20 kHz. It is believed that the natural frequency of measurement systems does not exceed the vibrating frequency. In another work by Rahim and co-workers (Rahim et al., 2008) showed that

pecking drilling method significantly reduces the thrust force and torque.

under flood coolant condition (Rahim et al., 2008)

flood coolant condition (Rahim et al., 2008)

workers found that the cutting force produced by high pressure internal cooling method was lower compared with the external cooling, which has a beneficial effect on workpiece deformation and hole quality (Lopez et al. , 2000).

Fig. 11. Thrust force when high speed drilling of Ti-6Al-4V under various coolant-lubricant conditions (Rahim & Sasahara, 2011)

The influence of drilling parameters has been assessed for different material characteristic and properties of titanium alloys (Mantle et al. ,1995). Result shows that the thrust force and torque for Ti-48Al-2Mn-2Nb were greater than Ti-6Al-4V. As shown in Figs. 13 and 14, the thrust force decreases as cutting speed increases (Rahim et al., 2008). At the same time, results also showed that low torque values were obtained at the highest cutting speed. This behavior is attributed to the reduction of the contact area between the tool-workpiece interface and the reduction of specific cutting energy. Moreover, with increase of cutting speed, the cutting temperature increases, subsequently reduced the material hardness. As a result, both the thrust force and torque are reduced. Meanwhile, the thrust force and torque values were significantly increased when the feed rate was increased as shown in Fig. 15 (Rahim & Sasahara, 2011). The thrust force and torque are strongly correlated with the chip thickness, which is associated with the feed rate (Liao et al., 2007). This is because high feed rate results in a larger cross sectional area of the undeformed chip, and, consequently, greater thrust force and torque are produced.

Fig. 12. Torque when high speed drilling of Ti-6Al-4V under various coolant-lubricant conditions (Rahim & Sasahara, 2011)

workers found that the cutting force produced by high pressure internal cooling method was lower compared with the external cooling, which has a beneficial effect on workpiece

Fig. 11. Thrust force when high speed drilling of Ti-6Al-4V under various coolant-lubricant

The influence of drilling parameters has been assessed for different material characteristic and properties of titanium alloys (Mantle et al. ,1995). Result shows that the thrust force and torque for Ti-48Al-2Mn-2Nb were greater than Ti-6Al-4V. As shown in Figs. 13 and 14, the thrust force decreases as cutting speed increases (Rahim et al., 2008). At the same time, results also showed that low torque values were obtained at the highest cutting speed. This behavior is attributed to the reduction of the contact area between the tool-workpiece interface and the reduction of specific cutting energy. Moreover, with increase of cutting speed, the cutting temperature increases, subsequently reduced the material hardness. As a result, both the thrust force and torque are reduced. Meanwhile, the thrust force and torque values were significantly increased when the feed rate was increased as shown in Fig. 15 (Rahim & Sasahara, 2011). The thrust force and torque are strongly correlated with the chip thickness, which is associated with the feed rate (Liao et al., 2007). This is because high feed rate results in a larger cross sectional area of the undeformed chip, and, consequently,

Fig. 12. Torque when high speed drilling of Ti-6Al-4V under various coolant-lubricant

deformation and hole quality (Lopez et al. , 2000).

conditions (Rahim & Sasahara, 2011)

greater thrust force and torque are produced.

conditions (Rahim & Sasahara, 2011)

Fig. 13. Comparison of thrust force when drilling Ti-6Al-4V using cemented carbide tool under flood coolant condition (Rahim et al., 2008)

Fig. 14. Comparison of torque when drilling Ti-6Al-4V using cemented carbide tool under flood coolant condition (Rahim et al., 2008)

Some researchers have tested several techniques in drilling of titanium alloys. A step feed drilling or intermittently decelerated feed drilling and vibratory drilling were conducted by Sakurai co-wrokers (Sakurai et al., 1992; Sakurai et al., 1996) to examine the cutting force and cutting characteristic of TiN coated cobalt HSS and oxide treatment nitridized cobalt HSS when drilling Ti-6Al-4V. Results of their study showed that step feed drilling contributed a lower thrust force and torque as compared to continuous conventional drilling. In addition, the thrust force and torque on TiN drills are lower than oxide treatment nitridized drills in both conventional and step feed drilling. As reported by Okamura and co-workers (Okamura et al., 2006), the non-vibration drilling shows a tremendous reduction on thrust force. However, the value tends to decrease once the vibration exceeded 20 kHz. It is believed that the natural frequency of measurement systems does not exceed the vibrating frequency. In another work by Rahim and co-workers (Rahim et al., 2008) showed that pecking drilling method significantly reduces the thrust force and torque.

Machinability of Titanium Alloys in Drilling 129

Fig. 17. System for measurement of the temperature in the workpiece (Zeilmann &

increases the friction and stresses, thus increasing the cutting temperature.

Cutting speed and feed rate are among the factors that contribute to the variation of temperature during drilling titanium alloys. The cutting temperature increases with the cutting speed. This corresponds with the high cutting energy, deformation strain rate as well as the heat flux (Rahim & Sasahara, 2011). Furthermore, drilling at high feed rate

The application of different cooling methods provides a variation in temperature results. For example, the maximum temperatures recorded for drilling with abundant emulsion through the interior of the tool stayed in the range of 22–32% of the values obtained with the application of MQL with an external nozzle as shown in Fig. 20 (Zeilmann & Weingartner, 2006). Comparing drilling with MQL applied with an external nozzle and dry drilling, the values obtained for the second condition were approximately 6% superior, ranging from 455 to 482 °C. Furthermore, flood and MQL conditions recorded a low workpiece temperature in comparison to the air blow condition as shown in Fig. 21 (Rahim & Sasahara, 2011).

Fig. 18. The thermocouple was inserted through the oil hole of internal coolant carbide drill

Weingaertner, 2006)

(Ozcelik & Bagci, 2006)

Fig. 15. (a) Thrust force and (b) torque for MQLSE and MQLPO under various cutting speed and feed rate (Rahim & Sasahara, 2011)

#### **5.3 Temperature**

Embedded thermocouples were one of the earliest technique used for the estimation of temperatures in various manufacturing and tribological applications. In order to use this technique, particularly in machining, a number of fine deep holes have to be made in the stationary part, namely the workpiece or the cutting tool, and the thermocouples are inserted in different locations in the interior of the part, with some of them as close to the surface as possible. In drilling process, the measurement of temperature by thermocouple wires can be done by embedding the wires in the workpiece and cutting tool as shown in Figs. 16, 17, 18 and 19, respectively. These methods are able to measure the workpiece and cutting tool temperature, especially when drilling titanium alloys.

Fig. 16. Thermocouple locations (Rahim & Sasahara, 2010a)

Fig. 15. (a) Thrust force and (b) torque for MQLSE and MQLPO under various cutting speed

Embedded thermocouples were one of the earliest technique used for the estimation of temperatures in various manufacturing and tribological applications. In order to use this technique, particularly in machining, a number of fine deep holes have to be made in the stationary part, namely the workpiece or the cutting tool, and the thermocouples are inserted in different locations in the interior of the part, with some of them as close to the surface as possible. In drilling process, the measurement of temperature by thermocouple wires can be done by embedding the wires in the workpiece and cutting tool as shown in Figs. 16, 17, 18 and 19, respectively. These methods are able to measure the workpiece and

cutting tool temperature, especially when drilling titanium alloys.

Fig. 16. Thermocouple locations (Rahim & Sasahara, 2010a)

and feed rate (Rahim & Sasahara, 2011)

**5.3 Temperature** 

Fig. 17. System for measurement of the temperature in the workpiece (Zeilmann & Weingaertner, 2006)

Cutting speed and feed rate are among the factors that contribute to the variation of temperature during drilling titanium alloys. The cutting temperature increases with the cutting speed. This corresponds with the high cutting energy, deformation strain rate as well as the heat flux (Rahim & Sasahara, 2011). Furthermore, drilling at high feed rate increases the friction and stresses, thus increasing the cutting temperature.

The application of different cooling methods provides a variation in temperature results. For example, the maximum temperatures recorded for drilling with abundant emulsion through the interior of the tool stayed in the range of 22–32% of the values obtained with the application of MQL with an external nozzle as shown in Fig. 20 (Zeilmann & Weingartner, 2006). Comparing drilling with MQL applied with an external nozzle and dry drilling, the values obtained for the second condition were approximately 6% superior, ranging from 455 to 482 °C. Furthermore, flood and MQL conditions recorded a low workpiece temperature in comparison to the air blow condition as shown in Fig. 21 (Rahim & Sasahara, 2011).

Fig. 18. The thermocouple was inserted through the oil hole of internal coolant carbide drill (Ozcelik & Bagci, 2006)

Machinability of Titanium Alloys in Drilling 131

As reported by Pujana and co-workers (Pujana et al., 2009), the cutting temperature was higher when using ultrasonic-assisted drilling in comparison to the non-vibration drilling. In this case, the higher the vibration amplitude, the higher the temperature variations. Okamura and co-authors (Okamura et al., 2006) have designed a low-frequency vibration drilling machine to drill a Ti-6Al-4V. They described the effect of low-frequency vibration drilling on cutting temperature. Results showed that, higher amplitude of 0.24 mm and frequency of 30 Hz exhibited lower cutting temperature as compared to non-vibration

Surface integrity is defined as the unimpaired or enhanced surface condition of a material resulting from the impact of a controlled manufacturing process (Field and Kahles, 1964). Damaged layer and surface integrity of the finished surface significantly influence the wear resistance, corrosion resistance and fatigue strength of the machined components. Surface integrity produced by metal removal operation can be categorized as geometrical surface integrity and physical surface integrity. To find the impact of the manufacturing process on the material properties both categories effects must be considered. Surface integrity aspects are very important, especially in aerospace industry with respect to the high degree of safety. Surface integrity is concerned primarily on the effect of the machining process on the changes in surface and sub-surface of the component which are categorized as surface

There are three essential parameters in a surface roughness; arithmetical mean deviation of the profile (Ra), maximum height of the profile (Rmax) and height of the profile irregularities in ten points (Rz). It is believed that the higher surface roughness value is responsible for the decrease of the fatigue strength on the machined surface. Significant improvement in surface roughness can be obtained when low feed rate and high cutting speed are employed. However, the response of surface roughness towards cutting speed was less significant when compared to feed rate. Sun and Guo (Sun & Guo, 2009), reported that surface roughness value increased with increase in feed rate and radial

Previous study showed that surface roughness value is lower at high cutting speed when drilling Ti-6Al-4V using carbide drills (Sharif & Rahim, 2007). During machining at high cutting speed, the cutting temperature increases due the small contact length between toolworkpiece interfaces. This could be due to the decrease in the value of coefficient of friction, which results in low friction at the tool-workpiece interface. These factors could contribute to the improvement in surface roughness values as shown in Fig. 22 (Rahim & Sasahara, 2010b). In addition, as the cutting speed increases, more heat is generated thus softening the workpiece material, which in turn improves the surface roughness. However, a low cutting speed may lead to the formation of built-up edge and hence deteriorates the machined surface. Investigation revealed that at high feed rate the surface roughness is poor, probably

due to the distinct feed marks produced at high feed rate (Rahim & Sasahara, 2010)

roughness, plastic deformation, residual stress and microhardness.

drilling.

**6. Surface integrity** 

**6.1 Surface roughness** 

depth-of-cut.

Fig. 19. Top view and coordinates of thermocouple tips on drill flank face (Li & Shih, 2007)

Fig. 20. Maximum temperature in the piece for different cutting fluids conditions when drilling Ti-6Al-4V using grade K10 cemented carbide tool (Zeilmann & Weingartner, 2006)

Fig. 21. Maximum workpiece temperature in high speed drilling of Ti-6Al-4V under various coolant-lubricant conditions (Rahim & Sasahara, 2011)

As reported by Pujana and co-workers (Pujana et al., 2009), the cutting temperature was higher when using ultrasonic-assisted drilling in comparison to the non-vibration drilling. In this case, the higher the vibration amplitude, the higher the temperature variations. Okamura and co-authors (Okamura et al., 2006) have designed a low-frequency vibration drilling machine to drill a Ti-6Al-4V. They described the effect of low-frequency vibration drilling on cutting temperature. Results showed that, higher amplitude of 0.24 mm and frequency of 30 Hz exhibited lower cutting temperature as compared to non-vibration drilling.

#### **6. Surface integrity**

130 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Fig. 19. Top view and coordinates of thermocouple tips on drill flank face (Li & Shih, 2007)

Fig. 20. Maximum temperature in the piece for different cutting fluids conditions when drilling Ti-6Al-4V using grade K10 cemented carbide tool (Zeilmann & Weingartner, 2006)

Fig. 21. Maximum workpiece temperature in high speed drilling of Ti-6Al-4V under various

coolant-lubricant conditions (Rahim & Sasahara, 2011)

Surface integrity is defined as the unimpaired or enhanced surface condition of a material resulting from the impact of a controlled manufacturing process (Field and Kahles, 1964). Damaged layer and surface integrity of the finished surface significantly influence the wear resistance, corrosion resistance and fatigue strength of the machined components. Surface integrity produced by metal removal operation can be categorized as geometrical surface integrity and physical surface integrity. To find the impact of the manufacturing process on the material properties both categories effects must be considered. Surface integrity aspects are very important, especially in aerospace industry with respect to the high degree of safety. Surface integrity is concerned primarily on the effect of the machining process on the changes in surface and sub-surface of the component which are categorized as surface roughness, plastic deformation, residual stress and microhardness.

#### **6.1 Surface roughness**

There are three essential parameters in a surface roughness; arithmetical mean deviation of the profile (Ra), maximum height of the profile (Rmax) and height of the profile irregularities in ten points (Rz). It is believed that the higher surface roughness value is responsible for the decrease of the fatigue strength on the machined surface. Significant improvement in surface roughness can be obtained when low feed rate and high cutting speed are employed. However, the response of surface roughness towards cutting speed was less significant when compared to feed rate. Sun and Guo (Sun & Guo, 2009), reported that surface roughness value increased with increase in feed rate and radial depth-of-cut.

Previous study showed that surface roughness value is lower at high cutting speed when drilling Ti-6Al-4V using carbide drills (Sharif & Rahim, 2007). During machining at high cutting speed, the cutting temperature increases due the small contact length between toolworkpiece interfaces. This could be due to the decrease in the value of coefficient of friction, which results in low friction at the tool-workpiece interface. These factors could contribute to the improvement in surface roughness values as shown in Fig. 22 (Rahim & Sasahara, 2010b). In addition, as the cutting speed increases, more heat is generated thus softening the workpiece material, which in turn improves the surface roughness. However, a low cutting speed may lead to the formation of built-up edge and hence deteriorates the machined surface. Investigation revealed that at high feed rate the surface roughness is poor, probably due to the distinct feed marks produced at high feed rate (Rahim & Sasahara, 2010)

Machinability of Titanium Alloys in Drilling 133

Fig. 23 also shows that there is hardening layer below the softened layer whose hardness depends on the cutting parameters (i.e cutting speed, feed rate, depth of cut) as well as mechanical and thermal interaction. It was generally observed in the work hardening region that the microhardness increases with increase in cutting speed and feed rate. An increase in microhardness of the surface layer, as a result of high feed rate could be associated by the high rubbing load between the tool and the machined surface and the consequent work

Fig. 23. Sub-surface hardness variations after drilling Ti-6Al-4V using MQLSE (Rahim &

the transformation of beta phase to alpha phase during drilling (Cantero et al., 2005).

Fig. 24. Microhardness variation beneath the surface produced when drilling Ti-5Al-4V-

In another work by Rahim and Sharif (Rahim & Sharif, 2006), it was reported that the hardness value underneath the drilled surface was higher than the average hardness of the bulk material when drilling Ti-5Al-4V-Mo/Fe (Fig. 24). Meanwhile, a significant changed of microhardness values were also observed underneath the machined surface. It was due to

hardening effect.

Sasahara, 2010b)

Mo/Fe (Rahim & Sharif, 2006)

Fig. 22. Comparison of surface roughness level obtained when drilling Ti-6Al-4V using TiAlN coated carbide tool under MQLSE and MQLPO (Rahim & Sasahara, 2010b)

Types of cutting fluid also influence the surface roughness of the machined surface. Under the MQL condition, vegetable oil (MQLPO: palm oil) exhibits better surface roughness than synthetic ester (MQLSE) as shown in Fig. 22 (Rahim & Sasahara, 2010b). It can be suggested that less heat is generated using palm oil thus provided enough time to cool and lubricate the tool-workpiece interface. Apparently, such reduction may attribute to better lubrication and shorter tool-chip contact length during drilling. Moreover, surface roughness measured by peck drilling is far better than conventional drilling method (Rahim et al, 2008).

#### **6.2 Microhardness**

The microhardness alterations observed during machining may be due to the effect of thermal, mechanical and chemical reaction. Many researchers believed that the workpiece material is subjected to work hardening and thermal softening effect during machining, especially at high cutting temperature and pressure (Che Haron, 2001; Ginting & Nouari, 2009). When machining titanium alloys, the hardness just beneath the machined surface was found to be softer than the bulk material hardness due to the thermal softening effect. However, when the depth below the machined surface increases, the hardness value starts to increase before reaching its peak value and finally drops gradually to the bulk material hardness as shown in Fig. 23. The increase in hardness value is directly associated with the effect of work hardening. This effect depends on the temperature, cutting time and the mechanism of internal stress relaxation (Ginting and Nouari, 2009).

Fig. 23 shows that the microhardness of the sub-surface at 0.025 mm underneath the machined surface was below the average base material hardness. This indicates that the machined surface experienced thermal softening effect or over aging due to the localized heating during the drilling process (Rahim & Sasahara, 2010b). Turning test of titanium alloys by Haron and Jawaid (Haron & Jawaid, 2005) and Ginting and Nouari (Ginting & Nouari, 2009) have also indicated significant drop of microhardness value near the surface of machined layer. They pointed out that the existence of high cutting temperature and high cutting pressure produced noticeable softening in the surface region.

Fig. 22. Comparison of surface roughness level obtained when drilling Ti-6Al-4V using TiAlN coated carbide tool under MQLSE and MQLPO (Rahim & Sasahara, 2010b)

by peck drilling is far better than conventional drilling method (Rahim et al, 2008).

mechanism of internal stress relaxation (Ginting and Nouari, 2009).

cutting pressure produced noticeable softening in the surface region.

**6.2 Microhardness** 

Types of cutting fluid also influence the surface roughness of the machined surface. Under the MQL condition, vegetable oil (MQLPO: palm oil) exhibits better surface roughness than synthetic ester (MQLSE) as shown in Fig. 22 (Rahim & Sasahara, 2010b). It can be suggested that less heat is generated using palm oil thus provided enough time to cool and lubricate the tool-workpiece interface. Apparently, such reduction may attribute to better lubrication and shorter tool-chip contact length during drilling. Moreover, surface roughness measured

The microhardness alterations observed during machining may be due to the effect of thermal, mechanical and chemical reaction. Many researchers believed that the workpiece material is subjected to work hardening and thermal softening effect during machining, especially at high cutting temperature and pressure (Che Haron, 2001; Ginting & Nouari, 2009). When machining titanium alloys, the hardness just beneath the machined surface was found to be softer than the bulk material hardness due to the thermal softening effect. However, when the depth below the machined surface increases, the hardness value starts to increase before reaching its peak value and finally drops gradually to the bulk material hardness as shown in Fig. 23. The increase in hardness value is directly associated with the effect of work hardening. This effect depends on the temperature, cutting time and the

Fig. 23 shows that the microhardness of the sub-surface at 0.025 mm underneath the machined surface was below the average base material hardness. This indicates that the machined surface experienced thermal softening effect or over aging due to the localized heating during the drilling process (Rahim & Sasahara, 2010b). Turning test of titanium alloys by Haron and Jawaid (Haron & Jawaid, 2005) and Ginting and Nouari (Ginting & Nouari, 2009) have also indicated significant drop of microhardness value near the surface of machined layer. They pointed out that the existence of high cutting temperature and high Fig. 23 also shows that there is hardening layer below the softened layer whose hardness depends on the cutting parameters (i.e cutting speed, feed rate, depth of cut) as well as mechanical and thermal interaction. It was generally observed in the work hardening region that the microhardness increases with increase in cutting speed and feed rate. An increase in microhardness of the surface layer, as a result of high feed rate could be associated by the high rubbing load between the tool and the machined surface and the consequent work hardening effect.

Fig. 23. Sub-surface hardness variations after drilling Ti-6Al-4V using MQLSE (Rahim & Sasahara, 2010b)

In another work by Rahim and Sharif (Rahim & Sharif, 2006), it was reported that the hardness value underneath the drilled surface was higher than the average hardness of the bulk material when drilling Ti-5Al-4V-Mo/Fe (Fig. 24). Meanwhile, a significant changed of microhardness values were also observed underneath the machined surface. It was due to the transformation of beta phase to alpha phase during drilling (Cantero et al., 2005).

Fig. 24. Microhardness variation beneath the surface produced when drilling Ti-5Al-4V-Mo/Fe (Rahim & Sharif, 2006)

Machinability of Titanium Alloys in Drilling 135

Creditable works have been carried out by many researchers in drilling of titanium alloys which resulted in significant improvements in the productivity of titanium parts. The application of drilling strategies, introduction of newly developed tool geometry and coolant conditions have improved the surface integrity and increased the tool life performance of the cutting tools in several folds. In general the following conclusions can be

i. Adhesion, attrition and diffusion are the operating tool wear mechanisms when drilling

ii. Flank wear, excessive chipping, cracking and tool breakage are the dominant tool

iii. The values of thrust force and torque decrease with increase in cutting speed. In

iv. Cutting speed and feed rate significantly affect the surface roughness of the machined surface whereby high cutting speed and low feed rate resulted in the better surface

v. Under various coolant-lubricant conditions, air blow produces higher cutting

vi. The machined surface deteriorates due to the effect of metallurgical changes and surface quality during drilling at high cutting speed, feed rate and under various

The authors wish to thank the Ministry of Higher Education of Malaysia and Research Management Center, UTM for their financial supports to this work through the Research University Grant (RUG) funding number Q.J130000.7124.02H43. Special gratitude is also extended to UTHM and TUAT, Japan and Technology and Mitsubishi Materials

Boyer, R. R. (1996). An Overview on the Use of Titanium in the Aerospace Industry. *Material* 

Brewer, W. D., Bird, R. K., & Wallace, T. A. (1998). Titanium alloys and Processing for High

Brinksmeier, E. (1990). Prediction of Tool Fracture in Drilling. *Annals of CIRP*, Vol. 39, No. 1,

Cantero, J. L., Tardio, M. M., Canteli, J.A., Marcos, M., & Miguelez, M. H., (2005). Dry

Choudhury, S. K., & Raju, G. (2000). Investigation into Crater Wear in Drilling. *Int. J. Mach.* 

Dolinsek, S., Sustarsic, B., & Kopac, J. (2001). Wear Mechanism of Cutting Tools in High

Speed Cutting Process. *Wear*, Vol. 250, pp. 349-352, ISSN 0043-1648

Speed Aircraft. *Materials Engineering & Science A*, Vol. 243, pp. 299-304, ISSN 0921-

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*Tools & Manuf*., Vol. 40, pp. 887-898, ISSN 0890-6955

contrast, these values increase significantly when the feed rate is increased.

**7. Conclusions** 

titanium alloy.

failure modes.

coolant conditions.

**8. Acknowledgments** 

**9. References** 

5093

0890-6955

finish.

drawn when drilling titanium alloys:

temperature as compared to other condistions.

Corporation for providing the cutting tools.

pp. 97-100, ISSN 0007-8506

#### **6.3 Sub-surface plastic deformation**

It is discernible that the surface and sub-surface of the machined surface are subjected to plastically deform. Sub-surface plastic deformation is particularly due to the effect of large strain, strain rate and temperature. In addition, a freshly cut surface may be burnished by a dull cutting tool, hence work hardened the machined surface. Jeelani and Ramakrishnan (Jeelani & Ramakrishnan, 1983) observed that the machined surface is severely damaged with the plastic flow in the direction of the tool motion. Meanwhile, Velasques and coworkers (Velasques et al., 2010) found that the severe deformation beneath the machined surface is associated with high cutting speed. Sub-surface plastic deformation area can be divided into three zones, namely highly perturbed region, a plastically deformed layer and unaffected zone. Normally, the sub-surface plastic deformations of microstructure of the machined surfaces are examined under the high magnification microscope in etched condition.

In most cases, plastic deformation occurs towards the spindle rotational direction. When drilling titanium alloys at higher cutting speeds and feed rates, a thicker plastic deformation can be observed. At this condition, the temperature between tool-chip interface increases thus sticking friction region occurred. Therefore, the combination of high cutting temperature and sticking friction contributed to the severe and noticeable subsurface plastic deformation (Rahim & Sasahara, 2010).

Fig. 25 shows an evidence of sub-surface plastic deformation when drilling Ti-5Al-4V-Mo/Fe (Rahim & Sharif, 2006). In this figure, the deformation is found to be severe after prolonged drilling. In this case, no white layer especially on the top of the machined surface is observed. The authors stated that high cutting force and temperature are the dominant factors which lead to the severe plastic deformation. Cantero and co-workers (Cantero et al., 2005) also found the same phenomenon and they concluded that the plastic deformations during machining are caused by mechanical forces from the cutting tool acting upon the work-piece. Additional deformation can occur as a consequence of temperature gradients due to localized heating of the machined surface area.

Fig. 25. Magnified view of the machined sub-surface when drilling Ti-5Al-4V-Mo/Fe (Rahim & Sharif, 2006)

#### **7. Conclusions**

134 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

It is discernible that the surface and sub-surface of the machined surface are subjected to plastically deform. Sub-surface plastic deformation is particularly due to the effect of large strain, strain rate and temperature. In addition, a freshly cut surface may be burnished by a dull cutting tool, hence work hardened the machined surface. Jeelani and Ramakrishnan (Jeelani & Ramakrishnan, 1983) observed that the machined surface is severely damaged with the plastic flow in the direction of the tool motion. Meanwhile, Velasques and coworkers (Velasques et al., 2010) found that the severe deformation beneath the machined surface is associated with high cutting speed. Sub-surface plastic deformation area can be divided into three zones, namely highly perturbed region, a plastically deformed layer and unaffected zone. Normally, the sub-surface plastic deformations of microstructure of the machined surfaces are examined under the high magnification microscope in etched

In most cases, plastic deformation occurs towards the spindle rotational direction. When drilling titanium alloys at higher cutting speeds and feed rates, a thicker plastic deformation can be observed. At this condition, the temperature between tool-chip interface increases thus sticking friction region occurred. Therefore, the combination of high cutting temperature and sticking friction contributed to the severe and noticeable subsurface plastic

Fig. 25 shows an evidence of sub-surface plastic deformation when drilling Ti-5Al-4V-Mo/Fe (Rahim & Sharif, 2006). In this figure, the deformation is found to be severe after prolonged drilling. In this case, no white layer especially on the top of the machined surface is observed. The authors stated that high cutting force and temperature are the dominant factors which lead to the severe plastic deformation. Cantero and co-workers (Cantero et al., 2005) also found the same phenomenon and they concluded that the plastic deformations during machining are caused by mechanical forces from the cutting tool acting upon the work-piece. Additional deformation can occur as a consequence of temperature gradients

Fig. 25. Magnified view of the machined sub-surface when drilling Ti-5Al-4V-Mo/Fe

**6.3 Sub-surface plastic deformation** 

deformation (Rahim & Sasahara, 2010).

(Rahim & Sharif, 2006)

due to localized heating of the machined surface area.

condition.

Creditable works have been carried out by many researchers in drilling of titanium alloys which resulted in significant improvements in the productivity of titanium parts. The application of drilling strategies, introduction of newly developed tool geometry and coolant conditions have improved the surface integrity and increased the tool life performance of the cutting tools in several folds. In general the following conclusions can be drawn when drilling titanium alloys:


#### **8. Acknowledgments**

The authors wish to thank the Ministry of Higher Education of Malaysia and Research Management Center, UTM for their financial supports to this work through the Research University Grant (RUG) funding number Q.J130000.7124.02H43. Special gratitude is also extended to UTHM and TUAT, Japan and Technology and Mitsubishi Materials Corporation for providing the cutting tools.

#### **9. References**


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**Part 3** 

**Surface Treatments of** 

**Titanium Alloys for Biomedical** 

**and Other Challenging Applications** 


## **Part 3**

**Surface Treatments of Titanium Alloys for Biomedical and Other Challenging Applications** 

138 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

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Ti6Al4V with minimal quantity of lubrication. *J. Mater. Process. Technol.*, Vol. 179,

**7** 

*Ukraine* 

**Chemico-Thermal Treatment** 

**of Titanium Alloys – Nitriding** 

*Physical-Mechanical Institute of National Academy of Sciences of Ukraine* 

Titanium and its alloys are widely used in aircraft, rocket production, shipbuilding, machine industry, chemical and food industry, medicine due to their high specific strength, good corrosion resistance and biological passivity. However, titanium has properties which limit its application as construction material. Particularly, tendency to surface adhesion and galling at friction results in its lowest wear resistance among construction materials. The application of titanium alloys in friction units and in the places of direct contact is impossible without additional surface treatment for higher strength. The corrosion resistance of titanium alloys is not often satisfactory. Therefore titanium alloys also need the

The chemical heat treatment, particularly nitriding, allows to extend the functionality of titanium alloys, enhancing the wear resistance and providing the high anticorrosion characteristics in aggressive media. However, the molecular nitrogen is a reactionless gas as a result of significant bond strength in molecule (Но=940 kJ/mole). Therefore it is very important to intensify the interaction between titanium and nitrogen and to elaborate the

**2. Basic regularities of high-temperature interaction of titanium alloys with** 

One of the ways to solve the problem of intensification of nitriding of titanium alloys is the high-temperature saturation based on temperature dependence of diffusion coefficient of nitrogen in titanium. Therefore we will consider the basic regularities of nitriding of

**2.1 Kinetic regularities of high-temperature interaction of titanium alloys with nitrogen**  Kinetics of nitrogen absorption by titanium alloys was widely studied by both the thermogravimetric analysis, when the mass change after different exposures in nitrogen at constant temperature is fixed and the manometric method, when the change of nitrogen pressure in the closed system is determined. The results of these studies in a wide temperature range (550...1600 oС) showed that the process of nitrogen absorption by

**1. Introduction** 

additional protection in the aggressive media.

relevant nitriding methods.

titanium alloys at high temperatures.

**nitrogen** 

Iryna Pohrelyuk and Viktor Fedirko

## **Chemico-Thermal Treatment of Titanium Alloys – Nitriding**

Iryna Pohrelyuk and Viktor Fedirko *Physical-Mechanical Institute of National Academy of Sciences of Ukraine Ukraine* 

### **1. Introduction**

Titanium and its alloys are widely used in aircraft, rocket production, shipbuilding, machine industry, chemical and food industry, medicine due to their high specific strength, good corrosion resistance and biological passivity. However, titanium has properties which limit its application as construction material. Particularly, tendency to surface adhesion and galling at friction results in its lowest wear resistance among construction materials. The application of titanium alloys in friction units and in the places of direct contact is impossible without additional surface treatment for higher strength. The corrosion resistance of titanium alloys is not often satisfactory. Therefore titanium alloys also need the additional protection in the aggressive media.

The chemical heat treatment, particularly nitriding, allows to extend the functionality of titanium alloys, enhancing the wear resistance and providing the high anticorrosion characteristics in aggressive media. However, the molecular nitrogen is a reactionless gas as a result of significant bond strength in molecule (Но=940 kJ/mole). Therefore it is very important to intensify the interaction between titanium and nitrogen and to elaborate the relevant nitriding methods.

#### **2. Basic regularities of high-temperature interaction of titanium alloys with nitrogen**

One of the ways to solve the problem of intensification of nitriding of titanium alloys is the high-temperature saturation based on temperature dependence of diffusion coefficient of nitrogen in titanium. Therefore we will consider the basic regularities of nitriding of titanium alloys at high temperatures.

#### **2.1 Kinetic regularities of high-temperature interaction of titanium alloys with nitrogen**

Kinetics of nitrogen absorption by titanium alloys was widely studied by both the thermogravimetric analysis, when the mass change after different exposures in nitrogen at constant temperature is fixed and the manometric method, when the change of nitrogen pressure in the closed system is determined. The results of these studies in a wide temperature range (550...1600 oС) showed that the process of nitrogen absorption by titanium is described by the parabolic dependence. That is the dependence of relative mass increase of the nitrided samples on time can be presented by a function:

$$(\Delta m/\text{S})^2 = K\,\tau,\tag{1}$$

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 143

change of mass increase of the nitrided samples of titanium alloys in time determined by treatment of nitriding isotherms of each investigated alloy in the double logarithmical

near to 2 that corresponds to the parabolic law of interaction (1) as well as for the unalloyed

The calculated constants (*K, Ko, E*) of the dependences *(1, 2)* at the use of corresponding mathematical models (Matychak et al., 2007, 2008, 2009, 2011) at the predetermined temperature and duration of saturation process allow to forecast the thickness of the phases formed during the nitriding process or determine time and temperature parameters of

The results of investigation of kinetics of saturation process of titanium alloys by nitrogen

Fig. 1. Kinetics of nitriding of titanium alloys : a - 950 оС; b - 900 оС; c - 1000 оС; d - 1100 оС; 1 - VТ1- 0; 2 - PТ-7М; 3 - VТ5- 1; 4 - ОТ4- 1; 5 - VТ6s; 6 - VТ5; 7 - VТ1-D; 8 - VТ3- 1; 9 - VТ6; 10 -

With the rise of temperature of isothermal exposure the influence of alloying elements on the nitriding process increases that confirmed by the increase of distance between the

nitriding at which the thicknesses of the formed layers would be predetermined.

testify that alloying influences on nitriding rate (fig. 1).

(cotangent of angle of inclination of the received straight lines) are

coordinates *ln(*

VТ20; 11 - VТ23.

titanium.

*m/S)-ln*

where *m = (m2 - m1)* – difference between sample's mass after and before nitriding; S – nitriding surface square; – nitriding duration; K – parabolic constant of nitriding rate.

The appreciable deviations from the parabolic law caused by the presence of negligible quantity of oxygen impurities in nitrogen are observed at the initial stage of reaction.

With the rise of temperature during isothermal exposure the intensity of interaction of titanium with nitrogen increases substantially. The parabolic constant of nitriding rate (K) is determined by the tangent of angle of inclination of straight lines of time dependence of square of mass increase. It characterizes quantitatively the relative intensity of saturation process (table 1). Values K increase practically on order with the rise of temperature on one hundred degrees (for example, at 900 оС *K* = 8,610-12 g2/сm4s and at 1000 оС *K =* 3,510-11 g2/сm4s).


Table 1. Parameters of temperature dependence of parabolic constant of nitriding rate К

The temperature dependence of parabolic constant of nitriding rate is described by Arrhenius equation:

$$K = K\_o \exp\left(-E/RT\right),\tag{2}$$

where *E* – activation energy of nitriding process, J/mole; *Ko* – preexponential multiplier, g2/сm4s; *R* – gas constant; *T* – temperature. Activation energy of nitriding determined by many researchers has different values for the certain temperature ranges and it increases with the rise of temperature of isothermal exposure as a rule. It allows to assert that process which determines the nitrogen absorption rate is changed with temperature.

The value *n* in law:

$$(\Delta m/\text{S})^n = K\pi \tag{3}$$

titanium is described by the parabolic dependence. That is the dependence of relative mass

*m/S)2 = K*

The appreciable deviations from the parabolic law caused by the presence of negligible

With the rise of temperature during isothermal exposure the intensity of interaction of titanium with nitrogen increases substantially. The parabolic constant of nitriding rate (K) is determined by the tangent of angle of inclination of straight lines of time dependence of square of mass increase. It characterizes quantitatively the relative intensity of saturation process (table 1). Values K increase practically on order with the rise of temperature on one hundred degrees (for example, at 900 оС *K* = 8,610-12 g2/сm4s and at 1000 оС *K =* 3,510-11 g2/сm4s).

quantity of oxygen impurities in nitrogen are observed at the initial stage of reaction.

Alloy <sup>К</sup> <sup>=</sup>Ко exp (Е/RT)

VТ1-0 0,11 214 ОТ4-1 0,42 229 VТ5-1 0,6 248 VТ20 0,16 241 PТ-7M 0,8 250 ОТ4 0,02 223 VТ6s 1,9 258 VТ6 0,9 267 VТ23 0,4 252 VТ32 0,4 240 VТ35 0,1 238 Table 1. Parameters of temperature dependence of parabolic constant of nitriding rate К

The temperature dependence of parabolic constant of nitriding rate is described by

where *E* – activation energy of nitriding process, J/mole; *Ko* – preexponential multiplier, g2/сm4s; *R* – gas constant; *T* – temperature. Activation energy of nitriding determined by many researchers has different values for the certain temperature ranges and it increases with the rise of temperature of isothermal exposure as a rule. It allows to assert that process

*m/S)n = K*

which determines the nitrogen absorption rate is changed with temperature.

*(* – nitriding duration; K – parabolic constant of nitriding rate.

Ко, g2/(cm4s) -Е, kJ/mole

*K = Ko exp (-E/RT),* (2)

(3)

*m = (m2 - m1)* – difference between sample's mass after and before nitriding; S –

*,* (1)

increase of the nitrided samples on time can be presented by a function:

where 

nitriding surface square;

Arrhenius equation:

The value *n* in law:

*(* change of mass increase of the nitrided samples of titanium alloys in time determined by treatment of nitriding isotherms of each investigated alloy in the double logarithmical coordinates *ln(m/S)-ln* (cotangent of angle of inclination of the received straight lines) are near to 2 that corresponds to the parabolic law of interaction (1) as well as for the unalloyed titanium.

The calculated constants (*K, Ko, E*) of the dependences *(1, 2)* at the use of corresponding mathematical models (Matychak et al., 2007, 2008, 2009, 2011) at the predetermined temperature and duration of saturation process allow to forecast the thickness of the phases formed during the nitriding process or determine time and temperature parameters of nitriding at which the thicknesses of the formed layers would be predetermined.

The results of investigation of kinetics of saturation process of titanium alloys by nitrogen testify that alloying influences on nitriding rate (fig. 1).

Fig. 1. Kinetics of nitriding of titanium alloys : a - 950 оС; b - 900 оС; c - 1000 оС; d - 1100 оС; 1 - VТ1- 0; 2 - PТ-7М; 3 - VТ5- 1; 4 - ОТ4- 1; 5 - VТ6s; 6 - VТ5; 7 - VТ1-D; 8 - VТ3- 1; 9 - VТ6; 10 - VТ20; 11 - VТ23.

With the rise of temperature of isothermal exposure the influence of alloying elements on the nitriding process increases that confirmed by the increase of distance between the

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 145

assist to enhance the surface relief. Alloying weakens the formation of surface relief owing

a b

The formation of surface relief worsens the quality of nitrided surface with the nitriding temperature rising (fig. 3). For example, nitriding of VТ23 alloy at 900°C results in change of surface roughness parameter (Ra) from 0,08 to 0,2 µm. After isothermal exposure in nitrogen at 950°C Ra is 0,4 µm, that is the surface roughness became worse on two classes. The substantial worsening of surface quality during nitriding at high temperatures complicates the obtaining of smooth surface. Therefore the use of nitriding with the purpose to increase the wear resistance of titanium foresees either limitation of process temperature ( 900 oC) or

The nitride film consists of only nitrides of base metal – -(TiN) and -(Ti2N). The grains of nitride phases have predominating orientations (table 2, 3). It is better expressed for -Ti2N grains, which are mainly oriented on planes [002]. It should be noted that the texture of phase is formed only during the process of nitride film growth (the redistribution of reflexes' intensity is not observed after short-term exposures when nitride film is thin).

Fig. 3. Surface roughness of VТ6 (а) and VТ22 (b) titanium alloys after nitriding.

to the rise of temperature of polymorphic transformation of alloys.

Fig. 2. Surface of nitrided VТ1-0 alloy: а – 850 oC, 12h; b – 950 oC, 8h.

additional surface treatment of the nitrided details.

isotherms of nitriding (fig. 1): difference in the mass increase of titanium alloys with different chemical composition at saturation by nitrogen increases.

As opposed to the unalloyed c.p.titanium the nitriding rate of titanium alloys, as a rule, is less. The mechanism allowed the alloying elements to decrease the nitrogen diffusion rate in titanium does not discussed in literature. However, as diffusion in titanium nitrides is interstitial, the influence of alloying elements effects either on the decrease of sizes of interstitial intervals in the titanium lattice or on their filling.

Thus, the process of high-temperature interaction of titanium with nitrogen is described by parabolic dependence which is the result of forming of chemical reaction products – nitrides on the metal surface that slows down the behavior in time. The presence of alloying elements in titanium does not change the process substantially and only slows its. Besides, as a result of heterogeneous reaction titanium–nitrogen the considerable dissolution of gas into the metal with formation of solid solution of nitrogen in and titanium (gasing) is observed. Therefore the study of nitriding process by only determination of general mass of absorbed nitrogen which includes both the nitride formation and gasing is incomplete and not quite correct. The differential estimation of contribution of both nitride formation and gasing during nitriding is necessary.

#### **2.2 Features of nitride formation at nitriding of titanium alloys**

The well coherent with matrix nitride film of the golden color is formed on the surface of titanium alloys during the isothermal exposure in the nitrogen at temperatures above 800 °С. The film can have different tints of base golden color (from bright to mat) which depends on temperature and duration of nitriding, chemical composition of the nitrides. The film loses the brightness at the temperature behind 900 oC. The thickness of the nitride film and the degree of saturation by nitrogen are stipulated by the time and temperature parameters of nitriding and chemical composition of the saturated material. It allows to assert that the change of colour gamut and reflection power of film depends on its thickness and degree of saturation by nitrogen because titanium nitrides, in particular TiN, is characterized by the wide homogeneity region (27…52 аt.%).

Nitride film formed at temperatures below 1000 oC repeats the contours of metallic matrix. In the case of the long exposures and temperatures below 1000 oC and above there are growths of film. On fig. 2a the characteristic topography of surface of the nitrided samples is presented: wavy inequalities forming net on surface, which, most probably, repeats the net of grains boundaries of the material matrix. These formations are most noticeable and reach the large sizes at the nitriding temperatures which are higher than temperature of polymorphic transformation (fig. 2b). phase transformation during the processes of heating and cooling causes the strain hardening, volume changes and formation of surface relief. The origin of considerable compression stresses at forming of nitride film causes the plastic deformation. It promotes the formation of quantities of inequalities. The surface topography, more or less expressed, is observed after nitriding at temperatures even lower than polymorphic transformation, and not arises after nitriding at the certain temperature. With the rise of nitriding temperature there is the growth of fragments of surface net like the growth of grain of titanium matrix. More active nitride formation on the grains boundaries promotes the forming of surface net, and processes, accompanying transformation,

isotherms of nitriding (fig. 1): difference in the mass increase of titanium alloys with

As opposed to the unalloyed c.p.titanium the nitriding rate of titanium alloys, as a rule, is less. The mechanism allowed the alloying elements to decrease the nitrogen diffusion rate in titanium does not discussed in literature. However, as diffusion in titanium nitrides is interstitial, the influence of alloying elements effects either on the decrease of sizes of

Thus, the process of high-temperature interaction of titanium with nitrogen is described by parabolic dependence which is the result of forming of chemical reaction products – nitrides on the metal surface that slows down the behavior in time. The presence of alloying elements in titanium does not change the process substantially and only slows its. Besides, as a result of heterogeneous reaction titanium–nitrogen the considerable dissolution of gas into the metal with formation of solid solution of nitrogen in and titanium (gasing) is observed. Therefore the study of nitriding process by only determination of general mass of absorbed nitrogen which includes both the nitride formation and gasing is incomplete and not quite correct. The differential estimation of contribution of both nitride formation and

The well coherent with matrix nitride film of the golden color is formed on the surface of titanium alloys during the isothermal exposure in the nitrogen at temperatures above 800 °С. The film can have different tints of base golden color (from bright to mat) which depends on temperature and duration of nitriding, chemical composition of the nitrides. The film loses the brightness at the temperature behind 900 oC. The thickness of the nitride film and the degree of saturation by nitrogen are stipulated by the time and temperature parameters of nitriding and chemical composition of the saturated material. It allows to assert that the change of colour gamut and reflection power of film depends on its thickness and degree of saturation by nitrogen because titanium nitrides, in particular TiN, is

Nitride film formed at temperatures below 1000 oC repeats the contours of metallic matrix. In the case of the long exposures and temperatures below 1000 oC and above there are growths of film. On fig. 2a the characteristic topography of surface of the nitrided samples is presented: wavy inequalities forming net on surface, which, most probably, repeats the net of grains boundaries of the material matrix. These formations are most noticeable and reach the large sizes at the nitriding temperatures which are higher than temperature of polymorphic transformation (fig. 2b). phase transformation during the processes of heating and cooling causes the strain hardening, volume changes and formation of surface relief. The origin of considerable compression stresses at forming of nitride film causes the plastic deformation. It promotes the formation of quantities of inequalities. The surface topography, more or less expressed, is observed after nitriding at temperatures even lower than polymorphic transformation, and not arises after nitriding at the certain temperature. With the rise of nitriding temperature there is the growth of fragments of surface net like the growth of grain of titanium matrix. More active nitride formation on the grains boundaries promotes the forming of surface net, and processes, accompanying transformation,

different chemical composition at saturation by nitrogen increases.

interstitial intervals in the titanium lattice or on their filling.

**2.2 Features of nitride formation at nitriding of titanium alloys** 

characterized by the wide homogeneity region (27…52 аt.%).

gasing during nitriding is necessary.

assist to enhance the surface relief. Alloying weakens the formation of surface relief owing to the rise of temperature of polymorphic transformation of alloys.

Fig. 2. Surface of nitrided VТ1-0 alloy: а – 850 oC, 12h; b – 950 oC, 8h.

The formation of surface relief worsens the quality of nitrided surface with the nitriding temperature rising (fig. 3). For example, nitriding of VТ23 alloy at 900°C results in change of surface roughness parameter (Ra) from 0,08 to 0,2 µm. After isothermal exposure in nitrogen at 950°C Ra is 0,4 µm, that is the surface roughness became worse on two classes. The substantial worsening of surface quality during nitriding at high temperatures complicates the obtaining of smooth surface. Therefore the use of nitriding with the purpose to increase the wear resistance of titanium foresees either limitation of process temperature ( 900 oC) or additional surface treatment of the nitrided details.

Fig. 3. Surface roughness of VТ6 (а) and VТ22 (b) titanium alloys after nitriding.

The nitride film consists of only nitrides of base metal – -(TiN) and -(Ti2N). The grains of nitride phases have predominating orientations (table 2, 3). It is better expressed for -Ti2N grains, which are mainly oriented on planes [002]. It should be noted that the texture of phase is formed only during the process of nitride film growth (the redistribution of reflexes' intensity is not observed after short-term exposures when nitride film is thin).

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 147

The thin outer layer of golden color consists of TiN and the thick inner layer of white color –

Due to the large brittleness of nitride layer the measuring of its quantitative characteristics is complicated. Therefore there are few articles on the kinetics of nitride formation. This explains the substantial spread and certain discordance of the results received by different researchers. The alloying elements of titanium alloys does not participate in the process of nitride formation. According to the thermodynamic activity of elements in relation to nitrogen (fig. 4), except of formation of titanium nitrides, zirconium nitrides are formed probably at nitriding of titanium alloys. The information about their formation is not found in literature, although at nitriding of alloys with 3...4 % Al, 8…12 % Zr, 1,2...2,6 %V the formation of phase (Ti, Zr)N

Alloying of titanium influences on the depth of nitride layer. In according with the recent

Fig. 4. Change of Gibbs thermodynamic potential GTo on 1 g-аt of titanium depending on temperature for the reactions of formation of some nitrides (Kiparisov & Levinskiy, 1972).

Fig. 5. Thickness of nitride layer on titanium alloys depending on duration of nitriding at

1000 оС.

results the nitrided area in (+)-alloys is less than in - and pseudo--alloys (fig. 5).

Ti2N. The perceptible growth of TiN is observed only at the long exposures.

with the lattice parameter of 0,4283 nm was fixed (Kiparisov & Levinskiy, 1972).


Table 2. The ratio І(111)/І(200) and coefficient of texture plane (200) Т(200)\* of TiNx (\*Т(200)= І(200)/ (I(200)+I(111)) (Hultman et al., 1995)).


Table 3. Relative intensity of diffraction reflexes (111) and (002) of Ti2N on the diffraction patterns from VТ6, VТ22 and Т110 alloys after nitriding.

оС, 1 h --- - оС, 5 h 1,3113 0,4327 1,0510 0,4876 оС, 10 h 1,1015 0,4758 1,0690 0,4893 оС, 1 h 1,2053 0,4535 - - оС, 5 h 1,2047 0,4537 0,9726 0,5069 оС, 10 h 0,9864 0,5034 1,0642 0,4845 оС, 1 h - - 1,105 0,4751 оС, 5 h 1,1129 0,4733 1,0209 0,4948 оС, 10 h 1,0759 0,4817 0,8165 0,5505 оС, 1 h 1,2308 0,4483 1,1349 0,4684 оС, 5 h 1,0634 0,4846 0,9272 0,5189 оС, 10 h 1,2248 0,4495 0,9865 0,5034 Table 2. The ratio І(111)/І(200) and coefficient of texture plane (200) Т(200)\* of TiNx (\*Т(200)= І(200)/

<sup>800</sup>оС, 1 h (111) <sup>254</sup> 2,7609 <sup>430</sup> 2,9861 (002) <sup>92</sup> <sup>144</sup>

<sup>800</sup>оС, 5 h (111) <sup>152</sup> 0,4199 <sup>268</sup> 0,4401 (002) <sup>362</sup> <sup>609</sup>

<sup>800</sup>оС, 10 h (111) <sup>236</sup> 0,3758 <sup>249</sup> 0,3522 (002) <sup>628</sup> <sup>707</sup>

<sup>850</sup>оС, 1 h (111) <sup>265</sup> 1,3731 <sup>345</sup> 1,1616 (002) <sup>193</sup> <sup>297</sup>

<sup>850</sup>оС, 5 h (111) <sup>251</sup> 0,5529 <sup>139</sup> 0,1154 (002) <sup>454</sup> <sup>1205</sup>

<sup>900</sup>оС, 1 h (111) <sup>316</sup> 1,5960 <sup>493</sup> 4,0410 (002) <sup>198</sup> <sup>122</sup>

<sup>900</sup>оС, 5 h (111) <sup>275</sup> 0,6643 <sup>406</sup> 0,7719 (002) <sup>414</sup> <sup>526</sup>

<sup>900</sup>оС, 10 h (111) <sup>268</sup> 0,8845 <sup>440</sup> 0,7666 (002) <sup>303</sup> <sup>574</sup>

<sup>950</sup>оС, 1 h (111) <sup>259</sup> - - - (002) - <sup>91</sup>

<sup>950</sup>оС, 10 h (111) <sup>387</sup> 3,5833 - - (002) <sup>108</sup> <sup>212</sup>

Table 3. Relative intensity of diffraction reflexes (111) and (002) of Ti2N on the diffraction

patterns from VТ6, VТ22 and Т110 alloys after nitriding.

<sup>950</sup>оС, 5 h (111) <sup>235</sup> 1,5359 <sup>216</sup> 1,1676 (002) <sup>153</sup> <sup>185</sup>

<sup>850</sup>оС, 10 h (111) <sup>121</sup> 0,1157 - - (002) <sup>1046</sup> <sup>2006</sup>

Alloys VТ6 VТ22 І(111)/І(200) Т(200) І(111)/І(200) Т(200)

> Alloys VТ6 VТ22 І, arb. units І(111)/І(200) І, arb. units І(111)/І(200)

Parameters of nitriding

(I(200)+I(111)) (Hultman et al., 1995)).

Reflexes Ti2N (hkl)

Parameters of nitriding

The thin outer layer of golden color consists of TiN and the thick inner layer of white color – Ti2N. The perceptible growth of TiN is observed only at the long exposures.

Due to the large brittleness of nitride layer the measuring of its quantitative characteristics is complicated. Therefore there are few articles on the kinetics of nitride formation. This explains the substantial spread and certain discordance of the results received by different researchers.

The alloying elements of titanium alloys does not participate in the process of nitride formation. According to the thermodynamic activity of elements in relation to nitrogen (fig. 4), except of formation of titanium nitrides, zirconium nitrides are formed probably at nitriding of titanium alloys. The information about their formation is not found in literature, although at nitriding of alloys with 3...4 % Al, 8…12 % Zr, 1,2...2,6 %V the formation of phase (Ti, Zr)N with the lattice parameter of 0,4283 nm was fixed (Kiparisov & Levinskiy, 1972).

Alloying of titanium influences on the depth of nitride layer. In according with the recent results the nitrided area in (+)-alloys is less than in - and pseudo--alloys (fig. 5).

Fig. 4. Change of Gibbs thermodynamic potential GTo on 1 g-аt of titanium depending on temperature for the reactions of formation of some nitrides (Kiparisov & Levinskiy, 1972).

Fig. 5. Thickness of nitride layer on titanium alloys depending on duration of nitriding at 1000 оС.

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 149

and -phases of titanium. For pseudo-- and (+)-alloys when the saturation process occurs at low ( 850 оС) temperatures and undurable exposures ( 5 h) the depth of gassaturated layer is decreased with the increase of coefficient of - stabilization of alloy. High temperatures and long exposures assist in the increase of depth of gas-saturated layer with

Fig. 7. Distribution of microhardness through cross section of surface layers after nitriding of - and pseudo-- (а), (+)- (b) and - (c) titanium alloys at temperatures of Т Т.

The morphology of gas-saturated layers after nitriding depends on temperature and metal's phase composition. Let's consider the influence of these factors on the morphology of the

At the temperatures of nitriding below the polymorphic transformation the morphology of gas-saturated layer does not depend on phase composition of alloys (fig. 8а). The layer consists of two parts. The first part contains -grains with high microhardness due to their

Fig. 8. Structure of gas-saturated layer of - and pseudo-- (a, b), (+)- (a, c) and - (d)

titanium alloys nitrided at the temperatures of Т Т (а) and Т Т (b-d).

*-solid solution stabilized by nitrogen*). The solubility of nitrogen in

the increase of coefficient of -stabilization.

gas-saturated layer.

strengthening by nitrogen (

The beginning of active nitride formation influences on the surface microhardness which depends on the nitrided material and ranges in 4,5…7,0 GPа. The rise of temperature leads to the increase of surface microhardness due to the activating of nitride formation (fig. 6).

Fig. 6. Dependence of surface microhardness of VT6 (a) and VT22 (b) alloys on duration and temperature of nitriding.

Thus, during the nitriding the growth of nitride film has columnar character (mainly on the grain boundaries). It assists in the rising of surface relief that worsens the surface quality at higher nitriding temperature. During the growth of nitrides the texture Ti2N is formed mainly on plane [002]. The change of thickness of nitride film with duration is described by parabolic dependence, thus the thickness of nitride film on - and (+)-titanium alloys is less than on -alloys. The texture growth increases with the rise of nitriding temperature resulting to the increase of imperfectness and heterogeneity of nitride film.

#### **2.3 Regularities of formation of gas saturated layer – Morphology of the nitrided layers**

At the high-temperature interaction of titanium alloys with nitrogen, except for formation of nitride area on the metal surface, nitrogen diffuses into the alloy, dissolves and forms the area with increased microhardness, so-called gas-saturated area. This area is identified as titanium with the increased lattice parameters (solid solution of nitrogen in -titanium). The grains of - solid solution are oriented mainly on plane [211]. The layers with the higher nitrogen concentration are characterized by the texture of -solid solution.

The process of gasing of titanium alloys is connected with the process of nitride formation. Surface microhardness (Hs ), strengthening level of gas-saturated layer (H = f (*ℓ*)) and its depth (*ℓ*) are determined by time and temperature parameters of nitriding and depend both on chemical and phase composition of alloys. The surface layer of - and pseudo--titanium alloys is the most strengthened (fig. 7). This effect weakens considerably at transition to (+)- and, especially, -alloys. Considerably bigger depth of gas-saturated layer of -alloys as compared to -alloys is determined by the values of nitrogen diffusion coefficients in -

The beginning of active nitride formation influences on the surface microhardness which depends on the nitrided material and ranges in 4,5…7,0 GPа. The rise of temperature leads to the increase of surface microhardness due to the activating of nitride formation (fig. 6).

Fig. 6. Dependence of surface microhardness of VT6 (a) and VT22 (b) alloys on duration and

Thus, during the nitriding the growth of nitride film has columnar character (mainly on the grain boundaries). It assists in the rising of surface relief that worsens the surface quality at higher nitriding temperature. During the growth of nitrides the texture Ti2N is formed mainly on plane [002]. The change of thickness of nitride film with duration is described by parabolic dependence, thus the thickness of nitride film on - and (+)-titanium alloys is less than on -alloys. The texture growth increases with the rise of nitriding temperature

resulting to the increase of imperfectness and heterogeneity of nitride film.

nitrogen concentration are characterized by the texture of -solid solution.

**2.3 Regularities of formation of gas saturated layer – Morphology of the nitrided** 

At the high-temperature interaction of titanium alloys with nitrogen, except for formation of nitride area on the metal surface, nitrogen diffuses into the alloy, dissolves and forms the area with increased microhardness, so-called gas-saturated area. This area is identified as titanium with the increased lattice parameters (solid solution of nitrogen in -titanium). The grains of - solid solution are oriented mainly on plane [211]. The layers with the higher

The process of gasing of titanium alloys is connected with the process of nitride formation.

depth (*ℓ*) are determined by time and temperature parameters of nitriding and depend both on chemical and phase composition of alloys. The surface layer of - and pseudo--titanium alloys is the most strengthened (fig. 7). This effect weakens considerably at transition to (+)- and, especially, -alloys. Considerably bigger depth of gas-saturated layer of -alloys as compared to -alloys is determined by the values of nitrogen diffusion coefficients in -

), strengthening level of gas-saturated layer (H = f (*ℓ*)) and its

temperature of nitriding.

Surface microhardness (Hs

**layers** 

and -phases of titanium. For pseudo-- and (+)-alloys when the saturation process occurs at low ( 850 оС) temperatures and undurable exposures ( 5 h) the depth of gassaturated layer is decreased with the increase of coefficient of - stabilization of alloy. High temperatures and long exposures assist in the increase of depth of gas-saturated layer with the increase of coefficient of -stabilization.

Fig. 7. Distribution of microhardness through cross section of surface layers after nitriding of - and pseudo-- (а), (+)- (b) and - (c) titanium alloys at temperatures of Т Т.

The morphology of gas-saturated layers after nitriding depends on temperature and metal's phase composition. Let's consider the influence of these factors on the morphology of the gas-saturated layer.

At the temperatures of nitriding below the polymorphic transformation the morphology of gas-saturated layer does not depend on phase composition of alloys (fig. 8а). The layer consists of two parts. The first part contains -grains with high microhardness due to their strengthening by nitrogen (*-solid solution stabilized by nitrogen*). The solubility of nitrogen in

Fig. 8. Structure of gas-saturated layer of - and pseudo-- (a, b), (+)- (a, c) and - (d) titanium alloys nitrided at the temperatures of Т Т (а) and Т Т (b-d).

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 151

The structure of gas-saturated layer of -titanium alloys does not depend on the nitriding temperature and is analogical to the structure of alloys nitrided below temperature of transformation when the structural difference between the alloy's matrix and second part of gas-saturated layer is absent (fig. 8). The first part is -phase stabilized by nitrogen with the

For -alloys the microhardness distribution of gas-saturated layer has the original regularity (fig. 8): the curve passes through a minimum on the boundary of the part of gas-saturated

For -titanium alloys the strengthening of surface layers is significantly lower than for alloys of other structural classes (microhardness distribution curves are in the region of lower

Thus, the basic characteristics of gas-saturated layer i.e depth and degree of strengthening of surface layers (surface microhardness, hardness redistribution), depend on the phase composition of nitrided material. The most strengthening of surface layers is proper for and pseudo--titanium alloys and substantially decreases at the transition to (+)- and, especially, to -alloys. The depth of gas-saturated layer of -alloys is considerably larger than depth of gas-saturated layer of -alloys. The morphology of gas-saturated layer of titanium alloys depends on nitriding temperature and phase composition of the nitrided alloy. For -alloys the morphology of gas-saturated layer does not depend on the nitriding

The strengthened surface layer consists of nitride and gas-saturated area. As it was shown above, it is the result of the high-temperature interaction of titanium with nitrogen. This interaction is accompanied with the redistribution of alloying elements in alloy's surface layers. Let's consider the general regularities of alloying elements redistribution at nitriding

The alloying elements of titanium alloys are categorized as - (Al), - (Mn, V, Mo, Cr, Fe, Nb, Si, W) - stabilizers and neutral reinforcers (Zr, Sn). During the saturation of Ti-alloys by nitrogen there is the redistribution of alloying elements between the nitrided layer and

The increase of electron concentration during the nitrogen dissolution leads to the decrease of solubility of alloying elements in titanium due to the limited solubility and also because the formation of continuous series of solid solutions. It assists in redistribution of alloying elements in the surface layers of titanium alloys: their separation from solid solution and

Thus, at thermodiffusion saturation by nitrogen there is diffusion of elements separated from solid solution into the alloy (fig. 9). The intensity of this diffusion is determined by the

The alloying elements have different solubility and diffusion coefficients in - and modifications of titanium. According to the calculations, taking into account the diffusion

temperature and thus is identical to alloys with other structures nitrided in -area.

boundaries decorated by initial -grains.

**2.4 Redistribution of alloying elements** 

matrix as well as in the gas-saturated layer.

solubility and diffusion mobility of alloying elements.

diffusion into titanium matrix.

layer stabilized by nitrogen.

hardness values) (fig. 7).

of titanium alloys.


At the nitriding temperatures higher than temperature of polymorphic transformation the gassaturated layer also consists of two parts. However the structure of every part is determined by phase composition of nitrided alloy (fig. 8, b-d). The gas-saturated layers of -, pseudo- and (+)-titanium alloys are separated by the phase boundary from the matrix which at the nitriding temperature was -titanium. During the cooling -phase is decomposed but the boundary fixed at the high temperature maintains at the room temperature as well as the structure of -phase formed in result of transformation. -phase which stabilized by nitrogen and transformed phase but saturated by nitrogen have significant differences. The morphology of gas-saturated layer of -titanium alloys does not almost depend on the nitriding temperature. That is caused by the absence of both polymorphic transformations of matrix and part of gas-saturated layer below the level of -phase stabilizing.

The first part of gas-saturated layer formed at the temperatures of ( + ) - - area is phase (grains of -solid solution of nitrogen in titanium). The high microhardness of this part is caused by large solubility of nitrogen in -titanium (21,5 аt.% in -titanium against 0,95 аt.% in -titanium at 1000 oC). It should be noted that microhardness of this part of gassaturated layer for - and pseudo--titanium alloys exceeds insignificantly the same layer for ( + )- and -alloys (18…7 GPа against 10...5 GPа). It is determined by the different nitrogen solubility in - and -phases of titanium. The second part of gas-saturated layer consists of metal transformed and enriched by nitrogen. It is separated from the first part by the phase boundary. For - and pseudo--titanium alloys this boundary is detected metallography as a dark band of high etching. For - and pseudo--titanium alloys this part is -grain of smaller size but with the increased degree of etching compared with structure of the first part. For ( + )-alloys it is mainly -phase (-plates) in -transformed structure (mixture of - and -phases). The second part of gas-saturated layer of ( + ) titanium alloys is often called "the transition area" between the gas-saturated layer and alloy's matrix because of the sharp structural difference as compared to the first part.

With the rise of nitriding temperature and duration the size of the second part of gas-saturated layer decreases, and the size of the first part increases. In addition, there is a coarsening of the structural components of both parts, and also a change of phase correlation in the direction of increasing the quantity of -phase. The typical microhardness redistribution through the gassaturated layer of - and pseudo--titanium alloys is shown in fig. 7. It represents the gradient of nitrogen concentration from the surface into the matrix. Some tendency to stabilization of microhardness in the second part of gas-saturated layer for (+)-alloy can be explained by different nitrogen solubility in - and -phases of titanium.


matrix and as a usual is less than that determined by measuring the microhardness.

matrix and part of gas-saturated layer below the level of -phase stabilizing.

At the nitriding temperatures higher than temperature of polymorphic transformation the gassaturated layer also consists of two parts. However the structure of every part is determined by phase composition of nitrided alloy (fig. 8, b-d). The gas-saturated layers of -, pseudo- and (+)-titanium alloys are separated by the phase boundary from the matrix which at the nitriding temperature was -titanium. During the cooling -phase is decomposed but the boundary fixed at the high temperature maintains at the room temperature as well as the structure of -phase formed in result of transformation. -phase which stabilized by nitrogen and transformed phase but saturated by nitrogen have significant differences. The morphology of gas-saturated layer of -titanium alloys does not almost depend on the nitriding temperature. That is caused by the absence of both polymorphic transformations of

The first part of gas-saturated layer formed at the temperatures of ( + ) - - area is phase (grains of -solid solution of nitrogen in titanium). The high microhardness of this part is caused by large solubility of nitrogen in -titanium (21,5 аt.% in -titanium against 0,95 аt.% in -titanium at 1000 oC). It should be noted that microhardness of this part of gassaturated layer for - and pseudo--titanium alloys exceeds insignificantly the same layer for ( + )- and -alloys (18…7 GPа against 10...5 GPа). It is determined by the different nitrogen solubility in - and -phases of titanium. The second part of gas-saturated layer consists of metal transformed and enriched by nitrogen. It is separated from the first part by the phase boundary. For - and pseudo--titanium alloys this boundary is detected metallography as a dark band of high etching. For - and pseudo--titanium alloys this part is -grain of smaller size but with the increased degree of etching compared with structure of the first part. For ( + )-alloys it is mainly -phase (-plates) in -transformed structure (mixture of - and -phases). The second part of gas-saturated layer of ( + ) titanium alloys is often called "the transition area" between the gas-saturated layer and

alloy's matrix because of the sharp structural difference as compared to the first part.

different nitrogen solubility in - and -phases of titanium.

With the rise of nitriding temperature and duration the size of the second part of gas-saturated layer decreases, and the size of the first part increases. In addition, there is a coarsening of the structural components of both parts, and also a change of phase correlation in the direction of increasing the quantity of -phase. The typical microhardness redistribution through the gassaturated layer of - and pseudo--titanium alloys is shown in fig. 7. It represents the gradient of nitrogen concentration from the surface into the matrix. Some tendency to stabilization of microhardness in the second part of gas-saturated layer for (+)-alloy can be explained by The structure of gas-saturated layer of -titanium alloys does not depend on the nitriding temperature and is analogical to the structure of alloys nitrided below temperature of transformation when the structural difference between the alloy's matrix and second part of gas-saturated layer is absent (fig. 8). The first part is -phase stabilized by nitrogen with the boundaries decorated by initial -grains.

For -alloys the microhardness distribution of gas-saturated layer has the original regularity (fig. 8): the curve passes through a minimum on the boundary of the part of gas-saturated layer stabilized by nitrogen.

For -titanium alloys the strengthening of surface layers is significantly lower than for alloys of other structural classes (microhardness distribution curves are in the region of lower hardness values) (fig. 7).

Thus, the basic characteristics of gas-saturated layer i.e depth and degree of strengthening of surface layers (surface microhardness, hardness redistribution), depend on the phase composition of nitrided material. The most strengthening of surface layers is proper for and pseudo--titanium alloys and substantially decreases at the transition to (+)- and, especially, to -alloys. The depth of gas-saturated layer of -alloys is considerably larger than depth of gas-saturated layer of -alloys. The morphology of gas-saturated layer of titanium alloys depends on nitriding temperature and phase composition of the nitrided alloy. For -alloys the morphology of gas-saturated layer does not depend on the nitriding temperature and thus is identical to alloys with other structures nitrided in -area.

#### **2.4 Redistribution of alloying elements**

The strengthened surface layer consists of nitride and gas-saturated area. As it was shown above, it is the result of the high-temperature interaction of titanium with nitrogen. This interaction is accompanied with the redistribution of alloying elements in alloy's surface layers. Let's consider the general regularities of alloying elements redistribution at nitriding of titanium alloys.

The alloying elements of titanium alloys are categorized as - (Al), - (Mn, V, Mo, Cr, Fe, Nb, Si, W) - stabilizers and neutral reinforcers (Zr, Sn). During the saturation of Ti-alloys by nitrogen there is the redistribution of alloying elements between the nitrided layer and matrix as well as in the gas-saturated layer.

The increase of electron concentration during the nitrogen dissolution leads to the decrease of solubility of alloying elements in titanium due to the limited solubility and also because the formation of continuous series of solid solutions. It assists in redistribution of alloying elements in the surface layers of titanium alloys: their separation from solid solution and diffusion into titanium matrix.

Thus, at thermodiffusion saturation by nitrogen there is diffusion of elements separated from solid solution into the alloy (fig. 9). The intensity of this diffusion is determined by the solubility and diffusion mobility of alloying elements.

The alloying elements have different solubility and diffusion coefficients in - and modifications of titanium. According to the calculations, taking into account the diffusion

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 153

and energy stability of system is achieved by the redistribution of other alloying elements. The similar selective redistribution, not so clear expressed, is observed for the systems, in which solubility and diffusion mobility of alloying elements differs significantly (for example, zirconium and molybdenum, tin and iron etc.). Near the interface gas (nitrogen) metal and afterwards near nitride - gas-saturated layer the areas (clusters) with the high concentration of aluminum are formed (fig. 10). It is possible that the redistribution and coagulation of aluminum will be over by establishing of short range ordering completing by decomposition of solid solution with formation of superstructure of 2-phase (Ti3Al) type.

a b Fig. 10. Image of surface layers of ОТ4-1 alloy in the characteristic rays КAl: a – 900°С, 100

With the rise of temperature of isothermal exposure the diffusion of alloying elements is activated because the diffusion coefficients increase. The active motion of alloying elements along grain boundaries leads to their loosening, nitrogen diffusion becomes accelerated and

The morphology of gas-saturated layer of titanium alloys is connected with the redistribution of alloying elements in the surface layers. There is the active diffusion of alloying elements from the most enriched by nitrogen part of layer. In that part with nitrogen concentration the redistribution is negligible: alloying elements are redistributed between - and -phases of titanium. Aluminum (-stabilizer) and, as a rule, the neutral reinforcers (zirconium, tin) are located in -phase of titanium, -stabilizers enrich -phase. Thus, the basic constituents of alloying elements redistribution at nitriding of titanium alloys are as follows: 1) separation of alloying elements from the hexagonal close-packed lattice of nitrogen solid solution in -titanium; 2) diffusion of alloying elements from the

The first process is controlled by the solubility and the second – by the diffusion constants. In result of alloying elements redistribution, as a rule: 1) a concentration of aluminum increases near the boundary nitride – gas-saturated areas; 2) the surface layers are depleted

The redistribution of alloying elements causes the structural and phase changes in the

In spite of the fact that at the increase of temperature the degree of surface strengthening increases continuously (thickness of both nitrided layer and its constituents, surface

surface layers of alloys, determining the morphology of the nitrided layer.

h; b – 950 °С, 8 h.

nitriding rate increases.

surface into the alloy's matrix.

by -stabilizing elements.

constants, the diffusion mobility of alloying elements is decreased in a sequence FeMnZrCrAlSnNbVMo. With regard to solubility, then the solubility of zirconium is unlimited, and tin and aluminum are characterized by high solubility in titanium. Vanadium, molybdenum and niobium are less soluble in -titanium but dissolve indefinitely in -titanium. Iron, chromium and manganese are limitedly solubles in -titanium and solubility in -titanium is small. Iron, manganese and chromium are redistributed the most actively in surface layers because their solubility is minimal and diffusion mobility is the highest. The solubility of zirconium, aluminum and tin in -titanium is high and diffusion mobility is low. Therefore, the substantial redistribution of these elements does not occur. Molybdenum, vanadium and niobium redistribute more active than zirconium, aluminum and tin but more weaker than iron, manganese and chromium.

Except for diffusion, there is the concentration of -stabilizing elements separated from solid solution of nitrogen in -titanium on the boundary of nitride - gas-saturated areas, even with high diffusion constants. These effects are caused by no occupied bonds between atoms located on the phases' interface and possible anomalous value of electron concentration in these areas.

Fig. 9. Scheme of redistribution of alloying elements in the gas-saturated layer of titanium alloys at nitriding (a) and images of surface layers of ОТ4-1 (b) and VT6s (c) alloys in the characteristic rays КMn and КV (1100 оС, 1 h).

Among these elements the special attention deserves aluminum. Since aluminum is stabilizer with high affinity to nitrogen (Но298 is -318,0 and -335,0 kJ/mole for AlN and TiN respectively), the increase of electron concentration of alloy during stabilization of hexagonal close-packed lattice of solid solution of nitrogen in -titanium does not influence on the solubility of aluminum and does not assist its diffusion. Releasing of lattice energy

constants, the diffusion mobility of alloying elements is decreased in a sequence FeMnZrCrAlSnNbVMo. With regard to solubility, then the solubility of zirconium is unlimited, and tin and aluminum are characterized by high solubility in titanium. Vanadium, molybdenum and niobium are less soluble in -titanium but dissolve indefinitely in -titanium. Iron, chromium and manganese are limitedly solubles in -titanium and solubility in -titanium is small. Iron, manganese and chromium are redistributed the most actively in surface layers because their solubility is minimal and diffusion mobility is the highest. The solubility of zirconium, aluminum and tin in -titanium is high and diffusion mobility is low. Therefore, the substantial redistribution of these elements does not occur. Molybdenum, vanadium and niobium redistribute more active than zirconium, aluminum

Except for diffusion, there is the concentration of -stabilizing elements separated from solid solution of nitrogen in -titanium on the boundary of nitride - gas-saturated areas, even with high diffusion constants. These effects are caused by no occupied bonds between atoms located on the phases' interface and possible anomalous value of electron concentration in

a

b c Fig. 9. Scheme of redistribution of alloying elements in the gas-saturated layer of titanium alloys at nitriding (a) and images of surface layers of ОТ4-1 (b) and VT6s (c) alloys in the

Among these elements the special attention deserves aluminum. Since aluminum is stabilizer with high affinity to nitrogen (Но298 is -318,0 and -335,0 kJ/mole for AlN and TiN respectively), the increase of electron concentration of alloy during stabilization of hexagonal close-packed lattice of solid solution of nitrogen in -titanium does not influence on the solubility of aluminum and does not assist its diffusion. Releasing of lattice energy

and tin but more weaker than iron, manganese and chromium.

characteristic rays КMn and КV (1100 оС, 1 h).

these areas.

and energy stability of system is achieved by the redistribution of other alloying elements. The similar selective redistribution, not so clear expressed, is observed for the systems, in which solubility and diffusion mobility of alloying elements differs significantly (for example, zirconium and molybdenum, tin and iron etc.). Near the interface gas (nitrogen) metal and afterwards near nitride - gas-saturated layer the areas (clusters) with the high concentration of aluminum are formed (fig. 10). It is possible that the redistribution and coagulation of aluminum will be over by establishing of short range ordering completing by decomposition of solid solution with formation of superstructure of 2-phase (Ti3Al) type.

Fig. 10. Image of surface layers of ОТ4-1 alloy in the characteristic rays КAl: a – 900°С, 100 h; b – 950 °С, 8 h.

With the rise of temperature of isothermal exposure the diffusion of alloying elements is activated because the diffusion coefficients increase. The active motion of alloying elements along grain boundaries leads to their loosening, nitrogen diffusion becomes accelerated and nitriding rate increases.

The morphology of gas-saturated layer of titanium alloys is connected with the redistribution of alloying elements in the surface layers. There is the active diffusion of alloying elements from the most enriched by nitrogen part of layer. In that part with nitrogen concentration the redistribution is negligible: alloying elements are redistributed between - and -phases of titanium. Aluminum (-stabilizer) and, as a rule, the neutral reinforcers (zirconium, tin) are located in -phase of titanium, -stabilizers enrich -phase.

Thus, the basic constituents of alloying elements redistribution at nitriding of titanium alloys are as follows: 1) separation of alloying elements from the hexagonal close-packed lattice of nitrogen solid solution in -titanium; 2) diffusion of alloying elements from the surface into the alloy's matrix.

The first process is controlled by the solubility and the second – by the diffusion constants.

In result of alloying elements redistribution, as a rule: 1) a concentration of aluminum increases near the boundary nitride – gas-saturated areas; 2) the surface layers are depleted by -stabilizing elements.

The redistribution of alloying elements causes the structural and phase changes in the surface layers of alloys, determining the morphology of the nitrided layer.

In spite of the fact that at the increase of temperature the degree of surface strengthening increases continuously (thickness of both nitrided layer and its constituents, surface

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 155

Fig. 11. Influence of thermocycling on intensity of interaction of titanium alloys with

on participating of certain alloying element in the forming of defect structure.

The efficiency of thermocycling at nitriding depends on the parameters of thermocyclic treatment (amplitude of thermal cycle, frequency of thermocycling) and increases with their

The maximal effect of thermocycling is proper for the temperature range of polymorphic transformation and correlates with the clear expressed effect of volume strengthening which increase at the rise of amplitude and frequency of thermocycling. Alloying, as a factor of intensification on interaction of titanium alloys with gas media, does not change generally the influence of thermocycling but changes only its intensity (slope of curves ) depending

The morphology of nitrided layer thermocycling as well as after the isothermal conditions is the thick nitride film ( 1 μm) and gas-saturated area. The difference from the isothermal nitriding is that the increase of degree of imperfectness of surface layers at thermocycling leads to the forming of surface nitride films with considerable deviation from stoichiometry, mainly the deficit of nonmetal component. Therefore the lattice parameter of TiN after isothermal saturation with the rise of temperature is decreased significantly while the cyclic change of temperature assists to reach the reverse dependence: lattice parameter of TiN

Taking into account the dependence of the lattice parameter of titanium mononitride on nitrogen content in the homogeneity region, the observed regularities allow to suppose that at nitriding in the conditions of thermocycling nitride with considerable deviation from stoichiometry with the deficit on nitrogen is formed on the surface. With displacement of temperature range of thermocycling into the range of lower temperatures the deviation from

It should be noted that forming of these nitride layers gives new possibilities in surface engineering of titanium alloys, in particular, at the complex modification of surface layers by the interstitial elements (Pohreliuk et al., 2007; Pohrelyuk et al., 2009, 2011; Yaskiv et al.,

nitrogen in different temperature intervals.

stoichiometry of surface nitride increases.

rising.

increases (fig. 12).

2011; Fedirko et al., 2009).

microhardness and gradient of nitrogen concentration on the cross section of surface layers increase), the use of temperature as the factor of intensification of saturation process has substantial limitations. In particular, the perceptible structural and concentration heterogeneity (on nitrogen and alloying elements) of surface layers caused by the active diffusion processes, and also the inconvertible grain growth of titanium matrix at the saturation temperatures of (+) - - areas result in the substantial decrease of fatigue life, plasticity of the nitrided details. The heightened requirements to these characteristics cause the limitation on the saturation temperature (-area) that does not always provides the provide level of surface strengthening (Н <sup>s</sup> 6…8 GPа; *ℓ* 100 μm). Moreover, with the increase of temperature the brittleness of nitrided layer increases catastrophically. In the result of thick nitride film forming the surface quality of the nitrided layer becomes worse (surface roughness increases, imperfection and heterogeneity of nitride film grow because the effect of growth texture increases). It influences negatively on the wear- and corrosion resistance of the nitrided details.

At present, it is actual to find out other factors of intensification which allow to provide the effective surface strengthening at lower nitriding temperatures and exclude the negative consequences of the influence of high nitriding temperatures on surface quality and level of mechanical characteristics.

#### **3. Nitriding of titanium alloys at thermocycling**

One of the ways to weaken the negative consequences of the high*-*temperature nitriding is to decrease the time of processing at high temperatures. It can be attained by nitriding in the conditions of thermocycling.

As opposed to the standard methods of the chemical heat treatment there are the additional sources of the influence on the structure at thermocycling. They are inherent only to the process of continuous change of temperature: phase transformations, gradient of temperature, thermal (volume) and interphase tensions caused by the difference of thermophysical characteristics of the phases. The accumulation of structural changes in the material leads to forming, moving and annihilation of point and linear defects, redistribution of distributions, forming of low-angle boundaries, migration of low-angle boundaries with absorption of defects, migration of grain boundaries between recrystallized grains with their coarsening at the simultaneous decrease of grain boundary and surface energies, by the redistribution of alloying elements etc. It results to the increase of mobility of impurity atoms and the acceleration of diffusion processes.

For titanium alloys the thermocyclic treatment is considered as a way to achieve of such structural changes which improve the level of mechanical properties. To estimate the intensity of saturation process in these conditions is impossible due to the absence of such investigations for titanium, although the interaction of steels, aluminum and nickel alloys with gases (carbon, nitrogen) at the cyclic change of temperature, pressure and gas composition is well studied.

Let's consider the regularities of interaction of titanium alloys with nitrogen in the conditions of thermocycling.

Nitriding is intensified at the cyclic change of temperature. It influences on the rise of mass increase of the samples, surface microhardness and depth of the nitrided layer as compared to the isothermal exposure at the middle temperature of thermal cycle (fig. 11).

microhardness and gradient of nitrogen concentration on the cross section of surface layers increase), the use of temperature as the factor of intensification of saturation process has substantial limitations. In particular, the perceptible structural and concentration heterogeneity (on nitrogen and alloying elements) of surface layers caused by the active diffusion processes, and also the inconvertible grain growth of titanium matrix at the saturation temperatures of (+) - - areas result in the substantial decrease of fatigue life, plasticity of the nitrided details. The heightened requirements to these characteristics cause the limitation on the saturation temperature (-area) that does not always provides the provide level of surface strengthening

<sup>s</sup> 6…8 GPа; *ℓ* 100 μm). Moreover, with the increase of temperature the brittleness of nitrided layer increases catastrophically. In the result of thick nitride film forming the surface quality of the nitrided layer becomes worse (surface roughness increases, imperfection and heterogeneity of nitride film grow because the effect of growth texture increases). It influences

At present, it is actual to find out other factors of intensification which allow to provide the effective surface strengthening at lower nitriding temperatures and exclude the negative consequences of the influence of high nitriding temperatures on surface quality and level of

One of the ways to weaken the negative consequences of the high*-*temperature nitriding is to decrease the time of processing at high temperatures. It can be attained by nitriding in the

As opposed to the standard methods of the chemical heat treatment there are the additional sources of the influence on the structure at thermocycling. They are inherent only to the process of continuous change of temperature: phase transformations, gradient of temperature, thermal (volume) and interphase tensions caused by the difference of thermophysical characteristics of the phases. The accumulation of structural changes in the material leads to forming, moving and annihilation of point and linear defects, redistribution of distributions, forming of low-angle boundaries, migration of low-angle boundaries with absorption of defects, migration of grain boundaries between recrystallized grains with their coarsening at the simultaneous decrease of grain boundary and surface energies, by the redistribution of alloying elements etc. It results to the increase of mobility

For titanium alloys the thermocyclic treatment is considered as a way to achieve of such structural changes which improve the level of mechanical properties. To estimate the intensity of saturation process in these conditions is impossible due to the absence of such investigations for titanium, although the interaction of steels, aluminum and nickel alloys with gases (carbon, nitrogen) at the cyclic change of temperature, pressure and gas

Let's consider the regularities of interaction of titanium alloys with nitrogen in the

Nitriding is intensified at the cyclic change of temperature. It influences on the rise of mass increase of the samples, surface microhardness and depth of the nitrided layer as compared

to the isothermal exposure at the middle temperature of thermal cycle (fig. 11).

negatively on the wear- and corrosion resistance of the nitrided details.

**3. Nitriding of titanium alloys at thermocycling** 

of impurity atoms and the acceleration of diffusion processes.

(Н

mechanical characteristics.

conditions of thermocycling.

composition is well studied.

conditions of thermocycling.

Fig. 11. Influence of thermocycling on intensity of interaction of titanium alloys with nitrogen in different temperature intervals.

The efficiency of thermocycling at nitriding depends on the parameters of thermocyclic treatment (amplitude of thermal cycle, frequency of thermocycling) and increases with their rising.

The maximal effect of thermocycling is proper for the temperature range of polymorphic transformation and correlates with the clear expressed effect of volume strengthening which increase at the rise of amplitude and frequency of thermocycling. Alloying, as a factor of intensification on interaction of titanium alloys with gas media, does not change generally the influence of thermocycling but changes only its intensity (slope of curves ) depending on participating of certain alloying element in the forming of defect structure.

The morphology of nitrided layer thermocycling as well as after the isothermal conditions is the thick nitride film ( 1 μm) and gas-saturated area. The difference from the isothermal nitriding is that the increase of degree of imperfectness of surface layers at thermocycling leads to the forming of surface nitride films with considerable deviation from stoichiometry, mainly the deficit of nonmetal component. Therefore the lattice parameter of TiN after isothermal saturation with the rise of temperature is decreased significantly while the cyclic change of temperature assists to reach the reverse dependence: lattice parameter of TiN increases (fig. 12).

Taking into account the dependence of the lattice parameter of titanium mononitride on nitrogen content in the homogeneity region, the observed regularities allow to suppose that at nitriding in the conditions of thermocycling nitride with considerable deviation from stoichiometry with the deficit on nitrogen is formed on the surface. With displacement of temperature range of thermocycling into the range of lower temperatures the deviation from stoichiometry of surface nitride increases.

It should be noted that forming of these nitride layers gives new possibilities in surface engineering of titanium alloys, in particular, at the complex modification of surface layers by the interstitial elements (Pohreliuk et al., 2007; Pohrelyuk et al., 2009, 2011; Yaskiv et al., 2011; Fedirko et al., 2009).

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 157

The crystallographic texture of titanium alloys depends on the chemical composition (alloying) of metal, degree of deformation, temperature and method of rolling, thickness of semi-finished rolled products, presence of gas-saturated layer etc. Heating of textured material also allows to change the texture (recrystallization, polygonization annealing, polymorphic transformation). Therefore in practice there is a great number of methods to operate the crystallographic texture

Having established the correlative dependences between crystallographic texture and processes of interaction of titanium alloys with gas media, it is possible to use the texture factor for operating of the intensity of physical and chemical processes in gas - metal system,

At the gasing of samples with t1 - "base" (0001)[10 1 0] (fraction of orientations 44 %) deformation texture (the plane of base of hexagonal close-packed lattice is parallel to the rolling plane) the rate of the increase of nitrogen concentration in titanium is higher than for samples with t2 - "prismatic" (10 1 0)[11 2 0] (fraction of orientations 50 %) texture that assists to form strengthened layers with different parameters (depth of area, surface hardness, gradient of hardness). During nitride formation with "base" texture the density of nucleation centers of nitride phases is larger and time to formation of continuous surface films is less than for the samples with "prismatic" texture (the plane of prism of hexagonal close-packed lattice is parallel to the rolling plane), that assists to form nitride films of different thickness. That is crystallographic texture of titanium alloys influences on the conditions of mass transfer on the gas - metal boundary and diffusion mobility of nitrogen. A schematically influence of crystallographic texture on the processes of interaction of titanium with nitrogen at different

Thus, the application of texture factor allows to influence on the intensity of nitride formation and gasing, changing the relation between the dimensions of nitride and gassaturated areas. Forming by preprocess of the primary crystallographic orientation ("base" or "prismatic" properly), it is possible to provide either higher level of surface strengthening

 Fig. 13. Influence of crystallographic texture on the processes of interaction of titanium with

and, consequently, to influence on improved characteristics of construction material.

Let's illustrate the influence of texture on the nitriding of titanium alloys.

phase-boundary conditions on the boundary gas - metal is presented on fig. 13.

or larger depth of nitrogen penetration in matrix.

nitrogen: а – at gasing; b - at nitride formation.

of titanium alloys allowung to form texture with set-up parameters.

Fig. 12. Lattice parameters of TiN after the isothermal nitriding and nitriding in the conditions of thermocycling. Change of lattice parameter of TiN depending on the content of nitrogen in the homogeneity region (Goldschmidt, 1967).

Thus, at thermocycling the nitriding process of titanium is intensified and reaches the maximum during processing in area of transition. The efficiency of application of thermocycling at nitriding depends on the parameters of thermocyclic treatment and rises with their increase. The temperature range of thermocycling determines the character of the surface strengthening. The nonstoichiometric nitride films with the deficit of nonmetal component are formed on the titanium surface.

The strength characteristics of titanium are improved after nitriding at thermocycling. The highest strengthening effect is observed at cycling in the area of temperatures of polymorphic transformation of titanium alloys and enhances with the increase of both amplitude and frequency of thermocycling.

#### **4. Influence of initial deformation texture on nitriding of titanium alloys**

The intensification of nitriding at thermocycling is based on the structural changes in material. The same changes in the structure of material is possible to provide before thermal heat treatment, for example, using material with deformation texture. Such approach is based on the dependence of diffusion constants on predominating crystallographic orientation (texture).

In practice the metallic materials are used, as a rule, in the polycrystalline state. Although all grains in homogeneous metal have the identical crystalline structure, however they differ in the mutual crystallographic orientation of axes. The analogue of crystallographic orientation of plane in monocrystal for polycrystal is the predominating orientation of grains (texture). One of the basic technological processes causing the formation of crystallographic texture, is plastic deformation. Formation of texture at plastic deformation occurs in the result of crystallographic planes turning in the process of sliding and twinning. In titanium the deformation occurs by sliding on the systems {10 1 0}<11 2 0>, {10 1 1}<11 2 0> and (0001)<11 2 0> (critical shear stress is minimal for plane {10 1 0} and maximal for basal plane) and by twinning on planes {10 1 2}, {11 2 2} and {11 2 } and causes corresponding deformation texture.

Fig. 12. Lattice parameters of TiN after the isothermal nitriding and nitriding in the conditions of thermocycling. Change of lattice parameter of TiN depending on the content

Thus, at thermocycling the nitriding process of titanium is intensified and reaches the maximum during processing in area of transition. The efficiency of application of thermocycling at nitriding depends on the parameters of thermocyclic treatment and rises with their increase. The temperature range of thermocycling determines the character of the surface strengthening. The nonstoichiometric nitride films with the deficit of nonmetal

The strength characteristics of titanium are improved after nitriding at thermocycling. The highest strengthening effect is observed at cycling in the area of temperatures of polymorphic transformation of titanium alloys and enhances with the increase of both

The intensification of nitriding at thermocycling is based on the structural changes in material. The same changes in the structure of material is possible to provide before thermal heat treatment, for example, using material with deformation texture. Such approach is based on the dependence of diffusion constants on predominating crystallographic orientation (texture). In practice the metallic materials are used, as a rule, in the polycrystalline state. Although all grains in homogeneous metal have the identical crystalline structure, however they differ in the mutual crystallographic orientation of axes. The analogue of crystallographic orientation of plane in monocrystal for polycrystal is the predominating orientation of grains (texture). One of the basic technological processes causing the formation of crystallographic texture, is plastic deformation. Formation of texture at plastic deformation occurs in the result of crystallographic planes turning in the process of sliding and twinning. In titanium the deformation occurs by sliding on the systems {10 1 0}<11 2 0>, {10 1 1}<11 2 0> and (0001)<11 2 0> (critical shear stress is minimal for plane {10 1 0} and maximal for basal plane) and by twinning on planes {10 1 2}, {11 2 2} and {11 2 } and causes corresponding

**4. Influence of initial deformation texture on nitriding of titanium alloys** 

of nitrogen in the homogeneity region (Goldschmidt, 1967).

component are formed on the titanium surface.

amplitude and frequency of thermocycling.

deformation texture.

The crystallographic texture of titanium alloys depends on the chemical composition (alloying) of metal, degree of deformation, temperature and method of rolling, thickness of semi-finished rolled products, presence of gas-saturated layer etc. Heating of textured material also allows to change the texture (recrystallization, polygonization annealing, polymorphic transformation). Therefore in practice there is a great number of methods to operate the crystallographic texture of titanium alloys allowung to form texture with set-up parameters.

Having established the correlative dependences between crystallographic texture and processes of interaction of titanium alloys with gas media, it is possible to use the texture factor for operating of the intensity of physical and chemical processes in gas - metal system, and, consequently, to influence on improved characteristics of construction material.

Let's illustrate the influence of texture on the nitriding of titanium alloys.

At the gasing of samples with t1 - "base" (0001)[10 1 0] (fraction of orientations 44 %) deformation texture (the plane of base of hexagonal close-packed lattice is parallel to the rolling plane) the rate of the increase of nitrogen concentration in titanium is higher than for samples with t2 - "prismatic" (10 1 0)[11 2 0] (fraction of orientations 50 %) texture that assists to form strengthened layers with different parameters (depth of area, surface hardness, gradient of hardness). During nitride formation with "base" texture the density of nucleation centers of nitride phases is larger and time to formation of continuous surface films is less than for the samples with "prismatic" texture (the plane of prism of hexagonal close-packed lattice is parallel to the rolling plane), that assists to form nitride films of different thickness. That is crystallographic texture of titanium alloys influences on the conditions of mass transfer on the gas - metal boundary and diffusion mobility of nitrogen. A schematically influence of crystallographic texture on the processes of interaction of titanium with nitrogen at different phase-boundary conditions on the boundary gas - metal is presented on fig. 13.

Thus, the application of texture factor allows to influence on the intensity of nitride formation and gasing, changing the relation between the dimensions of nitride and gassaturated areas. Forming by preprocess of the primary crystallographic orientation ("base" or "prismatic" properly), it is possible to provide either higher level of surface strengthening or larger depth of nitrogen penetration in matrix.

Fig. 13. Influence of crystallographic texture on the processes of interaction of titanium with nitrogen: а – at gasing; b - at nitride formation.

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 159

Fig. 14. Stages of nitriding of titanium alloys in nitrogen containing oxygen (a, b, c) and possibility of the intensification of process at the decrease of oxygen partial pressure.

slows down.

nitrogen to the gas - metal reaction area.

The decrease of pressure from 105 to 100 Pa (gas flow rate 0,03 l/min) with the rise of mass increase of samples causes the increase of depth of nitrided layer and the significant decrease of thickness of nitride film. With the increase of gas rarefaction to 10 Pa the nitride thickness is stabilized and the depth of nitrogen penetration into titanium is decreased. With the decrease of gas flow rate on one order of magnitude, the mass increase and depth of nitrided layer increase, and thickness of nitride film decreases. This effect is similar to the decrease of gas partial pressure. At the decrease of pressure to 0,1…1 Pa in order to intensify the nitriding it is necessary to change the nitrogen flow rate. Thus, with the decrease of nitrogen flow rate in the range of 0,03…0,003 l/min the growth of depth of nitrided layer

The observed regularities and general tendencies in the processes of saturation of titanium alloys in the rarefied dynamic nitrogen medium indicate that in the interval of rarefaction 0,1…10 Pa at the gas flow rate 0,03..0,003 l/min (specific leakage rate 710-2…710-4 Pаs-1) the kinetics of nitriding becomes receptive to the processes connected with supply of

The analysis of the results on the influence of nitrogen partial pressure and nitrogen supply rate on the mass increase of samples, surface strengthening (surface microhardness), depth of nitrided layer testifies that the providing of the indicated gas-dynamic parameters of gas medium allows the dynamic equilibrium between adsorbed and diffused nitrogen into the titanium matrix to be maintained in certain time interval. In such conditions the nitride film is not formed on the surface and the strengthened area is the solid solution of nitrogen in titanium. In due course, in the result of forming of diffusion layer and increase of nitrogen concentration on the gas – metal boundary to the necessary level for nitride formation, that

corresponds t\*, titanium nitride is fixed continually on titanium alloys (fig. 15).

The intensification of process at the use of the considered approach allows to decrease the temperature of treatment, and, consequently, to weaken the negative consequences of the influence of high saturation temperatures on the quality of surface and level of mechanical characteristics.

#### **5. Use of elements of vacuum technology at nitriding**

The analysis of results of nitriding of titanium alloys at high temperatures showed that the major reason which decelerates the diffusion of nitrogen into the matrix is the forming of thick nitride film with nitrogen diffusion coefficient less on 2 - 4 orders of magnitude than in matrix (DTiNN = 3,7610-12 сm2/s; D-TiN = 1,2910-10 сm2/s; D-TiN = 3,9210-8 сm2/s at 950 оС). Another reason is the presence of oxide films formed at technological operations of details' manufacturing and their heating to nitriding temperature due to the presence of oxygen impurities in nitrogen. Therefore, the possible ways of intensification of nitriding is to provide the corresponding conditions that allows to : 1) increase the nitrogen diffusion coefficient in nitride film or in general prevent its formation on the initial stages of nitriding; 2) favour the dissociation of existing oxide films and prevent the formation of new ones.

Let's consider some variants of realization of the above approaches.

#### **5.1 Nitriding in rarefied dynamic nitrogen atmosphere**

The parabolic character of kinetics of high-temperature interaction of titanium with nitrogen is caused by forming of nitride film on the surface. The amount of nitrogen diffused through nitride layer during its growth is decreased constantly preventing to penetration of nitrogen into the metal.

The calculated nitrogen diffusion rate in titanium and rate of nitrogen supply to the metal surface testify that under the certain conditions even all nitrogen molecules which get on surface can not be sufficient to provide the maximal flux of nitrogen atoms from surface into matrix. Except of it, not all nitrogen molecules, contacting with the surface of metal, interact with surface. In this case the processes connected with supply of nitrogen to the gas – metal reaction area become limiting. It allows to control the maximal nitrogen concentration on titanium surface and thus to provide the necessary nitrogen concentration for nitride formation. The absence of nitride film on the surface removes the diffusion barrier, and, consequently, penetration of nitrogen into titanium matrix intensifies. That is, the nitrogen partial pressure becomes the factor of intensification of nitriding process (fig. 14).

With the decrease of nitrogen partial pressure it is possible to provide the conditions when beginning of nitride film forming is shifted in time, that is at corresponding duration the nitride film is in general absent or its thickness is too thin. Thus, in the certain interval of nitrogen partial pressure the area of solid solution of nitrogen in -titanium on the surface is formed, providing the more uniform distribution of hardness in the diffusion layer and increasing the depth of nitrogen penetration.

Let's consider the general tendencies in the processes of gasing and nitride formation at nitriding of titanium alloys in the rarefied dynamic nitrogen medium.

The intensification of process at the use of the considered approach allows to decrease the temperature of treatment, and, consequently, to weaken the negative consequences of the influence of high saturation temperatures on the quality of surface and level of mechanical

The analysis of results of nitriding of titanium alloys at high temperatures showed that the major reason which decelerates the diffusion of nitrogen into the matrix is the forming of thick nitride film with nitrogen diffusion coefficient less on 2 - 4 orders of magnitude than in matrix (DTiNN = 3,7610-12 сm2/s; D-TiN = 1,2910-10 сm2/s; D-TiN = 3,9210-8 сm2/s at 950 оС). Another reason is the presence of oxide films formed at technological operations of details' manufacturing and their heating to nitriding temperature due to the presence of oxygen impurities in nitrogen. Therefore, the possible ways of intensification of nitriding is to provide the corresponding conditions that allows to : 1) increase the nitrogen diffusion coefficient in nitride film or in general prevent its formation on the initial stages of nitriding; 2) favour the dissociation of existing oxide films and prevent the

The parabolic character of kinetics of high-temperature interaction of titanium with nitrogen is caused by forming of nitride film on the surface. The amount of nitrogen diffused through nitride layer during its growth is decreased constantly preventing to penetration of nitrogen

The calculated nitrogen diffusion rate in titanium and rate of nitrogen supply to the metal surface testify that under the certain conditions even all nitrogen molecules which get on surface can not be sufficient to provide the maximal flux of nitrogen atoms from surface into matrix. Except of it, not all nitrogen molecules, contacting with the surface of metal, interact with surface. In this case the processes connected with supply of nitrogen to the gas – metal reaction area become limiting. It allows to control the maximal nitrogen concentration on titanium surface and thus to provide the necessary nitrogen concentration for nitride formation. The absence of nitride film on the surface removes the diffusion barrier, and, consequently, penetration of nitrogen into titanium matrix intensifies. That is, the nitrogen

With the decrease of nitrogen partial pressure it is possible to provide the conditions when beginning of nitride film forming is shifted in time, that is at corresponding duration the nitride film is in general absent or its thickness is too thin. Thus, in the certain interval of nitrogen partial pressure the area of solid solution of nitrogen in -titanium on the surface is formed, providing the more uniform distribution of hardness in the diffusion layer and

Let's consider the general tendencies in the processes of gasing and nitride formation at

partial pressure becomes the factor of intensification of nitriding process (fig. 14).

nitriding of titanium alloys in the rarefied dynamic nitrogen medium.

**5. Use of elements of vacuum technology at nitriding** 

Let's consider some variants of realization of the above approaches.

**5.1 Nitriding in rarefied dynamic nitrogen atmosphere** 

increasing the depth of nitrogen penetration.

characteristics.

formation of new ones.

into the metal.

Fig. 14. Stages of nitriding of titanium alloys in nitrogen containing oxygen (a, b, c) and possibility of the intensification of process at the decrease of oxygen partial pressure.

The decrease of pressure from 105 to 100 Pa (gas flow rate 0,03 l/min) with the rise of mass increase of samples causes the increase of depth of nitrided layer and the significant decrease of thickness of nitride film. With the increase of gas rarefaction to 10 Pa the nitride thickness is stabilized and the depth of nitrogen penetration into titanium is decreased. With the decrease of gas flow rate on one order of magnitude, the mass increase and depth of nitrided layer increase, and thickness of nitride film decreases. This effect is similar to the decrease of gas partial pressure. At the decrease of pressure to 0,1…1 Pa in order to intensify the nitriding it is necessary to change the nitrogen flow rate. Thus, with the decrease of nitrogen flow rate in the range of 0,03…0,003 l/min the growth of depth of nitrided layer slows down.

The observed regularities and general tendencies in the processes of saturation of titanium alloys in the rarefied dynamic nitrogen medium indicate that in the interval of rarefaction 0,1…10 Pa at the gas flow rate 0,03..0,003 l/min (specific leakage rate 710-2…710-4 Pаs-1) the kinetics of nitriding becomes receptive to the processes connected with supply of nitrogen to the gas - metal reaction area.

The analysis of the results on the influence of nitrogen partial pressure and nitrogen supply rate on the mass increase of samples, surface strengthening (surface microhardness), depth of nitrided layer testifies that the providing of the indicated gas-dynamic parameters of gas medium allows the dynamic equilibrium between adsorbed and diffused nitrogen into the titanium matrix to be maintained in certain time interval. In such conditions the nitride film is not formed on the surface and the strengthened area is the solid solution of nitrogen in titanium. In due course, in the result of forming of diffusion layer and increase of nitrogen concentration on the gas – metal boundary to the necessary level for nitride formation, that corresponds t\*, titanium nitride is fixed continually on titanium alloys (fig. 15).

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 161

the process duration. The possibility to provide the sufficient surface strengthening at the temperatures of -area excludes the negative aspects connected with high saturation temperatures, that is the quality of surface rises and the level of mechanical characteristics

Thus, with lowering of nitrogen partial pressure the nitride formation is suppressed. The absence of nitride film on the surface or substantially less its thickness weaken the diffusion barrier and penetration of nitrogen into titanium matrix intensifies. Nitriding in the rarefied dynamic nitrogen condition as compared to nitrogen of atmospheric pressure provides more uniform redistribution of hardness through the diffusion layer and more significant

The vacuum technology is widely used in practice of thermal heat treatment, including at nitriding. A brief annealing in vacuum (800 C, 2 h) before nitriding and final treatment of detail surface is recommended to conduct for distressing and prevention of warping (Samsonov & Epik, 1973). To decrease the brittleness of diffusion layer and increase the plasticity of alloys after nitriding on 10…15 % regardless of method it is recommended to conduct the additional annealing of details in vacuum at rarefaction of 410-2 Pa during 2 h

The vacuum annealing in this case is the separate technological process. However, such technological process can be used with better efficiency when applying the vacuum technology in nitriding, that is, realizing vacuum annealing not as the separate process but as the element of nitriding process. It allows to considerably shorten and simplify the treatment as in this case the additional processes of heating and cooling are not necessary

The treatment of titanium alloys in vacuum of 0,1..10 mPа when oxygen partial pressure is about 0,001…0,01 mPа excludes the possibility of formation of surface oxide films. Moreover, at the such oxygen partial pressures and corresponding time and temperature parameters it is possible to provide the conditions for dissociation and dissolution of natural oxide films before nitriding. The surface is activated and, as a result, at the subsequent inflow of nitrogen the adsorption and diffusion processes have been intensified. The upper limit of pressure of residual gases in vacuum is caused by the intensive processes of sublimation, contributing vacuum embittering of alloys's surface, and the lower one - by the

That is, providing before the nitrogen supply in the reaction furnace the pressure of the residual gases of 0,1…10 mPа provides the necessary conditions for the intensification of nitriding. At the rarefaction of 0,1…1 mPа and temperatures above 600..700 oС the dissolution of oxide films is began and the effective removal of internal tensions is occurred that determines the lower boundary of temperatures of vacuum treatment. To receive the optimal complex of the mechanical characteristics the use of temperatures above the temperature of polymorphic transformation is undesirable. This determines the upper

The use of vacuum technology elements before nitriding in above indicated temperatures interval assists in the increase of saturation of surface layers with nitrogen, depth of

(fig. 17a) and, consequently, to improve substantially its productivity.

active processes of oxidation and oxygen saturation.

boundary of temperatures range of vacuum treatment.

increases.

depth of nitrogen penetration.

**5.2 Some methods of nitriding improving** 

at 800 оС (Kiparisov & Levinskiy, 1972).

Fig. 15. Influence of temperature (T) and gas-dynamic parameters of nitrogen (р, d) on the phase-structural state of surface layers of titanium alloys.

At nitriding of titanium alloys with conservation of general tendencies the corresponding correctives in the process of saturation contributes the redistribution of alloying elements that influences on the absolute values of characteristics of the nitrided layers.

Nitriding in the rarefied nitrogen as compared to saturation in nitrogen of atmospheric pressure decreases the gradient of nitrogen concentration on the cross section of surface layers and increases the depth of penetration of nitrogen (in 1,3…2,3 time) and decreasing surface strengthening (fig. 16).

Fig. 16. Saturation temperature (а) and nitrogen partial pressure (b) as factors of the intensification of nitriding process of titanium alloys (arrows are direction of motion of phase boundaries in the areas of the identified parameters as factors of intensification).

As at lowering of nitrogen partial pressure the process of thermodiffusion saturation of titanium alloys intensifies, it makes possible to lower the nitriding temperature and decrease

Fig. 15. Influence of temperature (T) and gas-dynamic parameters of nitrogen (р, d) on the

At nitriding of titanium alloys with conservation of general tendencies the corresponding correctives in the process of saturation contributes the redistribution of alloying elements

Nitriding in the rarefied nitrogen as compared to saturation in nitrogen of atmospheric pressure decreases the gradient of nitrogen concentration on the cross section of surface layers and increases the depth of penetration of nitrogen (in 1,3…2,3 time) and decreasing

that influences on the absolute values of characteristics of the nitrided layers.

Fig. 16. Saturation temperature (а) and nitrogen partial pressure (b) as factors of the intensification of nitriding process of titanium alloys (arrows are direction of motion of phase boundaries in the areas of the identified parameters as factors of intensification).

As at lowering of nitrogen partial pressure the process of thermodiffusion saturation of titanium alloys intensifies, it makes possible to lower the nitriding temperature and decrease

phase-structural state of surface layers of titanium alloys.

surface strengthening (fig. 16).

the process duration. The possibility to provide the sufficient surface strengthening at the temperatures of -area excludes the negative aspects connected with high saturation temperatures, that is the quality of surface rises and the level of mechanical characteristics increases.

Thus, with lowering of nitrogen partial pressure the nitride formation is suppressed. The absence of nitride film on the surface or substantially less its thickness weaken the diffusion barrier and penetration of nitrogen into titanium matrix intensifies. Nitriding in the rarefied dynamic nitrogen condition as compared to nitrogen of atmospheric pressure provides more uniform redistribution of hardness through the diffusion layer and more significant depth of nitrogen penetration.

#### **5.2 Some methods of nitriding improving**

The vacuum technology is widely used in practice of thermal heat treatment, including at nitriding. A brief annealing in vacuum (800 C, 2 h) before nitriding and final treatment of detail surface is recommended to conduct for distressing and prevention of warping (Samsonov & Epik, 1973). To decrease the brittleness of diffusion layer and increase the plasticity of alloys after nitriding on 10…15 % regardless of method it is recommended to conduct the additional annealing of details in vacuum at rarefaction of 410-2 Pa during 2 h at 800 оС (Kiparisov & Levinskiy, 1972).

The vacuum annealing in this case is the separate technological process. However, such technological process can be used with better efficiency when applying the vacuum technology in nitriding, that is, realizing vacuum annealing not as the separate process but as the element of nitriding process. It allows to considerably shorten and simplify the treatment as in this case the additional processes of heating and cooling are not necessary (fig. 17a) and, consequently, to improve substantially its productivity.

The treatment of titanium alloys in vacuum of 0,1..10 mPа when oxygen partial pressure is about 0,001…0,01 mPа excludes the possibility of formation of surface oxide films. Moreover, at the such oxygen partial pressures and corresponding time and temperature parameters it is possible to provide the conditions for dissociation and dissolution of natural oxide films before nitriding. The surface is activated and, as a result, at the subsequent inflow of nitrogen the adsorption and diffusion processes have been intensified. The upper limit of pressure of residual gases in vacuum is caused by the intensive processes of sublimation, contributing vacuum embittering of alloys's surface, and the lower one - by the active processes of oxidation and oxygen saturation.

That is, providing before the nitrogen supply in the reaction furnace the pressure of the residual gases of 0,1…10 mPа provides the necessary conditions for the intensification of nitriding. At the rarefaction of 0,1…1 mPа and temperatures above 600..700 oС the dissolution of oxide films is began and the effective removal of internal tensions is occurred that determines the lower boundary of temperatures of vacuum treatment. To receive the optimal complex of the mechanical characteristics the use of temperatures above the temperature of polymorphic transformation is undesirable. This determines the upper boundary of temperatures range of vacuum treatment.

The use of vacuum technology elements before nitriding in above indicated temperatures interval assists in the increase of saturation of surface layers with nitrogen, depth of

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 163

Fig. 18. The phase-structural state of surface layers of nitrided titanium alloys: а – gassaturated area without nitride film; b, d – thin nitride film and gas-saturated area; c nitride islands and gas-saturated area; e – thick nitride film and gas-saturated area.

nitriding of titanium alloys in molecular nitrogen.

**in aqueous solutions of inorganic acids** 

Thus, the change of thermokinetic parameters of saturation, the use of vacuum technology elements and corresponding initial deformation texture allow to intensify the process of

**6. Surface engineering of titanium alloys by nitriding for corrosion protection** 

Thermodiffusion coatings, including nitride ones, protect the titanium alloys against corrosion by combining of covering and electrochemical mechanisms (Chukalovskaya et al., 1993). The covering mechanism is being realized by making of barrier layer on the metalmedium border and thus it depends on its dimension. The electrochemical mechanism is defined by electrochemical characteristics of contacting surface and thus it causes the tendency of system to disturb the balance. In other words, it leads to reactions between surface layer ions and medium. Thus, protective properties of coatings depend on their dimension and structural characteristics (such as uniformity, relief, amount of oxygen impurities). In aggressive medium every mechanism brings in own contribution in protection. The effective combination of these mechanisms is mandatory criterion to ensure the high protective properties of nitride layers. The high saturation temperature (950 oC) provides the high-quality of nitriding in commercially pure nitrogen medium. However, the increasing of nitride coating saturation by nitrogen as well as the increasing of coating dimension due to saturation temperature rising do not lead to the improving of protective properties, but quite the contrary these processes lead to the decrease of these properties due to roughness and defectiveness raising. The saturation temperature determines the changes of surface relief. The nitride film forming at the temperature lower then the

strengthened area, surface quality and more higher temperature of exposure in vacuum. With the increase of temperature the more smooth redistribution of hardness throughout surface layers is achieved.

On this stage the heating in vacuum provides the conditions which exclude the additional oxidation of titanium surface and intensifies the nitriding. The vacuum medium (~1 mPа) during the heating activates the surface in the result of dissociation of oxide films. In result, the supply of nitrogen even of technical purity (with oxygen content to 0,01 % vol.) at the saturation temperatures of 750…850 оС allows to realize the quality nitriding. Providing of low pressure of residual gases of vacuum before supply of nitrogen into the reaction furnace does not require the long-duration isothermal exposure ( 2 h). Only heating in vacuum causes the positive result because it excludes the forming of oxide film at heating.

Fig. 17. Efficiency of the use of vacuum technology elements at nitriding of titanium alloys: I – vacuum treatment; II - nitriding.

Thus, the proposed vacuum technology before nitriding (fig. 17b) assists the intensification of thermodiffusion saturation of titanium alloys with nitrogen, allows to decrease the purity requirements to nitrogen for oxygen impurities at high-temperature treatment (> 850 оС) and to realize nitriding at the temperatures of 750..850 оС.

Thus, at the use of one or another intensification factor it is possible to influence on the constituents of nitriding process – nitride formation and gasing. At thermocycling as well as at the high saturation temperatures both nitride formation and gasing intensify. At the use of vacuum technology elements (lowering of nitrogen partial pressure, heating before nitriding and exposure in vacuum) the nitride formation is slowed down but gasing is activated.

At the same time the nitrided layers formed on titanium alloys are not limited by single variant of structure (thick nitride film (1 μm) and gas-saturated area). Depending on the conditions of nitriding by the control of intensity of physical and chemical processes on the boundary gas - metal it is possible to form various phase-structural states of surface layers of the nitrided titanium (fig. 18) which allows to change the level of surface strengthening (surface hardness, depth of penetration of nitrogen, distribution of microhardness on the cross section of surface layers) in the wide range, to control by thickness, continuity, composition stoichiometry and content of oxygen impurities, and, consequently, to realize the surface engineering of titanium alloys at nitriding according to the requirements of exploitation.

strengthened area, surface quality and more higher temperature of exposure in vacuum. With the increase of temperature the more smooth redistribution of hardness throughout

On this stage the heating in vacuum provides the conditions which exclude the additional oxidation of titanium surface and intensifies the nitriding. The vacuum medium (~1 mPа) during the heating activates the surface in the result of dissociation of oxide films. In result, the supply of nitrogen even of technical purity (with oxygen content to 0,01 % vol.) at the saturation temperatures of 750…850 оС allows to realize the quality nitriding. Providing of low pressure of residual gases of vacuum before supply of nitrogen into the reaction furnace does not require the long-duration isothermal exposure ( 2 h). Only heating in vacuum

Fig. 17. Efficiency of the use of vacuum technology elements at nitriding of titanium alloys: I

Thus, the proposed vacuum technology before nitriding (fig. 17b) assists the intensification of thermodiffusion saturation of titanium alloys with nitrogen, allows to decrease the purity requirements to nitrogen for oxygen impurities at high-temperature treatment (> 850 оС)

Thus, at the use of one or another intensification factor it is possible to influence on the constituents of nitriding process – nitride formation and gasing. At thermocycling as well as at the high saturation temperatures both nitride formation and gasing intensify. At the use of vacuum technology elements (lowering of nitrogen partial pressure, heating before nitriding and exposure in vacuum) the nitride formation is slowed down but gasing is

At the same time the nitrided layers formed on titanium alloys are not limited by single variant of structure (thick nitride film (1 μm) and gas-saturated area). Depending on the conditions of nitriding by the control of intensity of physical and chemical processes on the boundary gas - metal it is possible to form various phase-structural states of surface layers of the nitrided titanium (fig. 18) which allows to change the level of surface strengthening (surface hardness, depth of penetration of nitrogen, distribution of microhardness on the cross section of surface layers) in the wide range, to control by thickness, continuity, composition stoichiometry and content of oxygen impurities, and, consequently, to realize the surface engineering of titanium alloys at nitriding according to the requirements of

causes the positive result because it excludes the forming of oxide film at heating.

surface layers is achieved.

– vacuum treatment; II - nitriding.

activated.

exploitation.

and to realize nitriding at the temperatures of 750..850 оС.

Fig. 18. The phase-structural state of surface layers of nitrided titanium alloys: а – gassaturated area without nitride film; b, d – thin nitride film and gas-saturated area; c nitride islands and gas-saturated area; e – thick nitride film and gas-saturated area.

Thus, the change of thermokinetic parameters of saturation, the use of vacuum technology elements and corresponding initial deformation texture allow to intensify the process of nitriding of titanium alloys in molecular nitrogen.

#### **6. Surface engineering of titanium alloys by nitriding for corrosion protection in aqueous solutions of inorganic acids**

Thermodiffusion coatings, including nitride ones, protect the titanium alloys against corrosion by combining of covering and electrochemical mechanisms (Chukalovskaya et al., 1993). The covering mechanism is being realized by making of barrier layer on the metalmedium border and thus it depends on its dimension. The electrochemical mechanism is defined by electrochemical characteristics of contacting surface and thus it causes the tendency of system to disturb the balance. In other words, it leads to reactions between surface layer ions and medium. Thus, protective properties of coatings depend on their dimension and structural characteristics (such as uniformity, relief, amount of oxygen impurities). In aggressive medium every mechanism brings in own contribution in protection. The effective combination of these mechanisms is mandatory criterion to ensure the high protective properties of nitride layers. The high saturation temperature (950 oC) provides the high-quality of nitriding in commercially pure nitrogen medium. However, the increasing of nitride coating saturation by nitrogen as well as the increasing of coating dimension due to saturation temperature rising do not lead to the improving of protective properties, but quite the contrary these processes lead to the decrease of these properties due to roughness and defectiveness raising. The saturation temperature determines the changes of surface relief. The nitride film forming at the temperature lower then the

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 165

the corrosion cracking of nitride films has been occurred and then the oxidation and

Change of quality, morphology and phase composition of nitride coatings influence on their

In 30 % HCl the increasing of thickness of nitride film improves the anticorrosion protection: kinetic curves of mass losses of nitrided specimens of significant thickness lie below. When thickness, caused by change of saturation temperature, increases the corrosion losses of light thin nitride films (5 m) formed at 900 oC are less in 1,2-1,4 times. Another words, the thickenings of nitride films caused by longer isothermal duration improve their protective properties, whereas when it caused by increasing of temperature of saturation – decrease. Losses of nitride films have been increased in 1,5 times (Fig. 19). Obviously, it connected with forming of surface relief. To confirm this assumption the influence of coating thickness was excluded. For that the nitride coatings of different roughness but similar thickness (10 m) have been forming. It has been achieved by nitriding at 950 oC for 10 h and at 1000 oC for 6 h. It was shown that increasing of roughness is been accompanied by increasing of

Fig. 19. Dependence of roughness (Ra, m), nitride coating thickness (h, m) and corrosion rate (К, mg/m2h) in 30% aqueous solution of chloric acid on nitriding duration at 950 oC.

Fig. 20. Dependence of corrosion rate in 30% aqueous solution of chloric acid on the

nitriding temperature and time parameters: a – 950 oC, 10 h; b – 1000 oC, 6h.

dissolution processes have being following.

corrosion rate near in 2 times (Fig. 20).

protective properties in chloric and sulphuric acids.

temperature of polymorphic transformation, only follows the material's matrix geometry. Nevertheless, at the temperatures higher than polymorphic transformation temperature, the relief fragments, such as burrs, form a grid and thus a roughness has been developed. The roughness is in 0,2…0,3 m more than after nitriding at 850 oC. The activity of nitrideforming on the grain boundaries at high temperatures assists the relief forming. Processes following transformation, such as deformation strengthening and three-dimensional changes, can only enhance the relief. It should be noted that the plastic deformation has certain influence. The plastic deformation is caused by significant residual stresses during the thick nitride film forming (Rolinski, 1988). The influence of temperature on the surface roughness is essential while the influence of isothermal duration is no significant because in this case the enhancing of surface relief is minimum. For instance, the roughness Ra after different durations (5 h and 10 h) is close: 1,06 and 1,09 m.

The investigations of influence of temperature-time parameters of nitriding on the nitride coatings dimension reveal that their thickness rises when the temperature and isothermal duration increase. For instance, after nitriding at the 950 oC thickness increases up to 3…4 m with duration change from 5 to 10 h as well as from 4 to 5 m at the duration 5 h but with the temperature increasing (from 850 oC to 950 oC) and from 5 to 6 m at duration 10 h.

At the same time low-temperature nitriding (lower than 950 oC) does not provide the forming of high quality phase composition since the surface becomes dark gold colors. XRD measurements show TiN and Ti2N reflexes as well as rutile TiO2 ones. It determines by high thermodynamic relationships between titanium and oxygen, because the active interaction each other begins at 200…300 oC while with nitrogen – at 500…600 oC. Thus the surface oxidation takes place before nitride-forming. Oxide films dissolve and dissociate only at the high temperature (upper than 850 oC).

Therefore, temperature-time and gaseous-dynamic parameters determine the dimension, quality and phase composition of nitride coating. The oxygen partial pressure and saturation temperature determine the purity of nitride coating. The roughness of nitride coating depends on saturation temperature. The thickness of nitride layer is grown with increasing of temperature and duration. Moreover the influence of temperature is more sufficient.

Since every part of nitride structure brings in own contribution in protection against corrosion, it can be possible to optimize the morphology of nitride coatings by manipulating of above-mentioned parameters to achieve the highest protective properties.

The aqueous solutions of inorganic acids dissolve the titanium and its alloys very actively (Kolotyrkin et al., 1982; Gorynin & Chechulin, 1990; Kelly, 1979). To prevent a significant corrosion losses the nitrides, carbides and borides coatings are formed, e.g. in chloric and sulphuric acids corrosion rate of nitrided titanium alloys decreases in hundred times (Kiparisov & Levinskiy, 1972; Tomashov et al., 1985; Fedirko et al., 1998). In the same time according to some studies, the corrosion mechanism has differences in chloric and sulphuric acids (Kolotyrkin et al., 1982; Brynza & Fedash, 1972; Sukhotin et al., 1990). It indicates that protective coatings must have a different structure and phase composition for use in these acids.

The corrosion of nitride coatings passes by the parabolic dependence. At first, the nitride layers are being dissolving and then the oxidation layers takes place. During the corrosion,

temperature of polymorphic transformation, only follows the material's matrix geometry. Nevertheless, at the temperatures higher than polymorphic transformation temperature, the relief fragments, such as burrs, form a grid and thus a roughness has been developed. The roughness is in 0,2…0,3 m more than after nitriding at 850 oC. The activity of nitrideforming on the grain boundaries at high temperatures assists the relief forming. Processes following transformation, such as deformation strengthening and three-dimensional changes, can only enhance the relief. It should be noted that the plastic deformation has certain influence. The plastic deformation is caused by significant residual stresses during the thick nitride film forming (Rolinski, 1988). The influence of temperature on the surface roughness is essential while the influence of isothermal duration is no significant because in this case the enhancing of surface relief is minimum. For instance, the roughness Ra after

The investigations of influence of temperature-time parameters of nitriding on the nitride coatings dimension reveal that their thickness rises when the temperature and isothermal duration increase. For instance, after nitriding at the 950 oC thickness increases up to 3…4 m with duration change from 5 to 10 h as well as from 4 to 5 m at the duration 5 h but with the temperature increasing (from 850 oC to 950 oC) and from 5 to 6 m at duration 10 h. At the same time low-temperature nitriding (lower than 950 oC) does not provide the forming of high quality phase composition since the surface becomes dark gold colors. XRD measurements show TiN and Ti2N reflexes as well as rutile TiO2 ones. It determines by high thermodynamic relationships between titanium and oxygen, because the active interaction each other begins at 200…300 oC while with nitrogen – at 500…600 oC. Thus the surface oxidation takes place before nitride-forming. Oxide films dissolve and dissociate only at the

Therefore, temperature-time and gaseous-dynamic parameters determine the dimension, quality and phase composition of nitride coating. The oxygen partial pressure and saturation temperature determine the purity of nitride coating. The roughness of nitride coating depends on saturation temperature. The thickness of nitride layer is grown with increasing of

Since every part of nitride structure brings in own contribution in protection against corrosion, it can be possible to optimize the morphology of nitride coatings by manipulating

The aqueous solutions of inorganic acids dissolve the titanium and its alloys very actively (Kolotyrkin et al., 1982; Gorynin & Chechulin, 1990; Kelly, 1979). To prevent a significant corrosion losses the nitrides, carbides and borides coatings are formed, e.g. in chloric and sulphuric acids corrosion rate of nitrided titanium alloys decreases in hundred times (Kiparisov & Levinskiy, 1972; Tomashov et al., 1985; Fedirko et al., 1998). In the same time according to some studies, the corrosion mechanism has differences in chloric and sulphuric acids (Kolotyrkin et al., 1982; Brynza & Fedash, 1972; Sukhotin et al., 1990). It indicates that protective coatings must have a different structure and phase composition for use in these

The corrosion of nitride coatings passes by the parabolic dependence. At first, the nitride layers are being dissolving and then the oxidation layers takes place. During the corrosion,

temperature and duration. Moreover the influence of temperature is more sufficient.

of above-mentioned parameters to achieve the highest protective properties.

different durations (5 h and 10 h) is close: 1,06 and 1,09 m.

high temperature (upper than 850 oC).

acids.

the corrosion cracking of nitride films has been occurred and then the oxidation and dissolution processes have being following.

Change of quality, morphology and phase composition of nitride coatings influence on their protective properties in chloric and sulphuric acids.

In 30 % HCl the increasing of thickness of nitride film improves the anticorrosion protection: kinetic curves of mass losses of nitrided specimens of significant thickness lie below. When thickness, caused by change of saturation temperature, increases the corrosion losses of light thin nitride films (5 m) formed at 900 oC are less in 1,2-1,4 times. Another words, the thickenings of nitride films caused by longer isothermal duration improve their protective properties, whereas when it caused by increasing of temperature of saturation – decrease. Losses of nitride films have been increased in 1,5 times (Fig. 19). Obviously, it connected with forming of surface relief. To confirm this assumption the influence of coating thickness was excluded. For that the nitride coatings of different roughness but similar thickness (10 m) have been forming. It has been achieved by nitriding at 950 oC for 10 h and at 1000 oC for 6 h. It was shown that increasing of roughness is been accompanied by increasing of corrosion rate near in 2 times (Fig. 20).

Fig. 19. Dependence of roughness (Ra, m), nitride coating thickness (h, m) and corrosion rate (К, mg/m2h) in 30% aqueous solution of chloric acid on nitriding duration at 950 oC.

Fig. 20. Dependence of corrosion rate in 30% aqueous solution of chloric acid on the nitriding temperature and time parameters: a – 950 oC, 10 h; b – 1000 oC, 6h.

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 167

The mass losses investigations of corrosion processes confirm the high protective properties

In 80% H2SO4 protective properties of nitride coating are changed: nitride layer without oxygen impurities is characterized by the lower corrosion resistance than oxynitride layer, e.g. more negative corrosion potential and higher current density at anodic dissolution (Fig. 22). Since the values of overstresses in both cases are similar the increasing of anodic characteristics of oxynitride coating should be related to addition modification of nitride coating by oxygen.

Fig. 22. Polarization curves of oxynitride (1) and nitride (2) coatings on the titanium surface

Indeed, the kinetic curves of mass losses of nitride coating without oxygen impurities are situated higher and corrosion rate is in 3…10 times bigger than the oxynitride layer ones. It indicates about the good protection properties of oxynitride coatings. The positive role of oxygen impurities at the providing of protective properties of nitride coatings in sulphuric acid should be explained by different activity of chlorine- and sulphate-iones in passivation

Differences in protective properties of nitride coatings in aqueous solutions of sulphuric and chloric acids indicate about necessity of differential approach to a protection against corrosion of titanium. Nitride coatings are more effective for use in chloric acid whereas

Oxygen impurities in nitride coatings increase the corrosion losses in chloric acid whereas it

Decreasing of nitride surface roughness at the simultaneous retaining of big thickness improves the protection properties in inorganic acids. More effective way to decrease the roughness is the decreasing of temperature of saturation. Increasing of thickness due to the large duration increases the protection properties but due to higher temperature –

The methods of thermodiffusion surface hardening of titanium alloys and, specifically, the procedures of thermal oxidation and nitriding have serious advantages over the other

**7. Influence of nitriding and oxidation on the wear of titanium alloys** 

of nitride coating without oxygen impurities (Table 4).

in 80% aqueous solution of sulphuric acid.

oxynitride coatings are more effective in sulphuric acid.

decreases the corrosion dissolution in sulphuric acid.

processes.

decreases ones.

Electrochemical measurements in 30 % HCl confirm above-mentioned regularities. The forms of anodic curves of nitride coating are similar to the nontreated cases. Nevertheless, the electrochemical values of nitride coating with lower roughness are better: corrosion potential (Ecor) becomes higher (+0,08 V versus +0,06 V) and corrosion current density (icor) becomes lower from 1,0 up to 0,2 A/m2 at the active dissolution and from 3,0 up to 0,8 A/m2 at the passive state (fig. 21). Decreasing of the growth of corrosion current density on the different intervals of anodic curve is obviously connected with the forming of oxide films of different composition during the polarization (Gorbachov, 1983).

Fig. 21. Polarization curves of nitrided titanium in 30% aqueous solution of chloric acid (1 - 950 оС, 5 h; 2 - 850 оС, 5 h).

In 30% HCl the oxygen impurities or oxides including in nitride layers after the nitriding at 850 oC increase the corrosion rate (Table 4). For instance, the corrosion potential of free of oxygen nitride coating is more positive than oxynitride coating ones (Fig. 21). The corrosion current density decreases up to 10-1 A/m2. The decreasing of anodic current density on the nitride coating in comparison with oxynitride ones indicates about a big braking of dissolution processes and confirms their advantage in protective properties. Cathode polarization passes by the hydrogen polarization mechanism. At the cathode polarization the nitride coating has lower current densities of cathode hydrogen depolarization. It indicates about the increasing of protection against electron conduction at the cathode depolarization of hydrogen. It is obviously that the hydrogen pickup of nitride coating is decrease sufficiently at that it decreases the hydrogen degradation of ones.


Table 4. Corrosion rate of nitride coatings of different morphology in aqueous solutions of inorganic acids

Electrochemical measurements in 30 % HCl confirm above-mentioned regularities. The forms of anodic curves of nitride coating are similar to the nontreated cases. Nevertheless, the electrochemical values of nitride coating with lower roughness are better: corrosion potential (Ecor) becomes higher (+0,08 V versus +0,06 V) and corrosion current density (icor) becomes lower from 1,0 up to 0,2 A/m2 at the active dissolution and from 3,0 up to 0,8 A/m2 at the passive state (fig. 21). Decreasing of the growth of corrosion current density on the different intervals of anodic curve is obviously connected with the forming of oxide

Fig. 21. Polarization curves of nitrided titanium in 30% aqueous solution of chloric acid (1 -

In 30% HCl the oxygen impurities or oxides including in nitride layers after the nitriding at 850 oC increase the corrosion rate (Table 4). For instance, the corrosion potential of free of oxygen nitride coating is more positive than oxynitride coating ones (Fig. 21). The corrosion current density decreases up to 10-1 A/m2. The decreasing of anodic current density on the nitride coating in comparison with oxynitride ones indicates about a big braking of dissolution processes and confirms their advantage in protective properties. Cathode polarization passes by the hydrogen polarization mechanism. At the cathode polarization the nitride coating has lower current densities of cathode hydrogen depolarization. It indicates about the increasing of protection against electron conduction at the cathode depolarization of hydrogen. It is obviously that the hydrogen pickup of nitride coating is

Nitride coating with large surface

Multiphase coating (mixture of

Table 4. Corrosion rate of nitride coatings of different morphology in aqueous solutions of

relief 5,2 12,0

nitrides and oxides or oxynitrides) 9,7 7,0

Corrosion rate, *К, mg/m2h* 30% HCl 80% H2SO4

decrease sufficiently at that it decreases the hydrogen degradation of ones.

Coating technique Morphology of coating

films of different composition during the polarization (Gorbachov, 1983).

950 оС, 5 h; 2 - 850 оС, 5 h).

Heating and cooling in nitrogen, 950 оС, 5 h

Heating and cooling in nitrogen, 850 оС, 5 h

inorganic acids

The mass losses investigations of corrosion processes confirm the high protective properties of nitride coating without oxygen impurities (Table 4).

In 80% H2SO4 protective properties of nitride coating are changed: nitride layer without oxygen impurities is characterized by the lower corrosion resistance than oxynitride layer, e.g. more negative corrosion potential and higher current density at anodic dissolution (Fig. 22). Since the values of overstresses in both cases are similar the increasing of anodic characteristics of oxynitride coating should be related to addition modification of nitride coating by oxygen.

Fig. 22. Polarization curves of oxynitride (1) and nitride (2) coatings on the titanium surface in 80% aqueous solution of sulphuric acid.

Indeed, the kinetic curves of mass losses of nitride coating without oxygen impurities are situated higher and corrosion rate is in 3…10 times bigger than the oxynitride layer ones. It indicates about the good protection properties of oxynitride coatings. The positive role of oxygen impurities at the providing of protective properties of nitride coatings in sulphuric acid should be explained by different activity of chlorine- and sulphate-iones in passivation processes.

Differences in protective properties of nitride coatings in aqueous solutions of sulphuric and chloric acids indicate about necessity of differential approach to a protection against corrosion of titanium. Nitride coatings are more effective for use in chloric acid whereas oxynitride coatings are more effective in sulphuric acid.

Oxygen impurities in nitride coatings increase the corrosion losses in chloric acid whereas it decreases the corrosion dissolution in sulphuric acid.

Decreasing of nitride surface roughness at the simultaneous retaining of big thickness improves the protection properties in inorganic acids. More effective way to decrease the roughness is the decreasing of temperature of saturation. Increasing of thickness due to the large duration increases the protection properties but due to higher temperature – decreases ones.

#### **7. Influence of nitriding and oxidation on the wear of titanium alloys**

The methods of thermodiffusion surface hardening of titanium alloys and, specifically, the procedures of thermal oxidation and nitriding have serious advantages over the other

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 169

after oxidation in a vacuum. At the same time, its fatigue life decreases by a factor of 3-6

= 1,0 % = 1,5 %

No Treatment Alloy σВ, МПа δ, % Fatigue life, cycles

1 No treatment VT1-0 450 34,5 2500 900

2 Oxidation VT1-0 374 21,7 1000 300

3 Nitriding VT1-0 395 9,3 300 50

Thus, the energy consumption, productivity, and the general level of the attained mechanical characteristics of workpieces for the procedure of oxidation are better than for nitriding. At the same time, the process of nitriding enables to set the higher degrees of surface hardening than oxidation due to the difference between the ionic radii of the interstitial elements (0,148 nm for nitrogen and 0,136 nm for oxygen). Moreover, titanium nitrides are characterized by much smaller Pilling-Bedworth ratios (1,1 for TiN and 1,7 for TiO2), coefficients of thermal expansion and residual stresses in the surface layer of nitrides, which excludes the possibility of violation of the continuity of the formed film (Strafford, 1979). In addition, the corrosion resistance of nitrided layer is higher than the corrosion resistance of the oxidized layer. Thus, the intensification of nitriding makes it possible to decrease the temperature and time of treatment and preserve the characteristics of toughness of the base material. This enables us to recommend this type of treatment as an

The diffusion coefficient of nitrogen in titanium nitride is in 2-4 times lower than in titanium. Indeed, at 850 °C we have DNTiN = 3,95·10-13 cm2×s-1, DNα-Ti = 1,81·10-11 cm2×s-1, and

As the nitrogen partial pressure decreases due to the weakening of the barrier effect of the nitride film or its complete vanishing, the process of nitriding intensifies. However, a vacuum of 102 Pa is insufficient for a positive effect attained only if the pressure of the dynamic atmosphere of nitrogen is as low as 0,1-1 Pa for a feed rate of the gas of 0,003 l/min (Fedirko & Pohrelyuk, 1995). The elements of vacuum technology (heating in a vacuum and preliminary vacuum annealing) also intensify the thermodynamic saturation of titanium with nitrogen. The low pressure of residual gases of rarefied atmosphere (~ 10 mPa) prior to the delivery of nitrogen in the stage of heating and, for a short period of time (≤ 2 h) at the saturation temperature, activates the titanium surface and promotes the dissociation of oxide films.

OT4-1 790 19,3 10100 2200 BT6s 1040 15,3 33700 6500

OT4-1 729 5,9 2600 250 BT6s 975 7,2 11700 2000

OT4-1 735 5,1 2700 230 BT6s 950 4,8 4200 1500

VT1-0 369 32,9 4000 770 OT4-1 740 18,4 3700 670 VT6s 943 14,2 13200 3100

depending on the level of strains (Table 5, rows 2 and 3).

Table 5. Mechanical characteristics of titanium alloys.

efficient and cost-effective method of surface hardening.

4 Nitriding in

DNβ-Ti = 9,03·10-9 cm2×s-1.

rarefied nitrogen

available methods under service conditions excluding the application of high contact stresses (>4 MPa) (Gorynin & Chechulin, 1990; Nazarenko et al., 1998). They are technologically simple, guarantee the reliable physical and chemical characteristics of treated surfaces, and do not require any additional technological procedures. Unlike the application of coatings, in this case, we have no problems connected with the adhesion between the hardened layers and the matrix, porosity, and, hence, with the sensitivity to aggressive media.

Let's analyze the wear resistance of titanium after thermal oxidation and nitriding.

To increase the output and efficiency of the technological processes of surface hardening, it is necessary to find out the ways of their intensification because the procedure of thermodiffusion saturation of titanium alloys, e.g., with interstitial impurities (oxygen, nitrogen, carbon, and boron), requires high temperatures and long duration. At present, the problem of intensification of the thermodiffusion saturation of titanium alloys with oxygen finds its solution in new technologies of thermal oxidation. Both the duration of process of getting the desired thickness and degree of hardening of the diffusion layer and its temperature can be decreased by applying of the procedures of boiling-bed and vacuum oxidation.

The thermodiffusion saturation of titanium alloys with oxygen from the boiling bed (650-800 °C, 4-7 h) is accelerated due to the activation of the surface of workpieces in contact (friction) with sand particles, which intensifies the processes of absorption and adsorption (Zavarov et al., 1985). In this case, we observe the formation of a hard (with a surface microhardness of 6.0-8.5GPa depending on the alloy) wear-resistant diffusion layer consisting of a 3-7-µmthick film of titanium dioxide (in the rutile modification) and an interstitial solid solution of oxygen in titanium with a thickness of 20-70 µm.

However, the non-uniform boiling of sand, its insufficient degree of dispersion and deviations from the required temperature conditions quite often lead to the formation of cavities on the titanium surface and exfoliation of the surface layer, which means that the corresponding workpieces must be rejected.

More stable results are attained in the process of thermal oxidation (700-1050 °C, 0.3-7 h) of titanium alloys in a vacuum (~ 0,1 Pa). In this case, the surface of workpieces is saturated with residual gases of vacuum media (oxygen, nitrogen, and carbon), which results in the formation of a hard wear-resistant layer consisting of two areas: a complex compound of titanium oxides, nitrides, and carbides and an interstitial solid solution of these elements in the matrix.

The procedure of nitriding of titanium alloys is carried out at temperatures of about 950 °C for 15-30 h either under the atmospheric pressure (105 Pa) or in rarefied nitrogen (≤ 102 Pa) (Gorynin & Chechulin, 1990). Long periods of holding at high temperatures result in the irreversible growth of grains in the titanium matrix accompanied by the formation of brittle surface layers and, hence, in a pronounced deterioration of the mechanical characteristics of nitrided workpieces. The characteristics of plasticity and fatigue life of the material prove to be especially sensitive to high-temperature treatment (Table 5). Thus, after nitriding in the indicated mode, the plasticity of unalloyed VT1-0 titanium becomes in 2.3 times lower than

available methods under service conditions excluding the application of high contact stresses (>4 MPa) (Gorynin & Chechulin, 1990; Nazarenko et al., 1998). They are technologically simple, guarantee the reliable physical and chemical characteristics of treated surfaces, and do not require any additional technological procedures. Unlike the application of coatings, in this case, we have no problems connected with the adhesion between the hardened layers and the matrix, porosity, and, hence, with the sensitivity to

To increase the output and efficiency of the technological processes of surface hardening, it is necessary to find out the ways of their intensification because the procedure of thermodiffusion saturation of titanium alloys, e.g., with interstitial impurities (oxygen, nitrogen, carbon, and boron), requires high temperatures and long duration. At present, the problem of intensification of the thermodiffusion saturation of titanium alloys with oxygen finds its solution in new technologies of thermal oxidation. Both the duration of process of getting the desired thickness and degree of hardening of the diffusion layer and its temperature can be decreased by applying of the procedures of boiling-bed and vacuum

The thermodiffusion saturation of titanium alloys with oxygen from the boiling bed (650-800 °C, 4-7 h) is accelerated due to the activation of the surface of workpieces in contact (friction) with sand particles, which intensifies the processes of absorption and adsorption (Zavarov et al., 1985). In this case, we observe the formation of a hard (with a surface microhardness of 6.0-8.5GPa depending on the alloy) wear-resistant diffusion layer consisting of a 3-7-µmthick film of titanium dioxide (in the rutile modification) and an interstitial solid solution of

However, the non-uniform boiling of sand, its insufficient degree of dispersion and deviations from the required temperature conditions quite often lead to the formation of cavities on the titanium surface and exfoliation of the surface layer, which means that the

More stable results are attained in the process of thermal oxidation (700-1050 °C, 0.3-7 h) of titanium alloys in a vacuum (~ 0,1 Pa). In this case, the surface of workpieces is saturated with residual gases of vacuum media (oxygen, nitrogen, and carbon), which results in the formation of a hard wear-resistant layer consisting of two areas: a complex compound of titanium oxides, nitrides, and carbides and an interstitial solid solution of these elements in

The procedure of nitriding of titanium alloys is carried out at temperatures of about 950 °C for 15-30 h either under the atmospheric pressure (105 Pa) or in rarefied nitrogen (≤ 102 Pa) (Gorynin & Chechulin, 1990). Long periods of holding at high temperatures result in the irreversible growth of grains in the titanium matrix accompanied by the formation of brittle surface layers and, hence, in a pronounced deterioration of the mechanical characteristics of nitrided workpieces. The characteristics of plasticity and fatigue life of the material prove to be especially sensitive to high-temperature treatment (Table 5). Thus, after nitriding in the indicated mode, the plasticity of unalloyed VT1-0 titanium becomes in 2.3 times lower than

Let's analyze the wear resistance of titanium after thermal oxidation and nitriding.

aggressive media.

oxidation.

the matrix.

oxygen in titanium with a thickness of 20-70 µm.

corresponding workpieces must be rejected.


after oxidation in a vacuum. At the same time, its fatigue life decreases by a factor of 3-6 depending on the level of strains (Table 5, rows 2 and 3).

Table 5. Mechanical characteristics of titanium alloys.

Thus, the energy consumption, productivity, and the general level of the attained mechanical characteristics of workpieces for the procedure of oxidation are better than for nitriding. At the same time, the process of nitriding enables to set the higher degrees of surface hardening than oxidation due to the difference between the ionic radii of the interstitial elements (0,148 nm for nitrogen and 0,136 nm for oxygen). Moreover, titanium nitrides are characterized by much smaller Pilling-Bedworth ratios (1,1 for TiN and 1,7 for TiO2), coefficients of thermal expansion and residual stresses in the surface layer of nitrides, which excludes the possibility of violation of the continuity of the formed film (Strafford, 1979). In addition, the corrosion resistance of nitrided layer is higher than the corrosion resistance of the oxidized layer. Thus, the intensification of nitriding makes it possible to decrease the temperature and time of treatment and preserve the characteristics of toughness of the base material. This enables us to recommend this type of treatment as an efficient and cost-effective method of surface hardening.

The diffusion coefficient of nitrogen in titanium nitride is in 2-4 times lower than in titanium. Indeed, at 850 °C we have DNTiN = 3,95·10-13 cm2×s-1, DNα-Ti = 1,81·10-11 cm2×s-1, and DNβ-Ti = 9,03·10-9 cm2×s-1.

As the nitrogen partial pressure decreases due to the weakening of the barrier effect of the nitride film or its complete vanishing, the process of nitriding intensifies. However, a vacuum of 102 Pa is insufficient for a positive effect attained only if the pressure of the dynamic atmosphere of nitrogen is as low as 0,1-1 Pa for a feed rate of the gas of 0,003 l/min (Fedirko & Pohrelyuk, 1995). The elements of vacuum technology (heating in a vacuum and preliminary vacuum annealing) also intensify the thermodynamic saturation of titanium with nitrogen.

The low pressure of residual gases of rarefied atmosphere (~ 10 mPa) prior to the delivery of nitrogen in the stage of heating and, for a short period of time (≤ 2 h) at the saturation temperature, activates the titanium surface and promotes the dissociation of oxide films.

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 171

The process nitriding of titanium alloys is accompanied by the formation of a typical surface pattern, which worsens the quality of the surfaces of the workpieces (the parameter *Ra* increases by 0,4-0,7 mm). During the friction, the microproasperities (elements of the pattern) are separated, move into the contact zone, and finally, penetrate into the relatively soft counterbody. These hard particles of nitrides play the role of an abrasive substance and make scratches of relatively small depth in the nitrided surface. Thus, to enhance the characteristics of wear resistance, it is necessary to improve the quality of the nitrided surface, which, in turn, depends on the temperature of nitriding (the lower is the saturation

The nitriding noticeably increases the antifriction characteristics of titanium in a friction couple with bronze. The friction coefficient equal to 0,18 is stable and independent of contact

The process of friction of the oxidized disk is accompanied both by the wear of its surface layers and the process of transfer of small pieces of soft bronze to the oxidized surface. This is connected with the adhesion of bronze particles to the disk leading to the formation of unstable secondary structures, which are destroyed and removed in the course of friction, and the process is repeated again. As a result, the time dependence of the mass of the oxidized disk is not monotonic (Fig. 23, curves 3 and 4). Under a load of 1 MPa, the processes of mass transfer of bronze and its fracture are practically balanced. As a result of fitting of the mating surfaces, the friction coefficient decreases from 0,23 to 0,14 and then stabilizes. Under a load of 2 MPa, the mass transfer of bronze is predominant and, in the course of time, the oxidized surface is more and more intensely rubbed with bronze. The friction coefficient is unstable and varies from 0,24 in the stage of fitting to 0,22-0,21 in the stationary mode. In aggressive media, the appearance of bronze on the oxidized titanium surface leads to the formation of galvanic couples and, thus, promotes the corrosion

The mass increment of the oxidized disk exceeds the degree of wear of the nitrided surface by one or even two orders of magnitude and the wear of the counterbodies is of the same order. Moreover, the wear of the counterbody is more intense in couples with the nitrided

In the friction couple of nitrided titanium with bronze, the wear of bronze exceeds the wear of the nitrided disk by more than three orders of magnitude. Under a load of 1 MPa, the mass loss of the shoe is a monotonically increasing function of time (Fig. 24, curve 1). Under a load of 2 MPa, this process becomes in 1,4-1,6 times more intense (curve 2). For a friction couple of oxidized titanium with bronze, the increase in the load leads to a more pronounced increase in the wear of the shoe (by a factor of 3,4-3,8) (Fig. 24, curves 3 and 4). Moreover, under a contact pressure of 1 MPa, the mass losses of the shoe in the process of friction with the nitrided disk are in 1,9 times greater than for the oxidized disk but, under a pressure of 2 MPa, the situation changes and the mass losses of the shoe coupled with the

The mass losses of the friction couples and their elements after different parts of the basic friction path have their own regularities. Let's now consider the wear of the counterbody in

and oxidized disks and, thus, determines the wear of the friction couples.

oxidized disk are in 1,6 times greater than for the nitrided disk.

the process of friction against the nitrided and oxidized disks.

temperature, the lower the roughness of the surface) (Fedirko & Pohrelyuk, 1995).

pressure.

processes.

These technological procedures decrease temperature and time of treatment down to 800- 900 °C and 5-10 h and, hence, significantly improve the mechanical characteristics of nitrided samples (Table 5, rows 2 and 4). Thus, the plasticity and fatigue life of nitrided VT1- 0 (c.p. titanium) are, respectively, in 1,5 and 2,6-4,0 (depending on the level of strains) times higher than in the case of oxidation.

The outlined procedure of nitriding leads to the formation of a hardened layer consisting of a TiN + Ti2N nitride film (≤ 1 µm) and a deep diffusion layer (100-180 µm).

Let's compare the wear resistance of the oxidized (in a vacuum of 0,1 Pa at 850 °C for 5 h) and nitrided [heating to 850 °C in a vacuum of 10 mPa, creation and maintenance of a dynamic atmosphere of nitrogen (1 Pa, 0.003 l/min) for 5 h, and cooling in nitrogen] titanium. The surface microhardness of nitrided and oxidized layers was 7,9 and 5,1 GPa and the thickness of hardened layer was 100 and 55 µm, respectively.

Wear tests were carried out for the case of boundary sliding friction with lubrication with a AMG-10 hydraulic fluid in an SMTs-2 friction-testing machine by using the disk-shoe mating scheme. The contact load was as large as 1 and 2 MPa. The counter body was made of bronze. The sliding velocity was equal to 0,6 m/s. The friction path was equal to 10 km. To make the contact area of the mating bodies not lower than 90%, the friction couple was run in on a path of 200 m. We analyzed the mass losses of the treated specimens and counter bodies, the friction coefficients, and wear depending on the lengths of the basic friction paths (1, 2, 2, 2, and 3 km).

It was discovered that the wear resistance of nitrided titanium is quite high. After testing, the mass losses of the disk for the indicated contact loads did not exceed 1 mg (Fig. 23, curves 1 and 2). The influence of the load is noticeable (i.e., for 2 MPa, the mass losses are higher) only for the first 1.5 h of operation of the friction couple (this corresponds to a friction path of about 3 km). After this, the mass losses under loads of 1 and 2 MPa are practically equal.

Fig. 23. Kinetics of mass changes of nitrided (1, 2) and oxidized (3, 4) disks of VT1-0 titanium in the process of friction with BrAZh9-4l bronze: 1, 3 – 1 MPa; 2, 4 – 2 MPa.

These technological procedures decrease temperature and time of treatment down to 800- 900 °C and 5-10 h and, hence, significantly improve the mechanical characteristics of nitrided samples (Table 5, rows 2 and 4). Thus, the plasticity and fatigue life of nitrided VT1- 0 (c.p. titanium) are, respectively, in 1,5 and 2,6-4,0 (depending on the level of strains) times

The outlined procedure of nitriding leads to the formation of a hardened layer consisting of

Let's compare the wear resistance of the oxidized (in a vacuum of 0,1 Pa at 850 °C for 5 h) and nitrided [heating to 850 °C in a vacuum of 10 mPa, creation and maintenance of a dynamic atmosphere of nitrogen (1 Pa, 0.003 l/min) for 5 h, and cooling in nitrogen] titanium. The surface microhardness of nitrided and oxidized layers was 7,9 and 5,1 GPa

Wear tests were carried out for the case of boundary sliding friction with lubrication with a AMG-10 hydraulic fluid in an SMTs-2 friction-testing machine by using the disk-shoe mating scheme. The contact load was as large as 1 and 2 MPa. The counter body was made of bronze. The sliding velocity was equal to 0,6 m/s. The friction path was equal to 10 km. To make the contact area of the mating bodies not lower than 90%, the friction couple was run in on a path of 200 m. We analyzed the mass losses of the treated specimens and counter bodies, the friction coefficients, and wear depending on the lengths of the basic friction

It was discovered that the wear resistance of nitrided titanium is quite high. After testing, the mass losses of the disk for the indicated contact loads did not exceed 1 mg (Fig. 23, curves 1 and 2). The influence of the load is noticeable (i.e., for 2 MPa, the mass losses are higher) only for the first 1.5 h of operation of the friction couple (this corresponds to a friction path of about

Fig. 23. Kinetics of mass changes of nitrided (1, 2) and oxidized (3, 4) disks of VT1-0 titanium

in the process of friction with BrAZh9-4l bronze: 1, 3 – 1 MPa; 2, 4 – 2 MPa.

3 km). After this, the mass losses under loads of 1 and 2 MPa are practically equal.

a TiN + Ti2N nitride film (≤ 1 µm) and a deep diffusion layer (100-180 µm).

and the thickness of hardened layer was 100 and 55 µm, respectively.

higher than in the case of oxidation.

paths (1, 2, 2, 2, and 3 km).

The process nitriding of titanium alloys is accompanied by the formation of a typical surface pattern, which worsens the quality of the surfaces of the workpieces (the parameter *Ra* increases by 0,4-0,7 mm). During the friction, the microproasperities (elements of the pattern) are separated, move into the contact zone, and finally, penetrate into the relatively soft counterbody. These hard particles of nitrides play the role of an abrasive substance and make scratches of relatively small depth in the nitrided surface. Thus, to enhance the characteristics of wear resistance, it is necessary to improve the quality of the nitrided surface, which, in turn, depends on the temperature of nitriding (the lower is the saturation temperature, the lower the roughness of the surface) (Fedirko & Pohrelyuk, 1995).

The nitriding noticeably increases the antifriction characteristics of titanium in a friction couple with bronze. The friction coefficient equal to 0,18 is stable and independent of contact pressure.

The process of friction of the oxidized disk is accompanied both by the wear of its surface layers and the process of transfer of small pieces of soft bronze to the oxidized surface. This is connected with the adhesion of bronze particles to the disk leading to the formation of unstable secondary structures, which are destroyed and removed in the course of friction, and the process is repeated again. As a result, the time dependence of the mass of the oxidized disk is not monotonic (Fig. 23, curves 3 and 4). Under a load of 1 MPa, the processes of mass transfer of bronze and its fracture are practically balanced. As a result of fitting of the mating surfaces, the friction coefficient decreases from 0,23 to 0,14 and then stabilizes. Under a load of 2 MPa, the mass transfer of bronze is predominant and, in the course of time, the oxidized surface is more and more intensely rubbed with bronze. The friction coefficient is unstable and varies from 0,24 in the stage of fitting to 0,22-0,21 in the stationary mode. In aggressive media, the appearance of bronze on the oxidized titanium surface leads to the formation of galvanic couples and, thus, promotes the corrosion processes.

The mass increment of the oxidized disk exceeds the degree of wear of the nitrided surface by one or even two orders of magnitude and the wear of the counterbodies is of the same order. Moreover, the wear of the counterbody is more intense in couples with the nitrided and oxidized disks and, thus, determines the wear of the friction couples.

In the friction couple of nitrided titanium with bronze, the wear of bronze exceeds the wear of the nitrided disk by more than three orders of magnitude. Under a load of 1 MPa, the mass loss of the shoe is a monotonically increasing function of time (Fig. 24, curve 1). Under a load of 2 MPa, this process becomes in 1,4-1,6 times more intense (curve 2). For a friction couple of oxidized titanium with bronze, the increase in the load leads to a more pronounced increase in the wear of the shoe (by a factor of 3,4-3,8) (Fig. 24, curves 3 and 4). Moreover, under a contact pressure of 1 MPa, the mass losses of the shoe in the process of friction with the nitrided disk are in 1,9 times greater than for the oxidized disk but, under a pressure of 2 MPa, the situation changes and the mass losses of the shoe coupled with the oxidized disk are in 1,6 times greater than for the nitrided disk.

The mass losses of the friction couples and their elements after different parts of the basic friction path have their own regularities. Let's now consider the wear of the counterbody in the process of friction against the nitrided and oxidized disks.

Chemico-Thermal Treatment of Titanium Alloys – Nitriding 173

Chukalovskaya Т.V., Myedova I.L., Tomashov N.D. at al. (1993). Corrosion properties and

Fedirko V. M. & Pohrelyuk I. M. (1995). Nitriding of Titanium and Its Alloys [in Ukrainian],

Fedirko V.M, Pohrelyuk I.M. & Yaskiv O.I. (1998). Corrosion resistance of nitrided titanium

Fedirko V.M., Pohrelyuk I.M. & Yaskiv O.I. (2009). Thermodiffusion multicomponent saturation of titanium alloys [in Ukrainian], Naukova Dumka, Kiev Goldfain V. I., Zuev A. M., Kablukov A. G. et al. (1977). Influence of friction and

Gorbachov A. K. (1983). Thermodynamics of reductive-oxidative processes in TiN - H2O

Gorynin I.V. & Chechulin B.B. (1990). Titanium in mechanical engineering [in Russian],

Hultman L., Sundgren J.E., Greene J.E., Bergstrom D.R. & Petrov I. (1995). J.Appl. Phys., Vol.

Kelly E.J. (1979). Anodic dissolution and passivation of titanium in acidic media. III.

Kiparisov S.S. & Levinskiy Yu. V. (1972). Nitriding of refractory metals [in Russian],

Kolotyrkin Ya.N., Novakovskiy V.N., Kuznetsova Ye. G. at al. (1982). Corrosion behavior of

Matychak Ya., Fedirko V., Prytula A. & Pohrelyuk I. (2007). Modeling of diffusion saturation

Matychak Ya., Fedirko V., Pohrelyuk I., Yaskiv O. & Tkachuk O. (2008). Modelling of

Matychak Ya. S., Pohrelyuk I. M. & Fedirko V. M. (2009). Thermodiffusion saturation of α-

Matychak Ya. S., Pohrelyuk I. M. & Fedirko V. M. (2011). Kinetic features of the process of

Nazarenko P. V., Polishchuk I. E. & Molyar A. G. (1998). Tribological properties of coatings

Pohreliuk I., Yaskiv O. & Fedirko V. (2007). Formation of carbonitride coatings on titanium

titanium in technological media of chemical industry [in Russian], NIITEChIM,

of titanium by interstitial elements under rarefied atmospheres. Defect and

diffusion saturation of (α+β) titanium alloys by nitrogen under rarefied medium.

titanium with nitrogen from a rarefied atmosphere. Materials Science, Vol. 45, No.

nitriding of (α + β)-titanium alloys, Materials Science, Vol. 46, No. 5, pp. 660-668,

on titanium alloys. Fiz.-Khim. Mekh. Mater., Vol. 34, No. 2, pp. 55-62 (in

through thermochemical treatment from carbon-nitrogen-oxygen-containing media. JOM Journal of the Minerals, Metals and Materials Society, Vol. 59, No. 6,

system. Zaschita metallov, Vol. 19, No. 2., pp. 253– 256 (in Russian)

Chloride solutions. J. Electrochem. Soc., Vol. 126, pp. 2064-2075

Zaschita metallov, No. 2, pp. 223- 230 (in Russian)

Induced Wear [in Russian], pp. 7-80, Nauka, Moscow Goldschmidt H. J. (1967). Interstitial alloys, Plenum Press, New York

Naukova Dumka, Kiev

Mechanical engineering, Moscow

Diffusion Forum, Vol. 261-262, pp. 47-54

1, pp. 72-83, ISSN: 1068-820X

ISSN: 1068-820X

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pp. 32–37

Defect and Diffusion Forum, Vol. 277, pp. 29-34

pp. 119–121

78., p. 5395

Moscow

Metallurgiya, Мoscow

electrochemical behavior of nitride layers on titanium surface in sulphuric acid.

alloys in aqueous solutions of hydrochloric acid. Materials Science, Vol. 34, No. 1.,

hydrogenation on the wear of titanium alloys, In: Investigation of Hydrogen-

Fig. 24. Kinetics of the mass losses of the counterbody of BrAZh9-4l bronze in the process of friction with nitrided (1, 2) and oxidized (3, 4) disks of VT1-0 titanium: 1, 3 – 1 MPa; 2, 4 – 2 MPa.

In the former case, the wear of the bronze shoe after fitting and the first reference section (200 and 1000 m) first increases, then stabilizes at a level of 0,12 g (0,21 g) under a contact load of 1 MPa (2 MPa) and, finally, begins to increase again. In the latter case, the wear of the bronze is different for different sections of the friction path and under different stresses. Thus, under a load of 1 MPa, the mass loss of the bronze shoe slightly increases after fitting and then stabilizes. Only at the very end of the tests, we observe a weak tendency to increase in the mass losses. Under a load of 2 MPa, the wear of bronze significantly increases after fitting, then decreases and, finally, stabilizes.

It seems possible that the intense wear of the counterbody in the process of friction against the nitrided or oxidized titanium surface is explained by the hydrogenation of bronze as a result of tribodestruction of the lubricant (Goldfain et al., 1977). This observation is confirmed by the formation of bronze powder, which may be caused by the dispersion of hydrated bronze.

Thus, as compared with recommended nitriding, the application of oxidation as a method for increasing the wear resistance of titanium alloys is reasonable only under low contact stresses (≤ 1 MPa). Under higher contact stresses, the processes of oxidation and nitriding are characterized by practically equal levels of energy consumption and productivity but the tribological and mechanical characteristics of nitrided titanium are better than the corresponding characteristics of oxidized titanium.

#### **8. References**

Brynza A.P. & Fedash V.P. (1972). About passivation mechanism of titanium in solutions of sulphuric and chloric acids, In: Noviy konstrukcyonniy material- titanium [in Russian], pp. 179-183, Nauka, Мoscow

Fig. 24. Kinetics of the mass losses of the counterbody of BrAZh9-4l bronze in the process of

In the former case, the wear of the bronze shoe after fitting and the first reference section (200 and 1000 m) first increases, then stabilizes at a level of 0,12 g (0,21 g) under a contact load of 1 MPa (2 MPa) and, finally, begins to increase again. In the latter case, the wear of the bronze is different for different sections of the friction path and under different stresses. Thus, under a load of 1 MPa, the mass loss of the bronze shoe slightly increases after fitting and then stabilizes. Only at the very end of the tests, we observe a weak tendency to increase in the mass losses. Under a load of 2 MPa, the wear of bronze significantly increases

It seems possible that the intense wear of the counterbody in the process of friction against the nitrided or oxidized titanium surface is explained by the hydrogenation of bronze as a result of tribodestruction of the lubricant (Goldfain et al., 1977). This observation is confirmed by the formation of bronze powder, which may be caused by the dispersion of hydrated bronze.

Thus, as compared with recommended nitriding, the application of oxidation as a method for increasing the wear resistance of titanium alloys is reasonable only under low contact stresses (≤ 1 MPa). Under higher contact stresses, the processes of oxidation and nitriding are characterized by practically equal levels of energy consumption and productivity but the tribological and mechanical characteristics of nitrided titanium are better than the

Brynza A.P. & Fedash V.P. (1972). About passivation mechanism of titanium in solutions of

sulphuric and chloric acids, In: Noviy konstrukcyonniy material- titanium [in

friction with nitrided (1, 2) and oxidized (3, 4) disks of VT1-0 titanium: 1, 3 – 1 MPa;

after fitting, then decreases and, finally, stabilizes.

corresponding characteristics of oxidized titanium.

Russian], pp. 179-183, Nauka, Мoscow

2, 4 – 2 MPa.

**8. References** 


**8** 

*Poland* 

Elzbieta Krasicka-Cydzik *University of Zielona Gora* 

**Anodic Layer Formation on Titanium** 

**and Its Alloys for Biomedical Applications** 

Properties of the oxide layers on titanium and its implant alloys can be tailored to desired applications by anodizing parameters. Electrochemical oxidation in various electrolytes and different polarization regimes may shape the morphology, structure and chemical composition of oxide layers to enhance the use of titanium materials in electronics, photovoltaic and medicine. Phosphate electrolytes play specific role in the anodizing process. Besides forming compact barrier layer they enable also to form porous and nanostructural oxide layers enriched with phosphates, which enhance their bioactivity.

The formation of anodic layers: thick or thin, compact or porous, gel-like and nanostructural on titanium and its alloys Ti6Al4V and Ti6Al7Nb in phosphoric acid solutions of different concentrations is described in this charter. Basing on morphological and chemical composition analysis (SEM, XPS) as well as on the electrochemical examination the influence of electrolyte composition on enrichment of surface oxide layers with phosphates and fluorides, enhancing their bioactivity, is presented. Studies to use Ti/titania systems as the platforms of the electrochemical biosensors to detect H2O2 and glucose proved the opportunity to use the nanotubular titania material as a platform for the 2rd generation

**2. Oxide layers on titanium and its implant alloys formed in H3PO4 solutions**  Anodic films formed on titanium and its alloys are of great interest due to the industrial applications of metal covered with oxide layers of various and unique properties [1-7]. These layers have been investigated extensively by many authors [8-11]. Thick oxide layers on titanium, obtained by anodizing, provide improved resistance to local corrosion [12]. Anodizing can result in the adsorption and incorporation of inorganic and organic, biologically important species, e.g. phosphate ions, into the oxide layer. Such surface layers, desirable for medical implants, are not only corrosion resistant in a biological environment, but also compatible with tissue response [13-15]. Anodizing titanium and its alloys has been investigated in a wide range of parameters [16-20], which include also the participation of

the electrolyte components, e.g. anions, in the formation of anodic films [21, 22]. At anodizing oxide layers are formed according to the following reaction [17, 23]:

**1. Introduction** 

biosensors.


### **Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications**

Elzbieta Krasicka-Cydzik *University of Zielona Gora Poland* 

#### **1. Introduction**

174 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Pohrelyuk I., Yaskiv O., Tkachuk O. & Lee Dong Bok. (2009). Formation of oxynitride layers

Pohrelyuk I. M., Tkachuk O. V. & Proskurnyak R. V. (2011). Corrosion resistance of the

Metals and Materials Society, Vol. 63, No. 6, pp. 35-40, ISSN: 1047-4838 Rolinski E. (1988). Effect of nitriding on the surface structure of titanium. J. Less-Common

Samsonov G.V., Epik А.P. (1973). Refractory coatings [in Russian], Metallurgiya, Мoscow Strafford K. N. (1979). A comparison of the high temperature nitridation and oxidation

Sukhotin V.N., Lygin S.A., Sunguchayeva I.A. & Gramanschikov V.A. (1990). Passivation of

Tomashov N.D., Chukalovskaya T.V., Myedova I.L. & Yegorov F.F. (1985). Corrosion and

Zavarov A. S., Baskakov A. P., & S. V. Grachev. (1985). Thermochemical Boiling-Bed

phosphoric acids. Zaschyta myetallov, No. 5., pp. 682-688 (in Russian) Yaskiv O., Pohrelyuk I., Fedirko V.M., Tkachuk O. & Lee Dong-Bok. (2011). Formation of

behaviour of metals. Corn Sci., Vol. 19, No. 1, pp. 49-62

Treatment [in Russian], Mashinostroenie, Moscow

Vol. 15, No 6, pp. 949–953., ISSN: 1598-9623

Metals, Vol. 141, No. 1., pp. L11 - L14.

(in Russian)

519, No19, pp. 6508 – 6514

on titanium alloys by gas diffusion treatment. Metals and Materials International,

nitrated Ti-6Al-4V titanium alloy in 0,9% NaCl, JOM Journal of the Minerals,

titanium in solutions of chloric acid. Zaschita metallov, Vol. 26, No. 1., pp. 128-131

anodic behaviour of carbide, nitride and boride of titanium in sulphuric and

oxynitrides on titanium alloys by gas diffusion treatment. Thin Solid Films, - Vol.

Properties of the oxide layers on titanium and its implant alloys can be tailored to desired applications by anodizing parameters. Electrochemical oxidation in various electrolytes and different polarization regimes may shape the morphology, structure and chemical composition of oxide layers to enhance the use of titanium materials in electronics, photovoltaic and medicine. Phosphate electrolytes play specific role in the anodizing process. Besides forming compact barrier layer they enable also to form porous and nanostructural oxide layers enriched with phosphates, which enhance their bioactivity.

The formation of anodic layers: thick or thin, compact or porous, gel-like and nanostructural on titanium and its alloys Ti6Al4V and Ti6Al7Nb in phosphoric acid solutions of different concentrations is described in this charter. Basing on morphological and chemical composition analysis (SEM, XPS) as well as on the electrochemical examination the influence of electrolyte composition on enrichment of surface oxide layers with phosphates and fluorides, enhancing their bioactivity, is presented. Studies to use Ti/titania systems as the platforms of the electrochemical biosensors to detect H2O2 and glucose proved the opportunity to use the nanotubular titania material as a platform for the 2rd generation biosensors.

#### **2. Oxide layers on titanium and its implant alloys formed in H3PO4 solutions**

Anodic films formed on titanium and its alloys are of great interest due to the industrial applications of metal covered with oxide layers of various and unique properties [1-7]. These layers have been investigated extensively by many authors [8-11]. Thick oxide layers on titanium, obtained by anodizing, provide improved resistance to local corrosion [12]. Anodizing can result in the adsorption and incorporation of inorganic and organic, biologically important species, e.g. phosphate ions, into the oxide layer. Such surface layers, desirable for medical implants, are not only corrosion resistant in a biological environment, but also compatible with tissue response [13-15]. Anodizing titanium and its alloys has been investigated in a wide range of parameters [16-20], which include also the participation of the electrolyte components, e.g. anions, in the formation of anodic films [21, 22].

At anodizing oxide layers are formed according to the following reaction [17, 23]:

$$\text{ZnMe} + \text{m} \to \text{H}\_2\text{O} \to \text{Me}\_n\text{O}\_m + 2 \text{ m} \to \text{H}^\* + 2 \text{m} \,\text{e} \tag{1}$$

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 177

polarization parameters however layers of different thickness and colorization are formed [11]. Due to the presence of phosphates in the anodic layers they are highly bioactive in comparison to oxides formed in other electrolytes. Just after 9 days in the SBF solution

a b c Fig. 2. Porous titania layers on Ti formed at 0.5 A/m2 in 0.5 M H3PO4 [22] (a,b) and HAp particles on anodic layer after 9 days in SBF solution [21, 25, 27] (c), JEOL JLM 5600 EDS

The bilayer structure of compact oxide covered with HAp particles can be demonstrated also in the impedance tests (Fig. 3). The first time constant in Bode diagrams in range of the high and intermediate frequencies confirmed the high *Rt* resistance of the barrier layer covering the anodized metal, thus giving the evidence of its high corrosion resistance. The second time constant corresponds to porous layer above the barrier one. Its lack in case of the Ti6Al4V ELI

Fig. 3. EIS spectra recorded after 9 days in SBF solution [21, 25, 27] for anodic layers on Ti,

To investigate the effect of phosphoric acid concentration (0.5 - 4 M) on the anodising of titanium and its alloys the galvanostatic and potentiodynamic techniques have been applied [11,19]. Particularly, the galvanostatic method with low current densities, up to 0.6 Am-2,

Ti6Al4V and Ti6Al7Nb formed at 0.5 A/m2 in 0.5 M H3PO4 [22]

indicates the different characteristics of the coating on this implant material.

(Simulated Body Fluid) they are covered with hydroxyapatite deposists (Fig. 2) [27].

Studies in this field [24, 25] have shown that, phosphate ions can be incorporated into the anodic layer on titanium and Ti-6Al-4V, and in turn stimulate the formation of the biocompatible hydroxyapatite [26].

Anodizing in phosphate solutions exhibit some advantages over other acid and base electrolytes. First of all less corrosive attack of phosphoric acid on titanium and its alloys, when compared with other acidic media, is related to the strong adsorption of phosphate anions on the surface [27,28]. Although unalloyed titanium is resistant [11] to naturally aerated pure solutions of phosphoric acid up to 30 % wt. concentration (~3.6 M) in phosphoric acid of lower concentrations (0.5-4 M), mainly non-dissociated acid molecules and H2PO4 of phosphate ions exist [29,30] and they exhibit a strong affinity or complexing power towards most metal cations.

Fig. 1. Adsorption of phosphates onto TiO2 [28] and effect of pH on phosphoric acid composition [30]

Thus, anodizing in phosphate solutions leads to the incorporation of phosphate ions into the oxide layers on titanium and Ti-6Al-4V [5-8] influencing their bioactivity and stimulating deposition of the biocompatible hydroxyapatite. The latter may be used to shape properties of titanium implant materials for medical purposes.

#### **2.1 Thin porous anodic layers**

Anodic layers on pure Ti and its alloys Ti6Al4V ELI and Ti6Al7Nb (ASTM F136-84), alloys in the annealed condition) can be formed by anodizing carried out at ambient pressure and room temperature in non-deaerated electrolyte solutions of 0.5 M H3PO4. Both techniques, the galvanostatic at anodizing current density values varied in the range of 0.1-0.5 Am-2 and the potentiostatic at up to 60V [22] are used. The oxide layers, 30-120nm thick, enriched with phosphorus are formed in these conditions. With mechanical and chemical pre-treatment applied to titanium and its alloys: Ti6Al4V ELI and Ti6Al7Nb (Timet Ltd, UK) [12-20], the layers obtained at 60V in phosphoric acid are golden and porous [11, 24-26] (Fig. 2) and they show very stable values of currents in passive region up to 1 V (SCE) (Fig.3) [11]. At other

Studies in this field [24, 25] have shown that, phosphate ions can be incorporated into the anodic layer on titanium and Ti-6Al-4V, and in turn stimulate the formation of the bio-

Anodizing in phosphate solutions exhibit some advantages over other acid and base electrolytes. First of all less corrosive attack of phosphoric acid on titanium and its alloys, when compared with other acidic media, is related to the strong adsorption of phosphate anions on the surface [27,28]. Although unalloyed titanium is resistant [11] to naturally aerated pure solutions of phosphoric acid up to 30 % wt. concentration (~3.6 M) in phosphoric acid of lower concentrations (0.5-4 M), mainly non-dissociated acid molecules

Fig. 1. Adsorption of phosphates onto TiO2 [28] and effect of pH on phosphoric acid

of titanium implant materials for medical purposes.

**2.1 Thin porous anodic layers** 

Thus, anodizing in phosphate solutions leads to the incorporation of phosphate ions into the oxide layers on titanium and Ti-6Al-4V [5-8] influencing their bioactivity and stimulating deposition of the biocompatible hydroxyapatite. The latter may be used to shape properties

Anodic layers on pure Ti and its alloys Ti6Al4V ELI and Ti6Al7Nb (ASTM F136-84), alloys in the annealed condition) can be formed by anodizing carried out at ambient pressure and room temperature in non-deaerated electrolyte solutions of 0.5 M H3PO4. Both techniques, the galvanostatic at anodizing current density values varied in the range of 0.1-0.5 Am-2 and the potentiostatic at up to 60V [22] are used. The oxide layers, 30-120nm thick, enriched with phosphorus are formed in these conditions. With mechanical and chemical pre-treatment applied to titanium and its alloys: Ti6Al4V ELI and Ti6Al7Nb (Timet Ltd, UK) [12-20], the layers obtained at 60V in phosphoric acid are golden and porous [11, 24-26] (Fig. 2) and they show very stable values of currents in passive region up to 1 V (SCE) (Fig.3) [11]. At other

of phosphate ions exist [29,30] and they exhibit a strong affinity or complexing

compatible hydroxyapatite [26].

power towards most metal cations.

and H2PO4-

composition [30]

n Me + m H2O MenOm + 2 m H+ + 2m *e*- (1)

polarization parameters however layers of different thickness and colorization are formed [11]. Due to the presence of phosphates in the anodic layers they are highly bioactive in comparison to oxides formed in other electrolytes. Just after 9 days in the SBF solution (Simulated Body Fluid) they are covered with hydroxyapatite deposists (Fig. 2) [27].

Fig. 2. Porous titania layers on Ti formed at 0.5 A/m2 in 0.5 M H3PO4 [22] (a,b) and HAp particles on anodic layer after 9 days in SBF solution [21, 25, 27] (c), JEOL JLM 5600 EDS

The bilayer structure of compact oxide covered with HAp particles can be demonstrated also in the impedance tests (Fig. 3). The first time constant in Bode diagrams in range of the high and intermediate frequencies confirmed the high *Rt* resistance of the barrier layer covering the anodized metal, thus giving the evidence of its high corrosion resistance. The second time constant corresponds to porous layer above the barrier one. Its lack in case of the Ti6Al4V ELI indicates the different characteristics of the coating on this implant material.

Fig. 3. EIS spectra recorded after 9 days in SBF solution [21, 25, 27] for anodic layers on Ti, Ti6Al4V and Ti6Al7Nb formed at 0.5 A/m2 in 0.5 M H3PO4 [22]

To investigate the effect of phosphoric acid concentration (0.5 - 4 M) on the anodising of titanium and its alloys the galvanostatic and potentiodynamic techniques have been applied [11,19]. Particularly, the galvanostatic method with low current densities, up to 0.6 Am-2,

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 179

M H3PO4. Polarization curves (Fig.5b), show that after an initial range of cathodic depassivation, samples reach the corrosion potential *Ecor*. Then an active–passive transition is observed with passivating currents the order of a few microamperes, which are typically observed during the passivation of titanium and its alloys. Anodic curves for 0.5 M and 3 M H3PO4 solutions (Fig.5b), illustrate quasi-passive behaviour in active-passive transitions, while curves for 1 and 2 M H3PO4, having the higher corrosion potential *Ecor*, do not show linear dependence in this potential region. The differences in anodic Tafel slopes are accompanied by a shift of the *Ecor* value in the positive direction by ~0.15V with the increase

Fig. 5. Corrosion potential values Ecor of Ti anodised at 0.5 A/m2 in 0.5-2 M H3PO4 in SBF solution [11,29] (a) and active-passive transition regions of polarization curves for titanium

SEM/EDS examination confirm the evidence of two-layered surface film. Such layers presented in Fig. 6, show the whole surface covered by a gel-like layer of H3PO4×0.5H2O.

of H3PO4 concentration to 2 M (Fig. 5b).

(scan 3 mV/s)in 0.5-4 M H3PO4 (b)

applied in order to minimize side effects (ie. oxygen evolution), and more importantly to determine processes responsible for the growth of the oxide layer on the anodised metal at the early stages of its formation, allowed to observe the abnormal behaviour of titanium at anodizing. In Fig. 4 the results presenting the minimum rates of potential growth at initial stages of galvanostatic anodising in 2 M H3PO4 solutions are shown.

Fig. 4. Surface response for the investigation of the effect of H3PO4 concentration and polarization current on the rate of potential growth dE/dt at galvanostatic anodizing of titanium [11]

Also the anodic polarization curves for titanium in electrolytes of different concentration show various shapes and different slopes in active-passive region (Fig.5).Potentiodynamic control at a comparable rate in 2 M H3PO4 applied to titanium and two of its implant alloys, Ti6Al4V ELI and Ti6Al7Nb, revealed a shift in corrosion potential toward the anodic direction with the lowest current densities in the passive region. This was possibly due to the effect of adsorption of phosphate ions onto the surface layer.

#### **2.2 Gel like layers on anodic titania**

Active-passive transition of titanium in phosphoric acid solutions of 0.5-4 M [11,19] reveals that the growth of anodic layer is affected proportionally by the applied anodic potential, but shows the unusual influence of electrolyte concentration. Under galvanostatic conditions at low current densities (0.1-0.6 Am-2) the slope of dE/dt shows the minimum at the concentration ~2 M H3PO4 (Fig. 4), which is resulted due to a coating of an oxide film by an additional gel-like layer during anodizing, [28, 29], similar to the one observed in other media on aluminum.

It was found that the active–passive transition was a process in which an inhibiting effect of phosphate ions on a dissolution of oxide layer was observed during anodizing [29]. Typical examples of voltage *vs* time transients (d*E/*d*t*) during the growth of the galvanostatic anodic oxide film on titanium for the current density of 0.5 A/m2 (Fig. 5a) show that the continuous linear growth of potential to the steady state, demonstrates the lowest value in 2

applied in order to minimize side effects (ie. oxygen evolution), and more importantly to determine processes responsible for the growth of the oxide layer on the anodised metal at the early stages of its formation, allowed to observe the abnormal behaviour of titanium at anodizing. In Fig. 4 the results presenting the minimum rates of potential growth at initial

Fig. 4. Surface response for the investigation of the effect of H3PO4 concentration and polarization current on the rate of potential growth dE/dt at galvanostatic anodizing of

the effect of adsorption of phosphate ions onto the surface layer.

**2.2 Gel like layers on anodic titania** 

media on aluminum.

Also the anodic polarization curves for titanium in electrolytes of different concentration show various shapes and different slopes in active-passive region (Fig.5).Potentiodynamic control at a comparable rate in 2 M H3PO4 applied to titanium and two of its implant alloys, Ti6Al4V ELI and Ti6Al7Nb, revealed a shift in corrosion potential toward the anodic direction with the lowest current densities in the passive region. This was possibly due to

Active-passive transition of titanium in phosphoric acid solutions of 0.5-4 M [11,19] reveals that the growth of anodic layer is affected proportionally by the applied anodic potential, but shows the unusual influence of electrolyte concentration. Under galvanostatic conditions at low current densities (0.1-0.6 Am-2) the slope of dE/dt shows the minimum at the concentration ~2 M H3PO4 (Fig. 4), which is resulted due to a coating of an oxide film by an additional gel-like layer during anodizing, [28, 29], similar to the one observed in other

It was found that the active–passive transition was a process in which an inhibiting effect of phosphate ions on a dissolution of oxide layer was observed during anodizing [29]. Typical examples of voltage *vs* time transients (d*E/*d*t*) during the growth of the galvanostatic anodic oxide film on titanium for the current density of 0.5 A/m2 (Fig. 5a) show that the continuous linear growth of potential to the steady state, demonstrates the lowest value in 2

titanium [11]

stages of galvanostatic anodising in 2 M H3PO4 solutions are shown.

M H3PO4. Polarization curves (Fig.5b), show that after an initial range of cathodic depassivation, samples reach the corrosion potential *Ecor*. Then an active–passive transition is observed with passivating currents the order of a few microamperes, which are typically observed during the passivation of titanium and its alloys. Anodic curves for 0.5 M and 3 M H3PO4 solutions (Fig.5b), illustrate quasi-passive behaviour in active-passive transitions, while curves for 1 and 2 M H3PO4, having the higher corrosion potential *Ecor*, do not show linear dependence in this potential region. The differences in anodic Tafel slopes are accompanied by a shift of the *Ecor* value in the positive direction by ~0.15V with the increase of H3PO4 concentration to 2 M (Fig. 5b).

Fig. 5. Corrosion potential values Ecor of Ti anodised at 0.5 A/m2 in 0.5-2 M H3PO4 in SBF solution [11,29] (a) and active-passive transition regions of polarization curves for titanium (scan 3 mV/s)in 0.5-4 M H3PO4 (b)

SEM/EDS examination confirm the evidence of two-layered surface film. Such layers presented in Fig. 6, show the whole surface covered by a gel-like layer of H3PO4×0.5H2O.

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 181

interface Cb. In the case of surface layer formed in 2 M H3PO4 the specimen is covered by a passive oxide film of higher impedance (Fig.7a). The significant increase of the resistance Rb and Rp values for the 2 M H3PO4 anodized samples, over those determined for the 0.5 M H3PO4 anodized titanium, confirm that the EIS results are complementary to those obtained

Titanium as metal very sensitive to the pre-treatment [33], due to the polishing, rinsing with water and drying, is usually covered by an air-formed oxide film and on immersion into acid solution shows potentials of the active-transition region [32-34]. However, in solutions of low pH may become active. In sulphate solution the anodic oxide film on titanium dissolves giving Ti3+ ions [11]. Typical activation behaviour and slow decrease in the opencircuit potential (–0.3 V SCE), is observed on immersing titanium into 1M HCl [35,36]. Titanium behaves differently in H3PO4 solutions. Although the values of *Ecor* ranging from - 0.1V to -0.6V (SCE) indicate in Pourbaix diagram [29] that the oxide film should dissolve to Ti3+, on immersion to H3PO4 solutions titanium shows the continuous shift of potential towards the anodic direction [32]. Such a tendency indicates that the rate of the anodic reactions is continually decreasing, as a result of the presence of an adsorbed, additional layer on the metal surface [37]. Thermodynamic data [11,33] for the potential E –0.8V

= 2 Ti2+ +3 H2O Eº = - 0.478 V - 0.177pH – 0.059 log(Ti2+) (1)

2 Ti(OH)3 + H2O=2 TiO2H2O + 2 H+ + 2e- Eº = -0.091V–0.059 pH (4)

2 H2PO4- + 2 H3O+ + (*n*-2) H2O 2 H3PO4 × *n* H2O (5)

In acidic solutions within the cathodic region, the only oxide dissolution is reaction (1). This reaction determines the potential-current changes in active-passive region of titanium in 0.5 M and 3-4 M H3PO4, whereas due to the shift of the corrosion potential *Ecor* towards the anodic direction for titanium immersed in 1-2 M H3PO4 electrolyte, its direct oxidation proceeds according to reactions 3 and 4. These results indicate that, the layer of phosphates (Fig. 6) blocks the oxide dissolution. Insight into the adsorption of phosphates to TiO2 surface revealed the hypothesis according to which they form covalent bonding to oxygen [38], or metal ions react with phosphate anions forming a gel of metal hydrophosphates [39]. Both proposed processes would lead to local increase of pH at the oxide surface and in consequence to the increase of concentration of dihydro-phosphate ions (E-pH diagram) [29]. Then, on phosphates covered titanium oxide electrode, the following gel like layer

This attribution agrees with Morligde's *et. al*. results on aluminum [40]. Both reactions: the phosphates adsorption and gel-like layer formation are non-faradaic, but are competitive towards the oxide dissolution (reaction 1) with regard to proton consumption. The advantageous effect of these two reactions on the anodizing, may be attributed to an inadequate supply of H+ ions to keeping up with the demand for the reaction (1) of oxide dissolution. The increasing coverage of the anodized titanium surface by phosphates ions

Eº = - 0.556 V – 0.059 pH (2)

Eº = - 0.139V – 0.059 pH (3)

by *Ecor* measurements and potentiodynamic polarization studies [11,31].

(SCE), indicate that the following reactions on titanium are likely:

Ti2O3 + H2O = 2 TiO2 + 2 H+ + 2 e-

Ti2O3 + 3 H2O = 2 TiO2·H2O + 2 H+ + 2 e-

Ti2O3 + 6 H+ + 2e-

formation could proceed

The SEM/EDS examinations reveal that thin films of anodic titania oxide are covered by gellike layer with crystalline phosphates nuclei inside. Phosphates deposits are few in layers formed in 0.5 M H3PO4, but numerous and uniformly dispersed in a surface oxide of sample anodised in 2 M H3PO4. However, the oxide and phosphates are covered with the additional layer consisting of 76.33.6 wt.% of phosphorus and 23.71.5wt.% of oxygen (Fig.6a and 6b).

Fig. 6. SEM micrographs of titanium surface anodized at 0.5 A/m2 in a,b) 0.5 M, c) 2 M H3PO4 [29], and d) HAp particles on anodic layer after 9 days in SBF solution [21, 27]

The EIS spectra of titanium after anodizing at 0.5 A/m2 in 0.5-2 M H3PO4 (Fig. 7), exhibit a behavior typical of a metallic material covered by a porous film which is exposed to an electrolytic environment [6]. Two time constants are seen in the spectra: the first in the highfrequency part arises from the ohmic electrolyte resistance and the impedance resulting from the penetration of the electrolyte through a porous film, and the second in lowfrequency part accounts for the processes at the substrate/electrolyte interface.

Fig. 7. Impedance spectra for titanium anodised in 2 M H3PO4 exposed to 0.9% NaCl solution a) Nyquist spectra, b) Bode diagrams and results oftheir fitting to c) equivalent circuit

The EIS data can be fitted to the equivalent circuit in Fig. 7c, which consists of a solution resistance Rs, the capacitance Cp of the barrier layer, the charge transfer resistance associated with the penetration of the electrolyte through the pores Rp; and the polarization resistance of the barrier Rb as well as the electrical double-layer capacitance at the substrate/electrolyte

The SEM/EDS examinations reveal that thin films of anodic titania oxide are covered by gellike layer with crystalline phosphates nuclei inside. Phosphates deposits are few in layers formed in 0.5 M H3PO4, but numerous and uniformly dispersed in a surface oxide of sample anodised in 2 M H3PO4. However, the oxide and phosphates are covered with the additional layer consisting of 76.33.6 wt.% of phosphorus and 23.71.5wt.% of oxygen (Fig.6a and 6b).

Fig. 6. SEM micrographs of titanium surface anodized at 0.5 A/m2 in a,b) 0.5 M, c) 2 M H3PO4 [29], and d) HAp particles on anodic layer after 9 days in SBF solution [21, 27]

frequency part accounts for the processes at the substrate/electrolyte interface.

The EIS spectra of titanium after anodizing at 0.5 A/m2 in 0.5-2 M H3PO4 (Fig. 7), exhibit a behavior typical of a metallic material covered by a porous film which is exposed to an electrolytic environment [6]. Two time constants are seen in the spectra: the first in the highfrequency part arises from the ohmic electrolyte resistance and the impedance resulting from the penetration of the electrolyte through a porous film, and the second in low-

Fig. 7. Impedance spectra for titanium anodised in 2 M H3PO4 exposed to 0.9% NaCl solution a) Nyquist spectra, b) Bode diagrams and results oftheir fitting to c) equivalent circuit

The EIS data can be fitted to the equivalent circuit in Fig. 7c, which consists of a solution resistance Rs, the capacitance Cp of the barrier layer, the charge transfer resistance associated with the penetration of the electrolyte through the pores Rp; and the polarization resistance of the barrier Rb as well as the electrical double-layer capacitance at the substrate/electrolyte interface Cb. In the case of surface layer formed in 2 M H3PO4 the specimen is covered by a passive oxide film of higher impedance (Fig.7a). The significant increase of the resistance Rb and Rp values for the 2 M H3PO4 anodized samples, over those determined for the 0.5 M H3PO4 anodized titanium, confirm that the EIS results are complementary to those obtained by *Ecor* measurements and potentiodynamic polarization studies [11,31].

Titanium as metal very sensitive to the pre-treatment [33], due to the polishing, rinsing with water and drying, is usually covered by an air-formed oxide film and on immersion into acid solution shows potentials of the active-transition region [32-34]. However, in solutions of low pH may become active. In sulphate solution the anodic oxide film on titanium dissolves giving Ti3+ ions [11]. Typical activation behaviour and slow decrease in the opencircuit potential (–0.3 V SCE), is observed on immersing titanium into 1M HCl [35,36]. Titanium behaves differently in H3PO4 solutions. Although the values of *Ecor* ranging from - 0.1V to -0.6V (SCE) indicate in Pourbaix diagram [29] that the oxide film should dissolve to Ti3+, on immersion to H3PO4 solutions titanium shows the continuous shift of potential towards the anodic direction [32]. Such a tendency indicates that the rate of the anodic reactions is continually decreasing, as a result of the presence of an adsorbed, additional layer on the metal surface [37]. Thermodynamic data [11,33] for the potential E –0.8V (SCE), indicate that the following reactions on titanium are likely:

$$\text{Ti}\_2\text{O}\_7 + 6\text{ H}^\* + 2\text{e}^\cdot = 2\text{ Ti}2^\circ + 3\text{ H}\_2\text{O} \quad \text{E}^\circ = -0.478\text{ V} - 0.177\text{pH} - 0.059\text{log(Ti}^2) \tag{1}$$

$$\text{Ti}\_2\text{O}\_5 + \text{H}\_2\text{O} = 2\text{TiO}\_2 + 2\text{ H}^+ + 2\text{ e}^- \to \text{ $\mathcal{C}^=-$ }.0.556\text{ V} - 0.059\text{ pH} \tag{2}$$

$$\text{Ti}\_2\text{O}\_3 + 3\text{ H}\_2\text{O} = 2\text{TiO}\_2\\\text{H}\_2\text{O} + 2\text{ H}^\* + 2\text{ e}^- \to \text{P}-0.139\text{V}-0.059\text{ pH}\tag{3}$$

$$2\text{ Ti(OH)}\_{3} + \text{H}\_{2}\text{O=2 TiO}\_{2}\text{H}\_{2}\text{O} + 2\text{ H}^{\*} + 2\text{e}^{-} \to \text{2e}^{\cdot} \text{ E}^{\circ} = \text{-0.091V-0.059 pH} \tag{4}$$

In acidic solutions within the cathodic region, the only oxide dissolution is reaction (1). This reaction determines the potential-current changes in active-passive region of titanium in 0.5 M and 3-4 M H3PO4, whereas due to the shift of the corrosion potential *Ecor* towards the anodic direction for titanium immersed in 1-2 M H3PO4 electrolyte, its direct oxidation proceeds according to reactions 3 and 4. These results indicate that, the layer of phosphates (Fig. 6) blocks the oxide dissolution. Insight into the adsorption of phosphates to TiO2 surface revealed the hypothesis according to which they form covalent bonding to oxygen [38], or metal ions react with phosphate anions forming a gel of metal hydrophosphates [39]. Both proposed processes would lead to local increase of pH at the oxide surface and in consequence to the increase of concentration of dihydro-phosphate ions (E-pH diagram) [29]. Then, on phosphates covered titanium oxide electrode, the following gel like layer formation could proceed

$$2\,\mathrm{H\_2PO\_4} + 2\,\mathrm{H\_3O^+} + \text{(n-2)}\,\mathrm{H\_2O} \to 2\,\mathrm{H\_3PO\_4} \times \,\mathrm{n}\,\mathrm{H\_2O} \tag{5}$$

This attribution agrees with Morligde's *et. al*. results on aluminum [40]. Both reactions: the phosphates adsorption and gel-like layer formation are non-faradaic, but are competitive towards the oxide dissolution (reaction 1) with regard to proton consumption. The advantageous effect of these two reactions on the anodizing, may be attributed to an inadequate supply of H+ ions to keeping up with the demand for the reaction (1) of oxide dissolution. The increasing coverage of the anodized titanium surface by phosphates ions

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 183

SEM observations (Fig.8) and EDS microanalysis indicate the presence of deposits dispersed on the surface of anodised titanium and its implant alloys. However, deposits are are nonuniformly dispersed on a surface. Titanium and its two alloys anodised in 0.5M H3PO4 are covered with very thin oxide layer, which includes numerous and more scattered Ca-O-P deposits of diameter varied from 200 to 800 nm, suggesting the heterogeneous nucleation of Ca-O-P on TiO2 covered surface. Althogh just after 24 hours deposits are seen on the surface of the 3 materials (more deposists on the Ti6Al4V alloy) the continuous films cover the whole surfaces after 9 days in SBF solution. At higher magnification it is seen that the film on titanium is formed of more flatter layer of deposits and broken layer of titanium oxides with titanium phosphates, whereas film on both alloys comprise small globules of Ca-O-P. The ratio

of Ca/P ranging from 1.26 to 1.42 corresponds to non-steochiometric hydroxyapatite.

Fig. 9. EIS) spectra recorded during immersion of titanium and Ti6Al6V samples anodized

Impedance (EIS) spectra (Fig. 9) recorded during immersion of the anodized titanium and the Ti6Al4V alloy in simulated body fluid (SBF) for titanium and Ti6Al4V alloy show changes in capacitance and structure of surface layers as well as differences between coatings on titanium and its alloy and confirm the SEM observations (Fig. 8). Titanium exhibits two-layered structure: the inner oxide layer is covered by an outer layer of more or less uniformly distributed various size Ca-P-O deposits. On contrary the Ti6Al4V alloy is coated by a more uniform and dense layer of deposits (Fig. 8 c,d) and much lower

in 0.5 M and 2 M H3PO4 in SBF solution [21,23]

concentration of titanium oxide on a surface [25].

with the electrolyte concentration provides the evidence of a direct influence of electrolyte anions in suppressing the formation of dissolved titanium ions. According to potential/pH diagram for P-Ti-H2O system [41], H3PO4×0.5H2O, the product of reaction 5, is stable thermodynamically in solutions of pH ranging to 3.

Thus, due to the applied anodizing conditions formation of either thin and porous oxide layer [11-23] or gel-like phosphates rich layer of H3PO4×0.5H2O [11, 24-29, 42], covering thicker oxide layer on titanium can be obtained.

#### **2.3 Bioactive layer**

Apart from mechanical properties and biocompatibility, which make titanium and its alloys the materials of choice for various applications (artificial hip and knee joints, dental prosthetics, vascular stents, heart valves) also enhancement of bone formation is desired feature of a metallic implant developed through adequate surface treatments to obtain proper osseointegration.

Fast deposition of hydroxyapatite (HAp) coatings on titanium and its alloys Ti6Al4V and Ti6Al7Nb substrates anodised in H3PO4 was observed [21,23,25]. Anodizing in 0.5 M H3PO4, which produces phosphates enriched porous sub-surface layer on of titanium and its alloys Ti6Al4V and Ti6Al7N [22] or anodizing in 2 M H3PO4 which generates phosphates rich gellike layer [31,42] may be used to enhance hydroxyapatite (HAp) deposition (Fig.8). For the latter anodic layers soaking the anodised substrates in simulated body fluid (SBF) resulted in the deposition of a uniform coating in 24 hours (Fig. 9). SEM and EDS investigations revealed that after 9 days thick coating consists of HAp globular of diameter varied from 100 to 300 nm aggregates. The Ca-O-P deposits merge in large clusters and they are seen in large numbers on both alloys, particularly on Ti6Al4V anodized in 2 M H3PO4.

Fig. 8. SEM micrographs of titanium (a,b) and its alloys: Ti6Al4V (c,d) Ti6Al7Nb (e,f) surface anodized at 0.5 A/m2 in 0.5 M H3PO4 after 24 h (abc) and 9 days (b,d,f) in SBF solution [21, 23]

with the electrolyte concentration provides the evidence of a direct influence of electrolyte anions in suppressing the formation of dissolved titanium ions. According to potential/pH diagram for P-Ti-H2O system [41], H3PO4×0.5H2O, the product of reaction 5, is stable

Thus, due to the applied anodizing conditions formation of either thin and porous oxide layer [11-23] or gel-like phosphates rich layer of H3PO4×0.5H2O [11, 24-29, 42], covering

Apart from mechanical properties and biocompatibility, which make titanium and its alloys the materials of choice for various applications (artificial hip and knee joints, dental prosthetics, vascular stents, heart valves) also enhancement of bone formation is desired feature of a metallic implant developed through adequate surface treatments to obtain

Fast deposition of hydroxyapatite (HAp) coatings on titanium and its alloys Ti6Al4V and Ti6Al7Nb substrates anodised in H3PO4 was observed [21,23,25]. Anodizing in 0.5 M H3PO4, which produces phosphates enriched porous sub-surface layer on of titanium and its alloys Ti6Al4V and Ti6Al7N [22] or anodizing in 2 M H3PO4 which generates phosphates rich gellike layer [31,42] may be used to enhance hydroxyapatite (HAp) deposition (Fig.8). For the latter anodic layers soaking the anodised substrates in simulated body fluid (SBF) resulted in the deposition of a uniform coating in 24 hours (Fig. 9). SEM and EDS investigations revealed that after 9 days thick coating consists of HAp globular of diameter varied from 100 to 300 nm aggregates. The Ca-O-P deposits merge in large clusters and they are seen in

a c e

b d f

Fig. 8. SEM micrographs of titanium (a,b) and its alloys: Ti6Al4V (c,d) Ti6Al7Nb (e,f) surface anodized at 0.5 A/m2 in 0.5 M H3PO4 after 24 h (abc) and 9 days (b,d,f) in SBF solution [21, 23]

large numbers on both alloys, particularly on Ti6Al4V anodized in 2 M H3PO4.

thermodynamically in solutions of pH ranging to 3.

thicker oxide layer on titanium can be obtained.

**2.3 Bioactive layer** 

proper osseointegration.

SEM observations (Fig.8) and EDS microanalysis indicate the presence of deposits dispersed on the surface of anodised titanium and its implant alloys. However, deposits are are nonuniformly dispersed on a surface. Titanium and its two alloys anodised in 0.5M H3PO4 are covered with very thin oxide layer, which includes numerous and more scattered Ca-O-P deposits of diameter varied from 200 to 800 nm, suggesting the heterogeneous nucleation of Ca-O-P on TiO2 covered surface. Althogh just after 24 hours deposits are seen on the surface of the 3 materials (more deposists on the Ti6Al4V alloy) the continuous films cover the whole surfaces after 9 days in SBF solution. At higher magnification it is seen that the film on titanium is formed of more flatter layer of deposits and broken layer of titanium oxides with titanium phosphates, whereas film on both alloys comprise small globules of Ca-O-P. The ratio of Ca/P ranging from 1.26 to 1.42 corresponds to non-steochiometric hydroxyapatite.

Fig. 9. EIS) spectra recorded during immersion of titanium and Ti6Al6V samples anodized in 0.5 M and 2 M H3PO4 in SBF solution [21,23]

Impedance (EIS) spectra (Fig. 9) recorded during immersion of the anodized titanium and the Ti6Al4V alloy in simulated body fluid (SBF) for titanium and Ti6Al4V alloy show changes in capacitance and structure of surface layers as well as differences between coatings on titanium and its alloy and confirm the SEM observations (Fig. 8). Titanium exhibits two-layered structure: the inner oxide layer is covered by an outer layer of more or less uniformly distributed various size Ca-P-O deposits. On contrary the Ti6Al4V alloy is coated by a more uniform and dense layer of deposits (Fig. 8 c,d) and much lower concentration of titanium oxide on a surface [25].

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 185

Fig. 10. Transients for potentiodynamic experiments recorded for titanium anodized to 20V (with scan rate 500mV/s) at various concentration of supporting electrolyte H3PO4, a) 1M, b)

Anodic titania nanotubes formed on titanium in 1-3 M H3PO4 with 0.4% wt. HF (Fig. 11) show the morphology of nanotubes on titanium which differ in diameter and the layer

2M, c) 3M with addition of 0.4% wt. HF [53]

thickness due to electrolyte concentration.

7600F) [53,54]

Fig. 11. SEM images of titania nanotubes formed anodically prepared at 20V for 2 h in aqueous solutions of H3PO4 ranging from 1 M to 3 M with 0.4% wt. HF (field emission JEOL

SEM observations (Fig. 11) confirmed formation of a highly organized nano-sized pores, ranging from 90 to 120 nm in all applied electrolytes. As apparent, the average nanotube diameter is slightly affected by the supporting electrolyte concentration. Also, the increase

**0 0.005 0.01 0.015 0.02 0.025 0.03 0.035 0.04 0.045**

**C u rren t [A ]**

> **0 5 10 15 20 Potential [V]**

### **3. Nanostructural oxide layer formed in phosphate solutions**

In the last 20 years anodizing has been also used as a method to form nanooxides on metal surfaces. Formation of self-organized titania nanotubes with high level organization of pores on large surfaces [43-48] became very useful technology applied to many purposes, i.e. to modification of surgical implant surfaces and to biomedical sensing. Titania nanotubes, just like barrier type titania, combine very well with osseous tissue and can be a perfect basis for osteoblasts in surgical implants. Studies focused on controlling the size and arrangement of pores [49-54], aiming at bone ingrowth and on use of titania nanotubes platform for biosensing, due to their capability to combine with e.g. enzymes, proteins or biological cells, brought promising results [55,56].

Formation of nanotubes at different polarization parameters in various electrolytes [57-60] as well as various scan rates during the very first seconds of anodizing, may help tailoring the oxides for effective implantation and improve their properties for biomedical and sensor applications. For the latter applications titania nanotubes require better ordering ie. controlled diameter achieved during improved oxide growth kinetics. For the last 5 years in several papers [61-64] it has been revealed that the value of polarisation determines the diameter of nanotubes. Every additional 5V of potential increases the nanotubes diameter of about 20nm, whereas the time of anodizing determines the length of the nanotube layer. Moreover, the low pH and organic aqueous solutions assure more regular shapes of nanotubes.

#### **3.1 Nanotubes on titanium in phosphate solutions**

Electrochemical oxidization of titanium can be carried out in electrolytes with or without HF additives [65-67]. Attempts to assess the optimal scan rate/fluoride concentration ratio for formation of structurally uniform nanotubes [60] revealed that 1M H3PO4+0.3% wt. HF is the most proper electrolyte for anodizing at 0.5Vs-1. To study the effect of phosphates concentration, layers of titania nanotubes were produced in electrolytes of different phosphoric acid concentration. Their properties as the future coatings on titanium for medical uses were characterized by SEM/EDS observations and capacitance tests in simulated body fluids. Formation of oxide layers on titanium in phosphoric acid solutions with additions of fluoride ions [50] at 25ºC, is usually carried out in two stages: the first stage potentiodynamic to the desired potential and the second stage, potentiostatic with fixed potential on electrodes for over 2 hours (Fig. 10).

Fig. 10 shows the behavior of titanium polarized from the OCP (Open Circuit Potential) to 20V with a sweep rate of 0.5Vs−1 in phosphoric acid solutions of different concentration (1M, 2M and 3M H3PO4) containing 0.4 wt.% HF. Flat polarization curve confirm passive behavior of titanium anodized in 2 and 3M H3PO4+0.4 wt.% HF, contrary to current transients recorded in 1M H3PO4+0.4 wt.% HF. The increase of current with potential in that region usually can be explained by the presence of some pores [9]. Polarization curves for more concentrated phosphate solutions show the presence of the anodic peaks, which can be ascribed to the oxygen evolution [6] followed by a broad passive region. By fixing the concentration of HF (fixing the dissolution rate) the decrease of current with potential indicates that oxide formation dominates over oxide dissolution at relatively higher field strengths or/and passive layer of phosphates is formed over nanotube titania in 2 and 3M H3PO4+0.4 wt.% HF solutions.

In the last 20 years anodizing has been also used as a method to form nanooxides on metal surfaces. Formation of self-organized titania nanotubes with high level organization of pores on large surfaces [43-48] became very useful technology applied to many purposes, i.e. to modification of surgical implant surfaces and to biomedical sensing. Titania nanotubes, just like barrier type titania, combine very well with osseous tissue and can be a perfect basis for osteoblasts in surgical implants. Studies focused on controlling the size and arrangement of pores [49-54], aiming at bone ingrowth and on use of titania nanotubes platform for biosensing, due to their capability to combine with e.g. enzymes, proteins or biological cells,

Formation of nanotubes at different polarization parameters in various electrolytes [57-60] as well as various scan rates during the very first seconds of anodizing, may help tailoring the oxides for effective implantation and improve their properties for biomedical and sensor applications. For the latter applications titania nanotubes require better ordering ie. controlled diameter achieved during improved oxide growth kinetics. For the last 5 years in several papers [61-64] it has been revealed that the value of polarisation determines the diameter of nanotubes. Every additional 5V of potential increases the nanotubes diameter of about 20nm, whereas the time of anodizing determines the length of the nanotube layer. Moreover, the low

Electrochemical oxidization of titanium can be carried out in electrolytes with or without HF additives [65-67]. Attempts to assess the optimal scan rate/fluoride concentration ratio for formation of structurally uniform nanotubes [60] revealed that 1M H3PO4+0.3% wt. HF is the most proper electrolyte for anodizing at 0.5Vs-1. To study the effect of phosphates concentration, layers of titania nanotubes were produced in electrolytes of different phosphoric acid concentration. Their properties as the future coatings on titanium for medical uses were characterized by SEM/EDS observations and capacitance tests in simulated body fluids. Formation of oxide layers on titanium in phosphoric acid solutions with additions of fluoride ions [50] at 25ºC, is usually carried out in two stages: the first stage potentiodynamic to the desired potential and the second stage, potentiostatic with

Fig. 10 shows the behavior of titanium polarized from the OCP (Open Circuit Potential) to 20V with a sweep rate of 0.5Vs−1 in phosphoric acid solutions of different concentration (1M, 2M and 3M H3PO4) containing 0.4 wt.% HF. Flat polarization curve confirm passive behavior of titanium anodized in 2 and 3M H3PO4+0.4 wt.% HF, contrary to current transients recorded in 1M H3PO4+0.4 wt.% HF. The increase of current with potential in that region usually can be explained by the presence of some pores [9]. Polarization curves for more concentrated phosphate solutions show the presence of the anodic peaks, which can be ascribed to the oxygen evolution [6] followed by a broad passive region. By fixing the concentration of HF (fixing the dissolution rate) the decrease of current with potential indicates that oxide formation dominates over oxide dissolution at relatively higher field strengths or/and passive layer of phosphates is formed over nanotube titania in 2 and 3M

**3. Nanostructural oxide layer formed in phosphate solutions** 

pH and organic aqueous solutions assure more regular shapes of nanotubes.

**3.1 Nanotubes on titanium in phosphate solutions** 

fixed potential on electrodes for over 2 hours (Fig. 10).

H3PO4+0.4 wt.% HF solutions.

brought promising results [55,56].

Fig. 10. Transients for potentiodynamic experiments recorded for titanium anodized to 20V (with scan rate 500mV/s) at various concentration of supporting electrolyte H3PO4, a) 1M, b) 2M, c) 3M with addition of 0.4% wt. HF [53]

Anodic titania nanotubes formed on titanium in 1-3 M H3PO4 with 0.4% wt. HF (Fig. 11) show the morphology of nanotubes on titanium which differ in diameter and the layer thickness due to electrolyte concentration.

Fig. 11. SEM images of titania nanotubes formed anodically prepared at 20V for 2 h in aqueous solutions of H3PO4 ranging from 1 M to 3 M with 0.4% wt. HF (field emission JEOL 7600F) [53,54]

SEM observations (Fig. 11) confirmed formation of a highly organized nano-sized pores, ranging from 90 to 120 nm in all applied electrolytes. As apparent, the average nanotube diameter is slightly affected by the supporting electrolyte concentration. Also, the increase

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 187

found to exist in two forms. The binding energies of the O 1s spectra corresponded to hydroxyl groups OH- (531.3eV) and oxygen in oxides O2- (529.9 eV). The presence of fluorine in the surface layer was confirmed by the F 1s spectra of binding energy 684.6 eV associated probably with Ti. As it is shown in Fig. 12 the concentrations of titanium corresponding to titanium dioxide and Ti IV phosphates are nearly the same in all tested samples, whereas the concentrations of the other elements vary with the composition of the anodizing electrolyte. Samples anodized in 1M H3PO4+0.3% wt. HF shows the highest amount of O2- (asTiO2) and fluorides, but the lowest amount of phosphates and hydroxyl ions. It means that previously recommended [6] conditions of uniform nanotubes formation on titanium implant materials favor titanium oxidation and enhance transport of fluorides in formed titania. The similarly high concentrations of relevant elements (Ti and O), together with the highest amount of phosphates of all controlled, are observed in samples anodized in 2-3M H3PO4+0.4% wt. HF. It indicates that the use of more concentrated phosphate electrolyte leads to the increase of phosphates adsorbed over the surface layer of nanotubes in

This is well known that at anodizing a competition between oxide formation and its dissolution exists, and that HF is the key factor, which causes the production of porous oxide layer. However, there is also a competition between phosphates and fluorides in the process of oxide nanotubes formation. These 2 anions differ in size, charge and rates of

> Diffusion coefficient [mol/m2s]

Fluoride ion **2.82 10-4 147** Phosphate ion **0.88 10-9 484**

Fig. 13. Values of the OCP for nanotubes formed in 1M, 2M and 3M H3PO4 [unpublished

Table 2. Values of diffusion coefficients and Van der Waals radius for fluoride and

Van der Waals radius (pm)

competition with much smaller and more mobile fluorides.

**3.1.1 Competition between phosphates and flurides** 

diffusion in oxides (Tabl. 2)

phosphate ions

results]

of the latter from 1M to 3M under fixed HF concentration results in significant decrease of nanotube layer thickness, from 760±35 nm to 590±35 nm, respectively.


Table 1. Diameters of nanotubes and the thickness of their layer on titanium formed in 1-3 M H3PO4 with 0.4% wt. HF

The XPS analysis (Fig. 12) revealed that the highest amount of fluorides in oxide surface layer was obtained in 1M H3PO4+0.3% wt. HF, but in this case the lowest amount of phosphates adsorbed above nanotubes was observed. Using higher concentrations of phosphoric acid 2-3 M H3PO4 Judging on the results of XPS analysis, the competition between fluorides and phosphates is observed during anodizing and the higher concentration of the latter is responsible for higher bioactivity of nanotubes formed in 2M H3PO4+0.4% wt. HF [53].

Fig. 12. Results of XPS analysis of nanotubes formed on titanium in 1-3 M H3PO4 solution containing different amount of fluorides, from 0.2% wt to 0.4% wt. HF VSW (Vacuum Systems Workshop, Ltd.) Kα Al (1486.6 eV) X-ray radiation working at power of 210 W (15 kV - voltage, 14 mA – emission current) [53].

The XPS spectra (Ti 2p, O 1s, P 2p and F 1s) revealed that nanotube layers consist of Ti, O, F and P species. The Ti 2p spectra for all samples showed only one doublet line. The position of the Ti 2p3/2 peak on the binding energy scale at 458.8 eV corresponded to titanium dioxide and Ti IV phosphates [16]. One type of phosphorus was revealed by the P 2p3/2 peak position at 133.3 eV associated to phosphate type species, indicating that species from the electrolyte are indeed adsorbed over the oxide film during anodizing. Oxygen was

of the latter from 1M to 3M under fixed HF concentration results in significant decrease of

Diameter (+/- 10nm)

1M 100nm 750nm 2M **110nm** 700nm 3M 120nm 590nm Table 1. Diameters of nanotubes and the thickness of their layer on titanium formed in 1-3 M

The XPS analysis (Fig. 12) revealed that the highest amount of fluorides in oxide surface layer was obtained in 1M H3PO4+0.3% wt. HF, but in this case the lowest amount of phosphates adsorbed above nanotubes was observed. Using higher concentrations of phosphoric acid 2-3 M H3PO4 Judging on the results of XPS analysis, the competition between fluorides and phosphates is observed during anodizing and the higher concentration of the latter is

responsible for higher bioactivity of nanotubes formed in 2M H3PO4+0.4% wt. HF [53].

Fig. 12. Results of XPS analysis of nanotubes formed on titanium in 1-3 M H3PO4 solution containing different amount of fluorides, from 0.2% wt to 0.4% wt. HF VSW (Vacuum Systems Workshop, Ltd.) Kα Al (1486.6 eV) X-ray radiation working at power of 210 W (15

The XPS spectra (Ti 2p, O 1s, P 2p and F 1s) revealed that nanotube layers consist of Ti, O, F and P species. The Ti 2p spectra for all samples showed only one doublet line. The position of the Ti 2p3/2 peak on the binding energy scale at 458.8 eV corresponded to titanium dioxide and Ti IV phosphates [16]. One type of phosphorus was revealed by the P 2p3/2 peak position at 133.3 eV associated to phosphate type species, indicating that species from the electrolyte are indeed adsorbed over the oxide film during anodizing. Oxygen was

kV - voltage, 14 mA – emission current) [53].

Thickness (+/- 35nm)

nanotube layer thickness, from 760±35 nm to 590±35 nm, respectively.

H3PO4 Concentration

H3PO4 with 0.4% wt. HF

found to exist in two forms. The binding energies of the O 1s spectra corresponded to hydroxyl groups OH- (531.3eV) and oxygen in oxides O2- (529.9 eV). The presence of fluorine in the surface layer was confirmed by the F 1s spectra of binding energy 684.6 eV associated probably with Ti. As it is shown in Fig. 12 the concentrations of titanium corresponding to titanium dioxide and Ti IV phosphates are nearly the same in all tested samples, whereas the concentrations of the other elements vary with the composition of the anodizing electrolyte. Samples anodized in 1M H3PO4+0.3% wt. HF shows the highest amount of O2- (asTiO2) and fluorides, but the lowest amount of phosphates and hydroxyl ions. It means that previously recommended [6] conditions of uniform nanotubes formation on titanium implant materials favor titanium oxidation and enhance transport of fluorides in formed titania. The similarly high concentrations of relevant elements (Ti and O), together with the highest amount of phosphates of all controlled, are observed in samples anodized in 2-3M H3PO4+0.4% wt. HF. It indicates that the use of more concentrated phosphate electrolyte leads to the increase of phosphates adsorbed over the surface layer of nanotubes in competition with much smaller and more mobile fluorides.

#### **3.1.1 Competition between phosphates and flurides**

This is well known that at anodizing a competition between oxide formation and its dissolution exists, and that HF is the key factor, which causes the production of porous oxide layer. However, there is also a competition between phosphates and fluorides in the process of oxide nanotubes formation. These 2 anions differ in size, charge and rates of diffusion in oxides (Tabl. 2)


Table 2. Values of diffusion coefficients and Van der Waals radius for fluoride and phosphate ions

Fig. 13. Values of the OCP for nanotubes formed in 1M, 2M and 3M H3PO4 [unpublished results]

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 189

phosphate concentration in anodizing electrolyte, accounting for the processes at the nanotube layer/electrolyte interface which can be associated with deposited products. As the changes between spectra occurred during the first hour of exposure to SBF solution, one can assume that the deposition processes on nanotube layers formed in 2-3M H3PO4+0.4%

The formation of porous metal oxides, ie. titania and alumina, is explained by a fieldenhanced model [11,61,63,69] that depends on the ability of ions to diffuse through the metal oxide. Thus, due to the large size the incorporation of the phosphate ions is difficult, but the increased fluoride concentration in solution leads to its ability to migrate and intercalate into the oxide films during the anodizing [70]. The XPS results show (Fig. 12), that the increased fluoride concentration is accompanied by the decreased phosphates and hydroxyl ions in adsorbed layer over nanotubes [53]. It correlates very well with the results of titanium anodizing in electrolyte not containing fluorides, where the gel-like protective

Titanium and its implant alloys, mainly ternary alloys of Ti-6Al-7Nb or, are widely used in biomedical implants and dental fields due to their unique mechanical, chemical properties,

Further improvement of the unique properties of nanotube anodic layers for medical applications, particularly for enhancement of bone in-growth [74] and biosensing [75] require not only the development of the formation method on two phase titanium alloys, but also providing the proper morphology and structure. Reported efforts to form anodic nanotube layers on Ti alloys such as Ti-6Al-7Nb, TiAl [76], or Ti45Nb [77] showed the formation of highly inhomogeneous surfaces due to selective dissolution of the less stable

Studies on development of nanotubes growth on the Ti6A4V [60] and Ti-6Al-7Nb alloys [54] were focused on varying the HF concentrations in the phosphoric acid media, in order to establish the pore size distribution and estimate the critical scan rate/concentration ratio for the initiation of nanopitting in compact oxide layer, which would be decisive for the

Among several parameters influencing the quality of nanotubes formed anodically, such as potential, time of anodizing, fluoride ions concentration and scan rate of polarization, particularly the last two seem to be determiners for nanotubes structure and morphology. As an example to show the effect of fluoride ions concentration on the morphology of nanotubes on the implant alloy, the anodizing of the two phase (α+β) Ti6Al7Nb alloy samples in 1 M H3PO4 containing 0.2%; 0.3% and 0.4 % wt. HF to 20V using scan rate 500mV/s and then holding them at that potential for further 2h in the same electrolyte, was performed. Nanotubes of diameter ranging from 50nm to 80nm, with thicker walls over βphase grains than over α- phase grains, were obtained. During the formation process, which includes two stages: the first potentiodynamic and the second potentiostatic (20V), different electrochemical behaviour was observed in electrolytes of various fluoride concentration.

wt. HF are quick.

layer of phosphates was formed over the oxide [42].

**3.2 Nanotubes on implant alloys in phosphate solutions** 

excellent corrosion resistance and biocompatibility [71-73].

formation of uniform nanotubes on both two-phase alloys.

phase and/or different reaction rates of the different phases of the alloys.

The effect of the phosphoric acid concentration in fluoride containing electrolytes on the properties of oxide nanotubes formed at anodizing was characterized by SEM/EDS observations and capacitance characteristics when immersed in simulated body fluids (SBF) in order to predict their behavior as the future coatings on titanium for biomaterial applications [53].

The values of the OCP (Fig. 13) for nanotubes formed in 1M, 2M and 3M H3PO4, each containing 0.4% wt. HF, measured at 25ºC in SBF solution 1 h after anodizing are -0.140 V, - 0.170V and -0.195V (SCE), respectively, indicate the observed earlier [68] decrease of the OCP of oxide produced in more concentrated phosphoric acid solution.

The impedance spectra for titanium anodized in 1-3M H3PO4+0.4% wt. HF were obtained at the OCP for frequency ranging from 105 to 0.18 Hz with ac amplitude 10 mV. The spectra recorded 1 hour after immersion in SBF solution show that variations in chemical composition of the surface layer over obtained nanotubes are confirmed by variations in capacitance characteristics.

The results of EIS tests (Fig.14) indicate nearly the same properties (similar impedance and – angle values) for ohmic resistance of the electrolyte and its penetration through nanotube films. However in the low frequency range, the impedance values are sensitive to the

Fig. 14. EIS spectra (a- Nyquist, b,c- Bode spectra) for titania nanotubes formed on titanium in 1-3 M H3PO4 solutions with addition of 0.4% wt. HF recorded at the OCP in SBF solution at 25ºC [53]

The effect of the phosphoric acid concentration in fluoride containing electrolytes on the properties of oxide nanotubes formed at anodizing was characterized by SEM/EDS observations and capacitance characteristics when immersed in simulated body fluids (SBF) in order to predict their behavior as the future coatings on titanium for biomaterial

The values of the OCP (Fig. 13) for nanotubes formed in 1M, 2M and 3M H3PO4, each containing 0.4% wt. HF, measured at 25ºC in SBF solution 1 h after anodizing are -0.140 V, - 0.170V and -0.195V (SCE), respectively, indicate the observed earlier [68] decrease of the

The impedance spectra for titanium anodized in 1-3M H3PO4+0.4% wt. HF were obtained at the OCP for frequency ranging from 105 to 0.18 Hz with ac amplitude 10 mV. The spectra recorded 1 hour after immersion in SBF solution show that variations in chemical composition of the surface layer over obtained nanotubes are confirmed by variations in

The results of EIS tests (Fig.14) indicate nearly the same properties (similar impedance and – angle values) for ohmic resistance of the electrolyte and its penetration through nanotube

Fig. 14. EIS spectra (a- Nyquist, b,c- Bode spectra) for titania nanotubes formed on titanium in 1-3 M H3PO4 solutions with addition of 0.4% wt. HF recorded at the OCP in SBF solution

a

b. c

films. However in the low frequency range, the impedance values are sensitive to the

OCP of oxide produced in more concentrated phosphoric acid solution.

applications [53].

capacitance characteristics.

at 25ºC [53]

phosphate concentration in anodizing electrolyte, accounting for the processes at the nanotube layer/electrolyte interface which can be associated with deposited products. As the changes between spectra occurred during the first hour of exposure to SBF solution, one can assume that the deposition processes on nanotube layers formed in 2-3M H3PO4+0.4% wt. HF are quick.

The formation of porous metal oxides, ie. titania and alumina, is explained by a fieldenhanced model [11,61,63,69] that depends on the ability of ions to diffuse through the metal oxide. Thus, due to the large size the incorporation of the phosphate ions is difficult, but the increased fluoride concentration in solution leads to its ability to migrate and intercalate into the oxide films during the anodizing [70]. The XPS results show (Fig. 12), that the increased fluoride concentration is accompanied by the decreased phosphates and hydroxyl ions in adsorbed layer over nanotubes [53]. It correlates very well with the results of titanium anodizing in electrolyte not containing fluorides, where the gel-like protective layer of phosphates was formed over the oxide [42].

#### **3.2 Nanotubes on implant alloys in phosphate solutions**

Titanium and its implant alloys, mainly ternary alloys of Ti-6Al-7Nb or, are widely used in biomedical implants and dental fields due to their unique mechanical, chemical properties, excellent corrosion resistance and biocompatibility [71-73].

Further improvement of the unique properties of nanotube anodic layers for medical applications, particularly for enhancement of bone in-growth [74] and biosensing [75] require not only the development of the formation method on two phase titanium alloys, but also providing the proper morphology and structure. Reported efforts to form anodic nanotube layers on Ti alloys such as Ti-6Al-7Nb, TiAl [76], or Ti45Nb [77] showed the formation of highly inhomogeneous surfaces due to selective dissolution of the less stable phase and/or different reaction rates of the different phases of the alloys.

Studies on development of nanotubes growth on the Ti6A4V [60] and Ti-6Al-7Nb alloys [54] were focused on varying the HF concentrations in the phosphoric acid media, in order to establish the pore size distribution and estimate the critical scan rate/concentration ratio for the initiation of nanopitting in compact oxide layer, which would be decisive for the formation of uniform nanotubes on both two-phase alloys.

Among several parameters influencing the quality of nanotubes formed anodically, such as potential, time of anodizing, fluoride ions concentration and scan rate of polarization, particularly the last two seem to be determiners for nanotubes structure and morphology. As an example to show the effect of fluoride ions concentration on the morphology of nanotubes on the implant alloy, the anodizing of the two phase (α+β) Ti6Al7Nb alloy samples in 1 M H3PO4 containing 0.2%; 0.3% and 0.4 % wt. HF to 20V using scan rate 500mV/s and then holding them at that potential for further 2h in the same electrolyte, was performed. Nanotubes of diameter ranging from 50nm to 80nm, with thicker walls over βphase grains than over α- phase grains, were obtained. During the formation process, which includes two stages: the first potentiodynamic and the second potentiostatic (20V), different electrochemical behaviour was observed in electrolytes of various fluoride concentration.

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 191

The typical current transients Fig. 16 recorded during the anodizing of the Ti-6Al-7Nb alloy in 1M H3PO4 containing 0.3 wt.% HF are similar to current transients observed during nanotube oxide layers formation on other alloys in other electrolytes [78]. As in previously described process the whole treatment consists of the potentiodynamic polarization from the OCP to 20V with a scan rate of 0.5Vs−1, followed by the potentiostatic polarization at 20V for further 2 h. However, contrary to constant current density increase observed at anodizing of pure Ti [49,50,79], during the potentiodynamic sweep to 20V at the alloy Ti-6Al-7Nb anodizing the current transients show 2 peaks: the first at about 2-3 V due to oxygen evolution and the second at about 4-6V linked probably to Al oxidation. In the potentiostatic stage of anodizing the current density for the alloy decreases until the end of the treatment, while in case of Ti a broad peak is seen at about 900 s of the anodizing (Fig. 16). According to [78,79] the broad peak, typically recorded in the potentiostatic stage of the process, indicates the dissolution of oxide before reaching final balance between both processes: oxide formation and oxide dissolution during nanotubes formation. Such the balance determines a steady-state oxide layer formation stage during anodizing of metals

Fig. 17. SEM images of nanotubes produced on the Ti6Al7Nb (a ,c, e α-phase; b, d, f- phase) by anodization at 20V for 2h in 1M H3PO4 containing (a),(b) 0.2%HF, (c),(d) 0.3% HF,

Small pits on β phase grains and regular nanotubes on α-phase are observed in 0.2 wt.% HF (Fig. 17a,b). Irregular tubes on β-phase and regular tubes on α-phase grains are seen after anodising in 0.3 wt. % HF (Fig. 17c,d). Both phases are covered with regular nanotubes in case of samples anodised in 0.4 wt. % HF (Fig. 17e,f), but on β- phase nanotube walls are thicker than on α- phase. Fig. 16 illustrates the dissolved oxide over α-phase on Ti6Al4V alloy and bigger size of nanotubes over β- phase in more concentrated phosphoric acid

[11].

(e),(f) 0.4% wt. HF

solution.

The implant alloy Ti6Al7Nb (Fig 15) of black α phase (hcp) and white β phase (bcc) irregular shape platelets forming variously oriented colonies, with the surface fraction of α and β phases 78% and 22%, respectively, was enriched with aluminum in oxides over α phase and enriched with niobium over β phase anodic nanotubes.

Fig. 15. Microstructure of the Ti-6Al-7Nb alloy [54]

Fig. 16. Current transient for potentiodymic and potentiostatic stages recorded at anodizing the Ti-6Al-7Nb alloy and titanium (for comparison) at 20V for 2h in 1M H3PO4 containing 0.3% HF (scan rate in the potentiodynamic stage 500mV/s) and current transients recorded during potentiodynamic stage of anodizing of the Ti6Al7Nb alloy in 1 M H3PO4 with different fluoride concentration, a) 0.2%HF b) 0.3%HF c) 0.4%HF [54]

The implant alloy Ti6Al7Nb (Fig 15) of black α phase (hcp) and white β phase (bcc) irregular shape platelets forming variously oriented colonies, with the surface fraction of α and β phases 78% and 22%, respectively, was enriched with aluminum in oxides over α phase and

enriched with niobium over β phase anodic nanotubes.

Fig. 15. Microstructure of the Ti-6Al-7Nb alloy [54]

**0 0.001 0.002 0.003 0.004 0.005 0.006 0.007 0.008 0.009 0.01**

**Current [A]**

**0 5 10 15 20 Potential [V]**

Fig. 16. Current transient for potentiodymic and potentiostatic stages recorded at anodizing the Ti-6Al-7Nb alloy and titanium (for comparison) at 20V for 2h in 1M H3PO4 containing 0.3% HF (scan rate in the potentiodynamic stage 500mV/s) and current transients recorded during potentiodynamic stage of anodizing of the Ti6Al7Nb alloy in 1 M H3PO4 with

different fluoride concentration, a) 0.2%HF b) 0.3%HF c) 0.4%HF [54]

The typical current transients Fig. 16 recorded during the anodizing of the Ti-6Al-7Nb alloy in 1M H3PO4 containing 0.3 wt.% HF are similar to current transients observed during nanotube oxide layers formation on other alloys in other electrolytes [78]. As in previously described process the whole treatment consists of the potentiodynamic polarization from the OCP to 20V with a scan rate of 0.5Vs−1, followed by the potentiostatic polarization at 20V for further 2 h. However, contrary to constant current density increase observed at anodizing of pure Ti [49,50,79], during the potentiodynamic sweep to 20V at the alloy Ti-6Al-7Nb anodizing the current transients show 2 peaks: the first at about 2-3 V due to oxygen evolution and the second at about 4-6V linked probably to Al oxidation. In the potentiostatic stage of anodizing the current density for the alloy decreases until the end of the treatment, while in case of Ti a broad peak is seen at about 900 s of the anodizing (Fig. 16). According to [78,79] the broad peak, typically recorded in the potentiostatic stage of the process, indicates the dissolution of oxide before reaching final balance between both processes: oxide formation and oxide dissolution during nanotubes formation. Such the balance determines a steady-state oxide layer formation stage during anodizing of metals [11].

Fig. 17. SEM images of nanotubes produced on the Ti6Al7Nb (a ,c, e α-phase; b, d, f- phase) by anodization at 20V for 2h in 1M H3PO4 containing (a),(b) 0.2%HF, (c),(d) 0.3% HF, (e),(f) 0.4% wt. HF

Small pits on β phase grains and regular nanotubes on α-phase are observed in 0.2 wt.% HF (Fig. 17a,b). Irregular tubes on β-phase and regular tubes on α-phase grains are seen after anodising in 0.3 wt. % HF (Fig. 17c,d). Both phases are covered with regular nanotubes in case of samples anodised in 0.4 wt. % HF (Fig. 17e,f), but on β- phase nanotube walls are thicker than on α- phase. Fig. 16 illustrates the dissolved oxide over α-phase on Ti6Al4V alloy and bigger size of nanotubes over β- phase in more concentrated phosphoric acid solution.

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 193

increases with fluoride concentration, but seems to reach the limit in these conditions for 0.3% wt. HF (Fig. 19). The highest current density (Fig. 14) is linked to the biggest nanotube diameters, as it was observed in case of pure titanium anodised in the same conditions [50].

> 0.3%HF [weight %]

Phases α β α β α β [weight %] Titanium 63.90 51.66 61.39 35.70 59.52 38.59 69.22 Oxygen 32.09 36.66 34.44 45.43 36.40 44.99 19.72 Aluminium 4.00 2.91 4.17 2.69 4.08 1.99 4.45 Niobium --- 8.77 --- 16.18 --- 14.43 6.53

Table 3. Results of EDS analysis of nanotube layers obtained by anodizing at 20V for 2h

transform into Nb2O5 oxide at 20V, according to the equations 6-7 [83,84]:

biological media immobilization which require morphologically uniform surface.

**4. Titania layers formed in phosphate solutions for biosensing** 

Due to chemical similarity of titanium and niobium [11, 82] electrochemical behaviour of the Ti-6Al-7Nb electrode should be qualitatively similar to that of the titanium and niobium electrodes in the potential range from -1 to 4V (SCE). Electrochemical oxidation of niobium electrode leads to formation of sub-oxides NbO and NbO2 at the OCP, which partly

 2NbO2 + H2O – 2e- → Nb2O5 + 2H+ (8) The dissolution process of niobium oxide (β-phase) (5) increases with increasing fluoride concentration [85], so the fluoride concentration is a crucial factor for nanotubes growth on Ti-6Al-7Nb. Structural and metallurgical aspects of the formation of self-organized anodic oxide nanotube layers on alloys are crucial for medical application to the advanced techniques of

The additional advantageous property of phosphate rich compact and nanotubular anodic oxide layers on titanium is its ability to attach enzymes, proteins or biological cells. To test such the possibility in order to apply anodic surfaces for H2O2 biosensing two electrodes were prepared: 1) the first electrode prepared by the electropolymerization of conducting polymer (PANI) on the surface of Ti/TiO2 (compact) electrode [55], 2) the second electrode was prepared by using titania nanotubes on titanium as a platform of the 3rd generation biosensor [56]. In both cases the HRP (horseradish peroxide) enzyme was immobilized on the sensing surface. By using either cyclic voltammetry or amperometric modes the feasibility and electrochemical parameters for H2O2 monitoring on Ti/TiO2 surface were checked in the simulated body fluid (SBF). Both electrodes were sensitive to H2O2, however the second electrode only in the presence of thionine as the mediator [56]. Two peaks seen on cyclic voltammograms (Fig. 18) for the Ti/TiO2 (nanotube) electrode with immobilized HRP, indicate the sensitivity of the prepared platform to the presence of H2O2 in the analyte.

0.4%HF [weight %]

Nb + H2O – 2e- → NbO + 2H+ (6)

NbO + H2O – 2e- → NbO2 + 2H+ (7)

Compact oxide

Fluoride concentration

0.2%HF [weight %]

in 1M H3PO4 containing 0.2; 0.3; 0.4% wt. HF [53]

Fig. 18. SEM images of nanotubes produced on the Ti6Al4V by anodization at 20V for 2h in 1M H3PO4 containing (a) 0.2%HF, (b) 0.4% wt. HF

According to EDS analysis nanotubes formed on the Ti-6Al-7Nb alloy showed that those films are predominately TiO2 with small amounts of Ti2O3, Al or Nb oxides (Table 3). Aluminium and niobium are present in their most stable oxidation states, Al2O3 and Nb2O3. The amount of alloying elements in the nanotube oxide layer was influenced by the underlying metal microstructure, where Nb was present in the β- phase and Al in the αphase [80].

Fig. 19. Results of XPS examination of surface layer of nanotubes formed on the Ti6Al7Nb alloy in 1 M H3PO4 with different fluoride concentration

In the combined SEM and XPS examinations [54] (Fig. 17 and 19) the highest intensities for all controlled elements and groups: titanium oxide and titanium phosphates (458.7eV), oxides (530eV), hydroxyl ions (531.6eV), phosphates (133.3eV) and fluorides (648.6eV), clearly confirm that the most advantageous scan rate and electrolyte composition for the formation of uniform nanotube layer on the Ti6Al7Nb alloy, are 0.5Vs-1 during potentiodynamic stage of anodizing in 1M H3PO4 containing 0.3% wt. HF. Interesting is that also the intensity of niobium (207.3eV) in the most stable of the niobium oxides Nb2O5 [81],

 Fig. 18. SEM images of nanotubes produced on the Ti6Al4V by anodization at 20V for 2h in

According to EDS analysis nanotubes formed on the Ti-6Al-7Nb alloy showed that those films are predominately TiO2 with small amounts of Ti2O3, Al or Nb oxides (Table 3). Aluminium and niobium are present in their most stable oxidation states, Al2O3 and Nb2O3. The amount of alloying elements in the nanotube oxide layer was influenced by the underlying metal microstructure, where Nb was present in the β- phase and Al in the α-

 Fig. 19. Results of XPS examination of surface layer of nanotubes formed on the Ti6Al7Nb

In the combined SEM and XPS examinations [54] (Fig. 17 and 19) the highest intensities for all controlled elements and groups: titanium oxide and titanium phosphates (458.7eV), oxides (530eV), hydroxyl ions (531.6eV), phosphates (133.3eV) and fluorides (648.6eV), clearly confirm that the most advantageous scan rate and electrolyte composition for the formation of uniform nanotube layer on the Ti6Al7Nb alloy, are 0.5Vs-1 during potentiodynamic stage of anodizing in 1M H3PO4 containing 0.3% wt. HF. Interesting is that also the intensity of niobium (207.3eV) in the most stable of the niobium oxides Nb2O5 [81],

1M H3PO4 containing (a) 0.2%HF, (b) 0.4% wt. HF

alloy in 1 M H3PO4 with different fluoride concentration

phase [80].


increases with fluoride concentration, but seems to reach the limit in these conditions for 0.3% wt. HF (Fig. 19). The highest current density (Fig. 14) is linked to the biggest nanotube diameters, as it was observed in case of pure titanium anodised in the same conditions [50].

Table 3. Results of EDS analysis of nanotube layers obtained by anodizing at 20V for 2h in 1M H3PO4 containing 0.2; 0.3; 0.4% wt. HF [53]

Due to chemical similarity of titanium and niobium [11, 82] electrochemical behaviour of the Ti-6Al-7Nb electrode should be qualitatively similar to that of the titanium and niobium electrodes in the potential range from -1 to 4V (SCE). Electrochemical oxidation of niobium electrode leads to formation of sub-oxides NbO and NbO2 at the OCP, which partly transform into Nb2O5 oxide at 20V, according to the equations 6-7 [83,84]:

$$\text{Nb} + \text{H}\_2\text{O} - 2\text{e}^\cdot \rightarrow \text{NbO} + 2\text{H}^+ \tag{6}$$

$$\text{NbO} + \text{H}\_2\text{O} - 2\text{e} \to \text{NbO}\_2 + 2\text{H}^\* \tag{7}$$

$$2\text{NbO}\_2 + \text{H}\_2\text{O} - 2\text{e}^\cdot \rightarrow \text{Nb}\_2\text{O}\_5 + 2\text{H}^\* \tag{8}$$

The dissolution process of niobium oxide (β-phase) (5) increases with increasing fluoride concentration [85], so the fluoride concentration is a crucial factor for nanotubes growth on Ti-6Al-7Nb. Structural and metallurgical aspects of the formation of self-organized anodic oxide nanotube layers on alloys are crucial for medical application to the advanced techniques of biological media immobilization which require morphologically uniform surface.

#### **4. Titania layers formed in phosphate solutions for biosensing**

The additional advantageous property of phosphate rich compact and nanotubular anodic oxide layers on titanium is its ability to attach enzymes, proteins or biological cells. To test such the possibility in order to apply anodic surfaces for H2O2 biosensing two electrodes were prepared: 1) the first electrode prepared by the electropolymerization of conducting polymer (PANI) on the surface of Ti/TiO2 (compact) electrode [55], 2) the second electrode was prepared by using titania nanotubes on titanium as a platform of the 3rd generation biosensor [56]. In both cases the HRP (horseradish peroxide) enzyme was immobilized on the sensing surface. By using either cyclic voltammetry or amperometric modes the feasibility and electrochemical parameters for H2O2 monitoring on Ti/TiO2 surface were checked in the simulated body fluid (SBF). Both electrodes were sensitive to H2O2, however the second electrode only in the presence of thionine as the mediator [56]. Two peaks seen on cyclic voltammograms (Fig. 18) for the Ti/TiO2 (nanotube) electrode with immobilized HRP, indicate the sensitivity of the prepared platform to the presence of H2O2 in the analyte.

Anodic Layer Formation on Titanium and Its Alloys for Biomedical Applications 195

of higher concentrations of phosphoric acid (2-3M H3PO4 with 0.4% wt. HF assures the formation of nanotubes containing the high concentration of both bioactivity enhancing elements, fluorides and adsorbed phosphates. The obtained titania nanotubes show the significantly higher bioactivity in vitro during the first hour of immersion in SBF in

Depending on fluoride ion concentrations in anodizing electrolyte morphologically different nanotubular layers have been obtained on both phases of two titanium alloys: the Ti6Al4V and the Ti6Al7Nb alloys. Self organized nanotubes grow on both phases ( and in 1M H3PO4 containing 0.4% wt. HF, though smaller pore size and thicker wall tubes are obtained on the β phase. The electrochemical behavior of both phases of the alloys differs due to fluoride concentrations which is the key parameter in controlling their morphology. Uniform nanotubes are obtained in 2M H3PO4 containing 0.3% wt. HF at scan rate of polarization 0.5Vs-1 during potentiodynamic stage of anodizing. Such conditions assure the highest fluoride and phosphate concentrations in surface layer of nanotubes on titanium and nanotubes containing niobium oxide on the Ti6Al7Nb alloy. Both features promise a proper coating for improved osteoblast cell adhesion on artificial implants and for

Funding from the Polish Ministry of Science and Higher Education under the N507 082 31/2009 project and by the National Centre of Research and Development under ERA-NET/MNT/TNTBIOSENS/1/2011 Project is gratefully acknowledged. My thanks go also to Professor Patrick Schmuki of University of Erlangen for inspiration on nanotubes and to my

[1] Trasatti S., Lodi G.: Electrodes of Conductive Metallic Oxides. ed. S. Trasatti, Chapt. B,

[3] Marciniak J.: Biomateriały (in Polish) Technical University of Upper Silesia Press (2002),

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#### **5. Conclusions**

Titanium surfaces can be modified by electrochemical treatment in the phosphoric acid solutions for better corrosion resistance, improved physicochemical and electrochemical properties and bioactivity. The formation of oxide layers enriched with phosphorus of 30- 120nm thick can be formed in 0.5 M H3PO4 at both galvanostatic anodizing current density values varied in the range of 0.1-0.5 Am-2 and potentiostatically at up to 60V giving yellowish layers porous on the surface. Due to the presence of phosphates they are highly bioactive in comparison to oxides formed in other electrolytes and are covered with hydroxyapatite deposists after 9 days in the SBF solution.

Anodizing in more concentrated 2M solutions of the phosphoric acid generates a gel-like film with thickness of about 100nm on titanium. The film, containing a large number of phosphates nuclei, exhibits its effectiveness to activate titanium surface for biomimetic coating of calcium phosphate. The electrochemically treated titanium was able to form uniform Ca–P coatings on titanium after 48 hour immersions in the SBF solution. The treatment is a simple method to generate bioactive metal surfaces, besides other methods such as alkaline treatment applied to titanium implant materials.

Electrochemical treatment in the phosphoric acid solutions with the addition of 0.2-0.4% wt. HF allows to form on titanium and its implant alloys nano-sized pores (nanotubes) in more concentrated phosphoric acid solutions (1-3M). Their morphology, electrochemical properties and chemical composition are in close relation with the anodic polarization parameters and with the concentration of both ions: phosphates and fluorides. The highest amount of fluorides in surface layer is obtained when using 1M H3PO4+0.3% wt. HF, but in this case the lowest amount of phosphates adsorbed above nanotubes is observed. The use of higher concentrations of phosphoric acid (2-3M H3PO4 with 0.4% wt. HF assures the formation of nanotubes containing the high concentration of both bioactivity enhancing elements, fluorides and adsorbed phosphates. The obtained titania nanotubes show the significantly higher bioactivity in vitro during the first hour of immersion in SBF in comparison to barrier titanium oxide.

Depending on fluoride ion concentrations in anodizing electrolyte morphologically different nanotubular layers have been obtained on both phases of two titanium alloys: the Ti6Al4V and the Ti6Al7Nb alloys. Self organized nanotubes grow on both phases ( and in 1M H3PO4 containing 0.4% wt. HF, though smaller pore size and thicker wall tubes are obtained on the β phase. The electrochemical behavior of both phases of the alloys differs due to fluoride concentrations which is the key parameter in controlling their morphology. Uniform nanotubes are obtained in 2M H3PO4 containing 0.3% wt. HF at scan rate of polarization 0.5Vs-1 during potentiodynamic stage of anodizing. Such conditions assure the highest fluoride and phosphate concentrations in surface layer of nanotubes on titanium and nanotubes containing niobium oxide on the Ti6Al7Nb alloy. Both features promise a proper coating for improved osteoblast cell adhesion on artificial implants and for biosensing.

#### **6. Acknowledgment**

194 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Fig. 20. SEM of Ti/TiO2 (nanotube) covered with HRP and cyclic voltammograms for Ti/TiO2/HRP in 0.1 M PBS (pH 6.8) in the presence of H2O2, scan rate 100 mV/s [56]

hydroxyapatite deposists after 9 days in the SBF solution.

such as alkaline treatment applied to titanium implant materials.

Titanium surfaces can be modified by electrochemical treatment in the phosphoric acid solutions for better corrosion resistance, improved physicochemical and electrochemical properties and bioactivity. The formation of oxide layers enriched with phosphorus of 30- 120nm thick can be formed in 0.5 M H3PO4 at both galvanostatic anodizing current density values varied in the range of 0.1-0.5 Am-2 and potentiostatically at up to 60V giving yellowish layers porous on the surface. Due to the presence of phosphates they are highly bioactive in comparison to oxides formed in other electrolytes and are covered with

Anodizing in more concentrated 2M solutions of the phosphoric acid generates a gel-like film with thickness of about 100nm on titanium. The film, containing a large number of phosphates nuclei, exhibits its effectiveness to activate titanium surface for biomimetic coating of calcium phosphate. The electrochemically treated titanium was able to form uniform Ca–P coatings on titanium after 48 hour immersions in the SBF solution. The treatment is a simple method to generate bioactive metal surfaces, besides other methods

Electrochemical treatment in the phosphoric acid solutions with the addition of 0.2-0.4% wt. HF allows to form on titanium and its implant alloys nano-sized pores (nanotubes) in more concentrated phosphoric acid solutions (1-3M). Their morphology, electrochemical properties and chemical composition are in close relation with the anodic polarization parameters and with the concentration of both ions: phosphates and fluorides. The highest amount of fluorides in surface layer is obtained when using 1M H3PO4+0.3% wt. HF, but in this case the lowest amount of phosphates adsorbed above nanotubes is observed. The use

**5. Conclusions** 

Funding from the Polish Ministry of Science and Higher Education under the N507 082 31/2009 project and by the National Centre of Research and Development under ERA-NET/MNT/TNTBIOSENS/1/2011 Project is gratefully acknowledged. My thanks go also to Professor Patrick Schmuki of University of Erlangen for inspiration on nanotubes and to my coworkers from the Biomedical Engineering Division at University of Zielona Góra.

#### **7. References**


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**9** 

*Malaysia* 

**Surface Modification** 

**Techniques for Biomedical Grade of** 

**and Ion Implantation Processes** 

*1Universiti Teknologi Malaysia, UTM Skudai, Johor,* 

E.M. Nazim1, R. Rosliza2, A. Shah1 and M.A. Hassan1

*2TATi University College, Malaysia, Kemaman, Terengganu,* 

**Titanium Alloys: Oxidation, Carburization** 

S. Izman1, Mohammed Rafiq Abdul-Kadir1, Mahmood Anwar1,

Titanium and titanium alloys are widely used in a variety of engineering applications, where the combination of mechanical and chemical properties is of crucial importance. Aerospace, chemical and automotive industries as well as the medical device manufacturers also benefited from the outstanding properties of titanium alloys. The wide spread of its uses in biomedical implants is mainly due to their well-established corrosion resistance and biocompatibility. However, not all titanium and its alloys can meet all of the clinical requirements for biomedical implants. For instance, it is reported that bare titaniumvanadium alloy has traces of vanadium ion release after long period exposure with body fluid (López et al., 2010). Excessive metal ions release into the body fluid and causes toxicity problems to the host body. A new group of titanium alloy such as Ti-Nb and Ti-Zr based are recently introduced in the market to overcome the toxicity of titanium-vanadium based alloy (Gutiérrez *et al.*, 2008). Although, these alloys have a high strength to weight ratio and good corrosion resistance and biocompatible, but it suffers from poor tribological properties which limits their usefulness to a certain extent especially when they are applied to joint movements. Wear debris generated from these articulation joints can induce inflammation problem and toxic effect to the human body. In biomedical point of view, post implantation is very crucial stage where the interaction between the implanted material surface and the biological environment in human body is critically evaluated. Either in the short or long run, the toxic effect becomes an issue to host body. Hence, the implant material surface has a strong role in the responses to the biological environment. In order to improve the biological and tribological properties of implant materials, surface modification is often required (Huang *et al.*, 2006, Kumar *et al.*, 2010b). This chapter embarks on the commonly used implant biomaterials, followed by general overview on the surface modification techniques for treating titanium alloy. The basic principles of oxidation, carburization and ion implantation methods and their developments are discussed in the following sections.

**1. Introduction** 


### **Surface Modification Techniques for Biomedical Grade of Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes**

S. Izman1, Mohammed Rafiq Abdul-Kadir1, Mahmood Anwar1, E.M. Nazim1, R. Rosliza2, A. Shah1 and M.A. Hassan1 *1Universiti Teknologi Malaysia, UTM Skudai, Johor, 2TATi University College, Malaysia, Kemaman, Terengganu, Malaysia* 

#### **1. Introduction**

200 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

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niobium oxide films, Electrochemica Acta, 1975, 20, 237-244

acid solutions, Thin Solid Films, 1997; 307, 141-147.

Titanium and titanium alloys are widely used in a variety of engineering applications, where the combination of mechanical and chemical properties is of crucial importance. Aerospace, chemical and automotive industries as well as the medical device manufacturers also benefited from the outstanding properties of titanium alloys. The wide spread of its uses in biomedical implants is mainly due to their well-established corrosion resistance and biocompatibility. However, not all titanium and its alloys can meet all of the clinical requirements for biomedical implants. For instance, it is reported that bare titaniumvanadium alloy has traces of vanadium ion release after long period exposure with body fluid (López et al., 2010). Excessive metal ions release into the body fluid and causes toxicity problems to the host body. A new group of titanium alloy such as Ti-Nb and Ti-Zr based are recently introduced in the market to overcome the toxicity of titanium-vanadium based alloy (Gutiérrez *et al.*, 2008). Although, these alloys have a high strength to weight ratio and good corrosion resistance and biocompatible, but it suffers from poor tribological properties which limits their usefulness to a certain extent especially when they are applied to joint movements. Wear debris generated from these articulation joints can induce inflammation problem and toxic effect to the human body. In biomedical point of view, post implantation is very crucial stage where the interaction between the implanted material surface and the biological environment in human body is critically evaluated. Either in the short or long run, the toxic effect becomes an issue to host body. Hence, the implant material surface has a strong role in the responses to the biological environment. In order to improve the biological and tribological properties of implant materials, surface modification is often required (Huang *et al.*, 2006, Kumar *et al.*, 2010b). This chapter embarks on the commonly used implant biomaterials, followed by general overview on the surface modification techniques for treating titanium alloy. The basic principles of oxidation, carburization and ion implantation methods and their developments are discussed in the following sections.

Surface Modification Techniques for Biomedical Grade of

various metallic biomaterials used for medical applications.

or trade name

Cast CoCrMo

Grade 2 "Phynox" Grade 1 "Elgiloy"

Ti-3Al-2.5W(α/β) Tikrutan (α/β) Ti 364(α/β) Ti 364ELI(α/β) Ti 367(α/β)

Ti-15Mo (Metastable Beta) "TMZF" (Metastable Beta) "Beta 3" (Metastable Beta) Ti-15Mo-5Zr-3Al(Metastable Beta) Ti-13Nb-13Zr (Metastable Beta) "TiOsteum" (Metastable Beta) Ti-45Nb (Metastable Beta)

Table 1. Metallic biomaterials for medical and surgical implants (Courtesy from ATI, Allvac,

Wrought CoCrMo, Alloy 1 Wrought CoCrMo, Alloy 2 Wrought CoCrMo,GADS

REX 734 XM-19 108

L-605 Syncoben

35N

CP-1(α) CP-2(α)

316L Stainless Steel 316L Stainless Steel ASTM Standard

ASTM F 138 ASTM F 745 ASTM F 1586 ASTM F 1314 F-04.12.35

ASTM F 75 ASTM F 1537 ASTM F 1537 ASTM F 1537 ASTM F 90 ASTM F 563 ASTM F 1058 ASTM F 1058 ASTM F 562

ASTM F 67 ASTM F 67 ASTM F 2146

ASTM F 1472 ASTM F 136 ASTM F 1295 ASTM F 2066 ASTM F 1813

ASTM F 1713

ASTM B 348




ISO Standard







ISO 5832-1

ISO 5832-9

ISO 5832-4 ISO 5832-12 ISO 5832-12

ISO 5832-5 ISO 5832-8 ISO 5832-7 ISO 5832-7 ISO 5832-6

ISO 5832-2 ISO 5832-2

ISO 5832-10 ISO 5832-3 ISO 5832-3 ISO 5832-11

ISO 5832-14

Material designation Common

Steel biomaterial : Fe-18Cr-14Ni-2.5Mo Fe-18Cr-12.5Ni-2.5Mo,Cast Fe-21Cr-10Ni-3.5Mn-2.5Mo Fe-22Cr-12.5Ni-5Mn-2.5Mo Fe-23Mn -21Cr-1Mo-1N

Cobalt base biomaterials: Co-28Cr-6Mo Casting alloy Co-28Cr-6Mo Wrought alloy#1 Co-28Cr-6Mo Wrought alloy#2 Co-28Cr-6Mo Wrought alloy#3 Co-20Cr-15W-10Ni-1.5Mn Co-20Ni-20Cr-5Fe-3.5Mo-3.5W-2Ti Co-19Cr-7Ni -14Fe-7Mo-1.5W-2Ti Co-20Cr-15Ni -15Fe-7Mo-2Mn

Co-35Ni-20Cr-10Mo

Ti-12Mo-6Zr-2Fe Ti-11.5Mo-6Zr-4.5Sn Ti-15Mo-5Zr-3Al Ti-13Nb-13Zr Ti-35Nb-7Zr-5Ta

Ti-45Nb

USA)

Ti Cp-1 Ti CP-2 Ti-3Al-2.5V Ti-5Al-2.5Fe Ti-6Al-4V Ti-6Al-4V ELI Ti-6Al-7Nb Ti-15Mo

Titanium base biomaterials:

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 203

other hand, biomedical grade 316L stainless steel possesses higher biocompatibility properties and it can be used as implant as well as for surgical devices. However, due to its heavy weight, this group of materials have been gradually replaced by other lighter biomaterials such as titanium alloys. Co based alloys is another metallic biomaterials but its high elastic modulus compared to bone causes stress shielding effect especially in load bearing applications (Kumar *et al.*, 2010b). Stress shielding effect is the reduction of bone density due to the removal of normal stress from bone by an implant. Among metallic materials, titanium and its alloys are considered as the most convincing materials in medical applications nowadays because they exhibit superior corrosion resistance and tissue acceptance when compared to stainless steels and Co-based alloys. Table 1 shows the

#### **2. Implant biomaterials**

Biomedical implant is defined as an artificial organ used for restoring the functionality of a damaged natural organ or tissue of the body (Liu *et al.*, 2004). In other words it is expected to perform the functions of the natural organ or tissue without adverse effect to other body parts (Andrew *et al.*, 2004). Fig.1 shows a typical hip and knee joint implants replacement. Biomedical engineering is a new discipline where engineering principles and design concepts are applied to improve healthcare diagnosis, monitoring and therapy by solving medical and biological science related problems. At least three different terminologies that always are being referred to biomedical implant materials, i.e. biocompatibility, biodegradable and biomaterials. Biocompatibility is defined as the immune rejection or inflammatory responses of the surrounding tissue systems to the presence of a foreign object in the body. Whereas biodegradable material means the implant material can easily decompose in the body. Their presence in the body is temporary and usually they degrade as a function of time, temperature or pressure. Biomaterials must possess biocompatibility and sometimes biodegradable properties. A typical example of simultaneously possess biocompatibility and biodegradable properties is drug delivery capsule where their presence to release drugs inside the body over a specific time without causing any toxic effect to the surrounding tissues (Hollinger, 2006). There are varieties of biomaterials such as metallic, ceramics and polymers that have been used as biomedical devices. Metallic biomaterials can be grouped as steels, cobalt and titanium based alloys. Among non-metallic polymeric based biomaterials are polyethylene terephthalate, polytetrafluoroethylene, ultrahigh molecular weight polyethylene (UHMWPE) and lactide-co-glycolide. While titania (TiO2), titanium carbide (TiC), titanium nitride (TiC), bioglass, hydroxyapatite (HA), silicon carbide (SiC) are typical examples of ceramic biomaterials. All biomaterials must be free from cytotoxicity. Plain steel would corrode easily in the body and become toxic. On the

Fig. 1. Total hip and knee implants replacements (Geetha, 2009)

Biomedical implant is defined as an artificial organ used for restoring the functionality of a damaged natural organ or tissue of the body (Liu *et al.*, 2004). In other words it is expected to perform the functions of the natural organ or tissue without adverse effect to other body parts (Andrew *et al.*, 2004). Fig.1 shows a typical hip and knee joint implants replacement. Biomedical engineering is a new discipline where engineering principles and design concepts are applied to improve healthcare diagnosis, monitoring and therapy by solving medical and biological science related problems. At least three different terminologies that always are being referred to biomedical implant materials, i.e. biocompatibility, biodegradable and biomaterials. Biocompatibility is defined as the immune rejection or inflammatory responses of the surrounding tissue systems to the presence of a foreign object in the body. Whereas biodegradable material means the implant material can easily decompose in the body. Their presence in the body is temporary and usually they degrade as a function of time, temperature or pressure. Biomaterials must possess biocompatibility and sometimes biodegradable properties. A typical example of simultaneously possess biocompatibility and biodegradable properties is drug delivery capsule where their presence to release drugs inside the body over a specific time without causing any toxic effect to the surrounding tissues (Hollinger, 2006). There are varieties of biomaterials such as metallic, ceramics and polymers that have been used as biomedical devices. Metallic biomaterials can be grouped as steels, cobalt and titanium based alloys. Among non-metallic polymeric based biomaterials are polyethylene terephthalate, polytetrafluoroethylene, ultrahigh molecular weight polyethylene (UHMWPE) and lactide-co-glycolide. While titania (TiO2), titanium carbide (TiC), titanium nitride (TiC), bioglass, hydroxyapatite (HA), silicon carbide (SiC) are typical examples of ceramic biomaterials. All biomaterials must be free from cytotoxicity. Plain steel would corrode easily in the body and become toxic. On the

Fig. 1. Total hip and knee implants replacements (Geetha, 2009)

**2. Implant biomaterials** 

other hand, biomedical grade 316L stainless steel possesses higher biocompatibility properties and it can be used as implant as well as for surgical devices. However, due to its heavy weight, this group of materials have been gradually replaced by other lighter biomaterials such as titanium alloys. Co based alloys is another metallic biomaterials but its high elastic modulus compared to bone causes stress shielding effect especially in load bearing applications (Kumar *et al.*, 2010b). Stress shielding effect is the reduction of bone density due to the removal of normal stress from bone by an implant. Among metallic materials, titanium and its alloys are considered as the most convincing materials in medical applications nowadays because they exhibit superior corrosion resistance and tissue acceptance when compared to stainless steels and Co-based alloys. Table 1 shows the various metallic biomaterials used for medical applications.


Table 1. Metallic biomaterials for medical and surgical implants (Courtesy from ATI, Allvac, USA)

Surface Modification Techniques for Biomedical Grade of

**Biochemical methods** Modification through silanized

as

**CVD (Chemical Vapour Deposition)** 

**Physical methods**  Thermal spray Flame spray HVOF DGUN

**PVD (Physical Vapour** 

**Ion implantation and** 

**Glow discharge plasma** 

**Deposition)**  Evaporation Ion plating Sputtering

**deposition**  Beam-line ion implantation

PIII

**treatment** 

(Liu *et al.*, 2004)

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 205

Improve wear resistance, corrosion resistance and blood compatibility

Induce specific cell and tissue response by means of surfaceimmobilized peptides, proteins, or growth factors

Improve wear resistance,

resistance and biological

Improve wear resistance, corrosion resistance and blood compatibility

Modify surface composition; improve wear, corrosion resistance, and biocompatibility

Clean, sterilize, oxide, nitride surface; remove native oxide layer

corrosion

properties

~1 µm of TiN, TiC, TiCN, diamond and diamond-like carbon thin film

titania, photochemistry, selfassembled monolayers, proteinresistance, etc.

~30 to 200 µm of coatings, such

titanium, HA, calcium silicate,

~1 µm of TiN, TiC, TiCN, diamond and diamond-like carbon thin film

~10 nm of surface modified layer and/or ~µm of thin film

~1 nm to ~100 nm of surface

Table 2. Summary of surface modification methods used for titanium and its alloys implants

Among the popular surface modification methods are mechanical, chemical, sol-gel, oxidation, carburization and ion implantation. The last three methods will be discussed in detailed in later sections. Mechanical surface treatments include machining, grinding, and blasting. These methods were discussed in depth elsewhere (Lausmaa *et al.*, 2001). The main goal of mechanical modification is to obtain particular surface roughness and topographies on implant surface. In general, mechanical surface treatments lead to rough structures which finally increase the surface area of implant. This condition is considered more favourable for the implant because it facilitates biomineralization process to take place (Sobieszczyk, 2010a). Surface roughness enhances cell attachment, proliferation and differentiation of osteogenic cells and is the key factor for the osseous integration of metallic implants. Among the mechanical methods, blasting is the most popular technique for achieving desired surface roughness on titanium. The common abrasive particles used as the blasting media are silicon carbide (SiC), alumina (Al2O3), biphasic calcium phosphates (BCP), hydroxyapatite and ß-Tricalcium phosphate (Citeau *et al.*, 2005). One of the disadvantages of blasting is that it may lead to surface contamination and local inflammatory reactions of surrounding tissues as a result of dissolution of abrasive particles

modified layer

Al2O3, ZrO2, TiO2

### **3. Overview of surface modification techniques**

Since all biomedical devices subject to extremely high clinical requirements, a thorough surface modification process is needed prior to implantation process into the human body. The main reasons to carry out various surface modification processes on implant materials for biomedical applications can be summarized as follows:


The proper surface modification techniques keep the excellent bulk attributes of titanium alloys, such as good fatigue strength, formability, machinability and relatively low modulus. It also improves specific surface properties required by different clinical requirements. Table 2 summarizes the typical surface modification schemes used to treat titanium and its alloys for implant.


Since all biomedical devices subject to extremely high clinical requirements, a thorough surface modification process is needed prior to implantation process into the human body. The main reasons to carry out various surface modification processes on implant materials

iii. Increase hardness of implant to reduce wear rate especially in articulation joint

The proper surface modification techniques keep the excellent bulk attributes of titanium alloys, such as good fatigue strength, formability, machinability and relatively low modulus. It also improves specific surface properties required by different clinical requirements. Table 2 summarizes the typical surface modification schemes used to treat titanium and its alloys

**Modified layer Objective** 

Produce specific surface topographies; clean and roughen surface; improve adhesion in bonding

Remove oxide scales and contamination Improve biocompatibility, bioactivity or bone conductivity

Improving biocompatibility, bioactivity or bone conductivity

Improve biocompatibility, bioactivity or bone conductivity

Produce specific surface topographies; improved corrosion resistance; improve biocompatibility, bioactivity or bone conductivity

i. Clean implant material surface from contaminations prior to implantation ii. Increase bioactivity, cell growth and tissue attachments after implantation

iv. Introduce passive layer to prevent excessive ion release into body environment

Rough or smooth surface formed by subtraction process

<10 nm of surface oxide layer

~1 µm of sodium titanate gel

~5 nm of dense inner oxide and porous outer layer

calcium phosphate, TiO2 and silica

adsorption and incorporation of electrolyte anions

**Sol–gel** ~10 µm of thin film, such as

**Anodic oxidation** ~10 nm to 40 µm of TiO2 layer,

**3. Overview of surface modification techniques** 

for biomedical applications can be summarized as follows:

applications

for implant.

v. Promote antibacterial effect

**Surface modification methods** 

**Mechanical methods** Machining Grinding Polishing Blasting

**Chemical methods** Chemical treatment Acidic

> treatment Alkaline treatment

 Hydrogen peroxide treatment

vi. Increase fatigue strength of implants


Table 2. Summary of surface modification methods used for titanium and its alloys implants (Liu *et al.*, 2004)

Among the popular surface modification methods are mechanical, chemical, sol-gel, oxidation, carburization and ion implantation. The last three methods will be discussed in detailed in later sections. Mechanical surface treatments include machining, grinding, and blasting. These methods were discussed in depth elsewhere (Lausmaa *et al.*, 2001). The main goal of mechanical modification is to obtain particular surface roughness and topographies on implant surface. In general, mechanical surface treatments lead to rough structures which finally increase the surface area of implant. This condition is considered more favourable for the implant because it facilitates biomineralization process to take place (Sobieszczyk, 2010a). Surface roughness enhances cell attachment, proliferation and differentiation of osteogenic cells and is the key factor for the osseous integration of metallic implants. Among the mechanical methods, blasting is the most popular technique for achieving desired surface roughness on titanium. The common abrasive particles used as the blasting media are silicon carbide (SiC), alumina (Al2O3), biphasic calcium phosphates (BCP), hydroxyapatite and ß-Tricalcium phosphate (Citeau *et al.*, 2005). One of the disadvantages of blasting is that it may lead to surface contamination and local inflammatory reactions of surrounding tissues as a result of dissolution of abrasive particles

Surface Modification Techniques for Biomedical Grade of

effective and less complicated equipment.

and oxygen gas (O2) the can be written as:

**4. Oxidation** 

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 207

Sol-gel consists of two terms, sol and gel. A sol can be defined as a colloidial suspension of very small solid particles in a continuous liquid. Gel can be defined as a substance that contains a continuous solid skeleton enveloping a continuous liquid phase (Brinker, 1990). The sol–gel process consists of five main steps: (1) hydrolysis and polycondensation; (2) gelation; (3) aging; (4) drying; (5) densification and crystallization (Piveteau *et al.*, 2001). Two different techniques usually used to carry out the sol–gel process: (i) spin coating technique and (ii) dip coating technique. In spin coating technique the specimens are spun to spread the coating solution on the substrate using centrifugal force where in dip coating specimens are dipped or submerged in the solution. The sol–gel process is popular for depositing thin (<10 µm) ceramic coatings (Liu *et al.*, 2004). In the biomedical area, the sol–gel process is considered new field. Sol-gel method capable of producing various types of coatings on titanium and titanium alloys for biomedical applications. Examples of these coatings are titanium oxide (TiO2), calcium phosphate (CaP), and TiO2–CaP composite. Sol–gel technique also has been applied for some silica-based coatings. It has a great potential to replace plasma spray for synthesizing the composite hydroxyapatite/titania coating on the titanium substrate with high adhesion and good bioactivity (Kim *et al.*, 2004). It is reported that plasma spray method results in chemical inhomogeneity and low crystallinity of HA coating on titanium alloys (Wang *et al.*, 2011). In contrast, sol-gel technique produces high crystalline HA microstructure and better chemical homogeneity due to ability to mix the calcium and phosphorus precursors at molecular-level. They also found that atomic diffusion accelerated when increasing the calcining temperature or prolong the calcining time. Other advantages of sol-gel method in comparison with other conventional thin layer oxidation processes are: i) low densification temperature, ii) better control of the homogeneity, chemical composition and crystalline structure of the thin coating, iii) Cost

Oxidation is a chemical reaction between metal and oxygen. This reaction occurs naturally. However, this reaction can be started with exciting the atoms by providing external energy. In simple way, an oxidation is defined as a chemical reaction by the interaction of metal with oxygen to form an oxide. The oxidation behaviour of a metal depends on various factors and the reaction mechanism usually quite complex. The phenomena started with adsorption of oxygen molecules from the atmosphere, and then followed by nucleation of oxides, formation of a thin oxide layer, finally growth to a thicker scale. During the growth process, nodule formation and scale spallation may also take place (Khanna, 2004) . The total chemical reaction for the formation of oxide (MaOb) by oxidation between metal (M)

The mechanism of oxidation process is illustrated in Fig. 2. The initial step started by the adsorption of gas on the clean metal surface during the metal-oxygen reaction. As the reaction proceeds, oxygen may dissolve in the metal forming an oxide on the surface either as a film or as oxide nuclei. The gas adsorption and initial oxide formation both are functions of various factors: (i) surface orientation, (ii) crystal defects at the surface, (iii) surface preparation, and (iv) impurities in both the metal and the gas. The oxides formed on

aM + (b/2) O2 = MaOb (1)

into the host bone (Gbureck *et al.*, 2003). A range of surface roughness (Ra = 0.5-1.5 μm) shows stronger bone response after the implantation compared to implants with smoother or rougher surface (Sobieszczyk, 2010b). This observation was in contrast with the findings by Fini *et al.* (2003), that rougher surface show encouraging results. Their results were confirmed in the vivo experiments using titanium implants having roughness of 16.5 – 21.4 μm inserted in the cortical and trabecular bone of goats.

Chemical methods include acid treatment, alkali treatment, sol–gel, oxidation, chemical vapour deposition (CVD), and biochemical modification. Following discussion is limited to the first four chemical methods. Since oxidation method itself is a big field, this technique is separately discussed in section 4.

Acid treatment is a popular surface treatment method to clean substrate surface by means of removing oxide and contamination. A mixed acids solution is frequently used for this purpose (Nanci *et al.*, 1998). It is also noted that TiO2 is the dominant oxide layers formed on the substrate due to high affinity of titanium to react with O2. These oxides need to be removed prior to other surface treatments such as HA coating, thermal oxidation or carburization and ion implantation. A recommended standard solution for acid treatment is composed of HNO3 and HF (ratio of 10 to 1 by volume) in distilled water. Hydrofluoric acid has natural tendency of quickly attack TiO2 in the acid solution and forms soluble titanium fluorides and hydrogen. This acid solution also can be used to minimize the formation of free hydrogen that prevent surface embrittlement occurs due to inclusion of hydrogen in titanium (American Society for Testing and Materials, 1997). A group of researchers investigated the decontamination efficiency to the Ti surface using three acids, Na2S2O8, H2SO4, and HCl (Takeuchi *et al.*, 2003). They found that HCl was the most effective decontamination agent among these three due to the capability to dissolve titanium salts easily without weakening Ti surfaces.

Alkali treatment can be simply defined as simple surface modification by alkali solution such as NaOH or KOH to form bioactive porous layer on substrate materials. Later, this method followed by thermal treatment to dehydrate and transform amorphous structure into porous crystalline. The combined treatment is called Alkali Heat Treatment (AHT). The alkali treatment process is started by immersing titanium alloy in a 5–10 M NaOH or KOH solution for 24 hr (Kim *et al.*, 1996). After that specimens have to rinse with distilled water followed by ultrasonic cleaning. It is then dried in an oven. Finally heat treatment is carried out by heating the specimens around 600–800 oC for 1 hr. The heat treatment is performed at very low pressure for avoiding oxidation of titanium at high temperature. The porous surface formed on treated titanium surface disclosed the formation of sodium titanate hydrogel on the titanium substrate. It was observed that after thermal treatment, a large quantity of crystalline sodium titanate with rutile and anatase precipitated. Bioactive bonelike apatite was obtained on the surface after soaking the treated titanium in simulated body fluid (SBF) for 4 weeks. They found that bone-like apatite layer which is bioactive can be formed on other surfaces such as bioglass, hydroxyapatite and glass–ceramic by using this method. It is noted that bioglass, hydroxyapatite and glass–ceramic all are the examples of bioactive ceramics. Recently, a group of researchers investigated the effect of substrate surface roughness on alkali treated CP-Ti for apatite formation after immerging in SBF solution for seven days (Ravelingien *et al.*, 2010). They found that apatite formation increased with the moderate surface roughness. However, very smooth surface (< 0.5 µm) causes sudden decrease in apatite formation.

Sol-gel consists of two terms, sol and gel. A sol can be defined as a colloidial suspension of very small solid particles in a continuous liquid. Gel can be defined as a substance that contains a continuous solid skeleton enveloping a continuous liquid phase (Brinker, 1990). The sol–gel process consists of five main steps: (1) hydrolysis and polycondensation; (2) gelation; (3) aging; (4) drying; (5) densification and crystallization (Piveteau *et al.*, 2001). Two different techniques usually used to carry out the sol–gel process: (i) spin coating technique and (ii) dip coating technique. In spin coating technique the specimens are spun to spread the coating solution on the substrate using centrifugal force where in dip coating specimens are dipped or submerged in the solution. The sol–gel process is popular for depositing thin (<10 µm) ceramic coatings (Liu *et al.*, 2004). In the biomedical area, the sol–gel process is considered new field. Sol-gel method capable of producing various types of coatings on titanium and titanium alloys for biomedical applications. Examples of these coatings are titanium oxide (TiO2), calcium phosphate (CaP), and TiO2–CaP composite. Sol–gel technique also has been applied for some silica-based coatings. It has a great potential to replace plasma spray for synthesizing the composite hydroxyapatite/titania coating on the titanium substrate with high adhesion and good bioactivity (Kim *et al.*, 2004). It is reported that plasma spray method results in chemical inhomogeneity and low crystallinity of HA coating on titanium alloys (Wang *et al.*, 2011). In contrast, sol-gel technique produces high crystalline HA microstructure and better chemical homogeneity due to ability to mix the calcium and phosphorus precursors at molecular-level. They also found that atomic diffusion accelerated when increasing the calcining temperature or prolong the calcining time. Other advantages of sol-gel method in comparison with other conventional thin layer oxidation processes are: i) low densification temperature, ii) better control of the homogeneity, chemical composition and crystalline structure of the thin coating, iii) Cost effective and less complicated equipment.

#### **4. Oxidation**

206 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

into the host bone (Gbureck *et al.*, 2003). A range of surface roughness (Ra = 0.5-1.5 μm) shows stronger bone response after the implantation compared to implants with smoother or rougher surface (Sobieszczyk, 2010b). This observation was in contrast with the findings by Fini *et al.* (2003), that rougher surface show encouraging results. Their results were confirmed in the vivo experiments using titanium implants having roughness of 16.5 – 21.4

Chemical methods include acid treatment, alkali treatment, sol–gel, oxidation, chemical vapour deposition (CVD), and biochemical modification. Following discussion is limited to the first four chemical methods. Since oxidation method itself is a big field, this technique is

Acid treatment is a popular surface treatment method to clean substrate surface by means of removing oxide and contamination. A mixed acids solution is frequently used for this purpose (Nanci *et al.*, 1998). It is also noted that TiO2 is the dominant oxide layers formed on the substrate due to high affinity of titanium to react with O2. These oxides need to be removed prior to other surface treatments such as HA coating, thermal oxidation or carburization and ion implantation. A recommended standard solution for acid treatment is composed of HNO3 and HF (ratio of 10 to 1 by volume) in distilled water. Hydrofluoric acid has natural tendency of quickly attack TiO2 in the acid solution and forms soluble titanium fluorides and hydrogen. This acid solution also can be used to minimize the formation of free hydrogen that prevent surface embrittlement occurs due to inclusion of hydrogen in titanium (American Society for Testing and Materials, 1997). A group of researchers investigated the decontamination efficiency to the Ti surface using three acids, Na2S2O8, H2SO4, and HCl (Takeuchi *et al.*, 2003). They found that HCl was the most effective decontamination agent among these three due to the capability to dissolve titanium salts

Alkali treatment can be simply defined as simple surface modification by alkali solution such as NaOH or KOH to form bioactive porous layer on substrate materials. Later, this method followed by thermal treatment to dehydrate and transform amorphous structure into porous crystalline. The combined treatment is called Alkali Heat Treatment (AHT). The alkali treatment process is started by immersing titanium alloy in a 5–10 M NaOH or KOH solution for 24 hr (Kim *et al.*, 1996). After that specimens have to rinse with distilled water followed by ultrasonic cleaning. It is then dried in an oven. Finally heat treatment is carried out by heating the specimens around 600–800 oC for 1 hr. The heat treatment is performed at very low pressure for avoiding oxidation of titanium at high temperature. The porous surface formed on treated titanium surface disclosed the formation of sodium titanate hydrogel on the titanium substrate. It was observed that after thermal treatment, a large quantity of crystalline sodium titanate with rutile and anatase precipitated. Bioactive bonelike apatite was obtained on the surface after soaking the treated titanium in simulated body fluid (SBF) for 4 weeks. They found that bone-like apatite layer which is bioactive can be formed on other surfaces such as bioglass, hydroxyapatite and glass–ceramic by using this method. It is noted that bioglass, hydroxyapatite and glass–ceramic all are the examples of bioactive ceramics. Recently, a group of researchers investigated the effect of substrate surface roughness on alkali treated CP-Ti for apatite formation after immerging in SBF solution for seven days (Ravelingien *et al.*, 2010). They found that apatite formation increased with the moderate surface roughness. However, very smooth surface (< 0.5 µm)

μm inserted in the cortical and trabecular bone of goats.

separately discussed in section 4.

easily without weakening Ti surfaces.

causes sudden decrease in apatite formation.

Oxidation is a chemical reaction between metal and oxygen. This reaction occurs naturally. However, this reaction can be started with exciting the atoms by providing external energy. In simple way, an oxidation is defined as a chemical reaction by the interaction of metal with oxygen to form an oxide. The oxidation behaviour of a metal depends on various factors and the reaction mechanism usually quite complex. The phenomena started with adsorption of oxygen molecules from the atmosphere, and then followed by nucleation of oxides, formation of a thin oxide layer, finally growth to a thicker scale. During the growth process, nodule formation and scale spallation may also take place (Khanna, 2004) . The total chemical reaction for the formation of oxide (MaOb) by oxidation between metal (M) and oxygen gas (O2) the can be written as:

$$\mathbf{a}\mathbf{M} + \begin{pmatrix} \mathbf{b}/2 \end{pmatrix} \mathbf{O}\_2 = \mathbf{M}\_\mathbf{a} \mathbf{O}\_\mathbf{b} \tag{1}$$

The mechanism of oxidation process is illustrated in Fig. 2. The initial step started by the adsorption of gas on the clean metal surface during the metal-oxygen reaction. As the reaction proceeds, oxygen may dissolve in the metal forming an oxide on the surface either as a film or as oxide nuclei. The gas adsorption and initial oxide formation both are functions of various factors: (i) surface orientation, (ii) crystal defects at the surface, (iii) surface preparation, and (iv) impurities in both the metal and the gas. The oxides formed on

Surface Modification Techniques for Biomedical Grade of

and αM (element activity in alloy) = γM . XM ; γM = the activity coefficient of metal in the alloy ;

XM = mole fraction of metal in the alloy; P(O2) = Partial pressure of the oxygen gas.

oxide, the equation 3 becomes

**4.2 Thermal oxidation** 

structures (Krishna *et al.*, 2005).

αMO2 and αM = Activities of the oxide and the metal respectiely

Where,

2004).

ΔGo = Gibbs free energy R = Universal gas constant T= Absolute temperature

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 209

If a value coefficient γM is not available, ideal behaviour is assumed and γM is assigned the value of unity. Assuming element activity of the solid constituents is unit, i.e. the metal and

Or, P (O2) = exp (ΔGo / RT) (5) Therefore, equation 5 can be used to determine the partial pressure of oxygen required for any metal to form oxide at any temperature from the standard free formation energy. Standard free energy for the formation of oxides is a function of temperature. This can be obtained from the Ellingham/Richardson diagrams which is mentioned elsewhere (Khanna,

Thermal oxidation occurs when metals or alloys are heated in a highly oxidizing atmosphere such as air or in the presence of oxygen. It is one of the cost- effective surface modification methods to deliberately generate a barrier oxide layer on titanium alloy. Thermal oxidation treatment aims for obtaining a ceramic coating, mainly focussed on rutile structure. Particularly, oxidation at temperature above 200 °C promotes the development of a crystalline oxide film. Many researchers reported that the thermally formed oxide layer enables increment in hardness, wear resistance and corrosion resistance of titanium and its alloy (Borgioli *et al.*, 2005, Kumar *et al.*, 2009). This protective oxide layer also reduce ion release inside body fluid and thus helps the body from metal toxicity (López *et al.*, 2010). During thermal oxidation process, titanium can easily reacts with air due to its affinity to oxygen. Three types of oxides structure can be produced through this method, which are rutile, brookite or anatase structures. Among the three, rutile structure is more preferable for several reasons. Rutile structure is more inert to bacterial attack (Bloyce *et al.*, 1998), has high hardness and low friction coefficient that can reduce wear as compared to the other two

Many researchers investigate thermal oxidation method to solve ion release issues through increasing the corrosion resistance of titanium alloy. A group of researchers studied the chemical composition of oxide layer produced by thermal oxidation on vanadium free TiNb and TiZr based alloys (López *et al.*, 2001). Their aim was to reduce ion release as well as improving corrosion resistance for better biocompatibility. They reported that Ti, Al and Zr based oxide dominate the surface where small amount of Nb based oxide formed. They also

ΔGo = - RT ln P (O2) (4)

surface separates the metal and the gas and sometimes act as a barrier for further oxide formation. This barrier oxide is called protective oxide layer. The oxide can be continuous film or porous structure. Oxides can also be liquid or volatile at high temperature. In general, the reaction mechanism for a specific metal will be a function of several factors: (i) pre-treatment and surface preparation of the metal, (ii) temperature, (iii) gas composition, (iv) pressure and (v) required time of reaction (Kofstad, 1988). The oxidation mechanism can be generalised both at room temperature as well as at high temperature. The basic difference between oxidation at room temperature and high temperature is the reaction rate. At room temperature reaction rate is very slow where at high temperature the rate is accelerated. There are various types of oxidation for surface modification of biomedical grade titanium alloy such as (i) Thermal oxidation, (ii) Anodic oxidation, (iii) Micro-arc oxidation (MAO). These techniques are discussed separately in 4.2, 4.3 and 4.4 respectively.

Fig. 2. Scale Formation during high temperature metal oxidation: (a) O2 gas absorption, (b) O2 dissolution, (c) Thin oxide film formation, (d) Oxide layer growth, (e) Thick oxide layer (Kofstad, 1988)

#### **4.1 Mechanism of oxidation based on thermodynamic point of view**

In oxidation, the chemical reaction between a metal (M) and the oxygen gas (O2) can be written as:

$$\text{M (s)} + \text{O}\_2\text{ (g)} = \text{MO}\_2\text{ (s)}\tag{2}$$

In thermodynamic point of view, if oxygen potential in the environment is greater than the oxygen partial pressure in equilibrium with the oxide then an oxide will form on the surface of that metal. This equilibrium oxygen pressure is determined from the standard free energy of formation of the oxide. This equilibrium oxygen pressure is also called the dissociation pressure of the oxide in equilibrium with the metal. From equation 2, the standard free energy of the oxidation can be written as,

$$
\Delta \text{G}^{\circ} = -\text{RT} \ln \left( \mathfrak{a}\_{\text{MCO2}} / \mathfrak{a}\_{\text{M}} \cdot \mathrm{P}\_{\text{(O2)}} \right) \tag{3}
$$

Where,

208 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

surface separates the metal and the gas and sometimes act as a barrier for further oxide formation. This barrier oxide is called protective oxide layer. The oxide can be continuous film or porous structure. Oxides can also be liquid or volatile at high temperature. In general, the reaction mechanism for a specific metal will be a function of several factors: (i) pre-treatment and surface preparation of the metal, (ii) temperature, (iii) gas composition, (iv) pressure and (v) required time of reaction (Kofstad, 1988). The oxidation mechanism can be generalised both at room temperature as well as at high temperature. The basic difference between oxidation at room temperature and high temperature is the reaction rate. At room temperature reaction rate is very slow where at high temperature the rate is accelerated. There are various types of oxidation for surface modification of biomedical grade titanium alloy such as (i) Thermal oxidation, (ii) Anodic oxidation, (iii) Micro-arc oxidation (MAO). These techniques are discussed separately in

Fig. 2. Scale Formation during high temperature metal oxidation: (a) O2 gas absorption, (b) O2 dissolution, (c) Thin oxide film formation, (d) Oxide layer growth, (e) Thick oxide layer

In oxidation, the chemical reaction between a metal (M) and the oxygen gas (O2) can be

In thermodynamic point of view, if oxygen potential in the environment is greater than the oxygen partial pressure in equilibrium with the oxide then an oxide will form on the surface of that metal. This equilibrium oxygen pressure is determined from the standard free energy of formation of the oxide. This equilibrium oxygen pressure is also called the dissociation pressure of the oxide in equilibrium with the metal. From equation 2, the standard free

M (s) + O2 (g) = MO2 (s) (2)

(d)

(e)

ΔGo = - RT ln ( αMO2 / αM . P(O2)) (3)

**4.1 Mechanism of oxidation based on thermodynamic point of view** 

(c)

(b)

(a)

4.2, 4.3 and 4.4 respectively.

(Kofstad, 1988)

written as:

energy of the oxidation can be written as,

ΔGo = Gibbs free energy R = Universal gas constant T= Absolute temperature αMO2 and αM = Activities of the oxide and the metal respectiely and αM (element activity in alloy) = γM . XM ; γM = the activity coefficient of metal in the alloy ; XM = mole fraction of metal in the alloy; P(O2) = Partial pressure of the oxygen gas.

If a value coefficient γM is not available, ideal behaviour is assumed and γM is assigned the value of unity. Assuming element activity of the solid constituents is unit, i.e. the metal and oxide, the equation 3 becomes

$$
\Delta \mathbf{G}^{\diamond} = \text{-RT } \ln \mathbf{P}\_{\text{(O2)}} \tag{4}
$$

Or, P (O2) = exp (ΔGo / RT) (5)

Therefore, equation 5 can be used to determine the partial pressure of oxygen required for any metal to form oxide at any temperature from the standard free formation energy. Standard free energy for the formation of oxides is a function of temperature. This can be obtained from the Ellingham/Richardson diagrams which is mentioned elsewhere (Khanna, 2004).

#### **4.2 Thermal oxidation**

Thermal oxidation occurs when metals or alloys are heated in a highly oxidizing atmosphere such as air or in the presence of oxygen. It is one of the cost- effective surface modification methods to deliberately generate a barrier oxide layer on titanium alloy. Thermal oxidation treatment aims for obtaining a ceramic coating, mainly focussed on rutile structure. Particularly, oxidation at temperature above 200 °C promotes the development of a crystalline oxide film. Many researchers reported that the thermally formed oxide layer enables increment in hardness, wear resistance and corrosion resistance of titanium and its alloy (Borgioli *et al.*, 2005, Kumar *et al.*, 2009). This protective oxide layer also reduce ion release inside body fluid and thus helps the body from metal toxicity (López *et al.*, 2010). During thermal oxidation process, titanium can easily reacts with air due to its affinity to oxygen. Three types of oxides structure can be produced through this method, which are rutile, brookite or anatase structures. Among the three, rutile structure is more preferable for several reasons. Rutile structure is more inert to bacterial attack (Bloyce *et al.*, 1998), has high hardness and low friction coefficient that can reduce wear as compared to the other two structures (Krishna *et al.*, 2005).

Many researchers investigate thermal oxidation method to solve ion release issues through increasing the corrosion resistance of titanium alloy. A group of researchers studied the chemical composition of oxide layer produced by thermal oxidation on vanadium free TiNb and TiZr based alloys (López *et al.*, 2001). Their aim was to reduce ion release as well as improving corrosion resistance for better biocompatibility. They reported that Ti, Al and Zr based oxide dominate the surface where small amount of Nb based oxide formed. They also

Surface Modification Techniques for Biomedical Grade of

the adhesion strength of oxide layer formed on the titanium substrate.

Anodic oxidation is an electrochemical reaction which is a combined phenomenon of diffusion between oxygen and metal ion. In this technique, the metal ion is driven by an electric field. This phenomenon leads to oxide layer formation on the surface of anode (Liu *et al.*, 2004). Anodic oxidation can be used for producing different types of protective oxide layer on different metals. Common electrolytes used in the process are various diluted acids such as H2SO4, H3PO4, acetic acid, etc. The main advantage of anodic oxidation compared to

**4.3 Anodic oxidation** 

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 211

team of researchers made an effort to address the stress shielding issue of TiNb and TiZr based implant alloys (Munuera *et al.*, 2007). They studied the surface structural properties by evaluating the nanoscale elastic properties of oxide layers at various oxidation times. They found that most cases the Young moduli of the oxide layer are lower than 65 GPA and in some cases it is almost near to bone i.e. 20 GPA. In other study, the nanomechanical properties of oxide scale (hardness and Young modulus) was also investigated (Cáceres *et al.*, 2008). They noticed that TiZr based alloy shows increment in hardness and Young modulus after thermal oxidation. However, TiNb shows reduction in hardness and Young modulus at prolonged oxidation duration which is near to bone. Several researchers also investigated the effective oxidation parameters to produce rutile structure through thermal oxidation. As mentioned earlier, rutile structure is more preferable compared to other structures due to better resistant to bacterial attack and also having higher hardness. A group of researchers carried out thermal oxidation process on Ti–6Al–7Nb, Ti–13Nb–13Zr, and Ti–15Zr–4Nb at 750 °C for 24 hours (López *et al.*, 2003). They reported that TiNbZr based alloys present a thicker scale with rutile structure. Other group of researchers studied the effect of producing rutile structure on AISI 316L coated with titanium (Krishna *et al.*, 2005). They found that the presence of rutile structure improves the hardness and corrosion resistance. Another group of workers investigated the effect of thermal oxidation temperature on the Commercial Pure Titanium (CP-Ti) (Kumar *et al.*, 2010a). They reported that rutile structure can be obtained at 800 oC after continuous heating for 24 hours. Another group of researchers carried out experiments to investigate the effects of different pickling times as well as temperature on the adhesion strength of oxide layer formed on the Ti-6Al-4V after oxidation process (S. Izman *et al.*, 2011a). It was revealed that the thickness of oxide layer increases with pickling time but the adhesion strengths become lower. It was also found that the adhesion strength of oxide layer formed on Ti substrate surface increases with the increase of temperature while the thickness of the oxide layer decreased within 40 oC pickling temperature. Izman *et al* took an attempt to evaluate the effect of thermal oxidation temperature on surface morphology and structure of the Ti13Nb13Zr biomedical material (S. Izman *et al.*, 2011c). It is noted that all thermally oxidized samples exhibit the presence of oxides without spallation regardless of the thermal oxidation temperatures. Surface morphology of oxidized substrates changes from smooth to nodular particles-like shape when the oxidation temperature increases from low to high. Rutile structure dominants the surface area when the substrate is thermally oxidized at 850 °C. In summary, thermal oxidation is a simple and low cost method to produce protective oxide layer with rutile structure on titanium alloys. Studies show that the Young modulus of rutile structure is near to that of the bone (less than 65 GPa) and has antibacterial effect, better corrosion and wear resistance. Despite these encouraging properties, limited works have been reported on

explored further to study the chemical composition of the deeper oxide layer surface and found that rutile structure dominates in oxide layer of TiZr based alloy (López *et al.*, 2002). Another group of workers investigated the oxide structure to develop thicker oxide layer for improving the corrosion resistance as well as biocompatibility (Morant *et al.*, 2003). They found that oxide layer was compact and uniform with the granular structure in TiNb based alloy and longitudinal groove structure in TiZr based alloy. It was investigated the corrosion- wear responses by increasing hardness through thermal oxidation of CP-Ti and Ti-6Al-4V (Dearnley *et al.*, 2004). They found that corrosion-wear resistance improved by oxidation where surface of oxidized Ti-6Al-4V is harder than CP-Ti.

Some investigators study the effect of oxidation time on surface roughness of oxide layer through this method to improve surface structure for better corrosion resistance (Gutiérrez *et al.*, 2006). It is noted that higher surface roughness will provide better cell adhesion. They observed that surface roughness increases with the increase of oxidation time. A group of researchers investigated in depth chemical composition of oxide layer to understand the diffusion of elements in the substrate during oxidation (Gutiérrez *et al.*, 2008). Their motivation was to produce thick oxide layer for reducing ion release as well as corrosion protection. They observed that TiZr based alloy showed thicker oxide layer than TiNb based alloy but less homogenous. Another group of researchers studied extensively to optimize the oxidation temperature and time (Kumar *et al.*, 2010b). Their objectives were to produce well adherence oxide layer with rutile structure for improving corrosion resistance and biocompatibility. They observed that best corrosion resistance achieved by oxidation at 650 oC for 24 hr and the hardness increased threefold at 650 oC for 48 hr compare to bare metal.

Excessive wear rate is another issue that limit the usage of titanium alloy in various articulation applications. To address the wear resistance issue, another group of researchers investigated the effect of oxidation and temperature on hardness of the oxide layer formed through thermal oxidation on CP-Ti (Yan *et al.*, 2004). They found that thickness of the oxide layer increases with increasing temperature or time and hardness also increases accordingly. Another team of researchers studied the effect of temperature on adhesion and hardness of the oxide layer through thermal oxidation (Rastkar *et al.*, 2005). Their aim was to improve sliding wear resistance by providing hard surface on TiAl based alloy. They observed that higher temperature oxide layer is non-adherent where lower temperature produced adherent oxide layer and also hardness increases with the oxidation temperature increment. Other group of workers also investigated the effect of oxidation time on oxide layer produced through this method (Guleryuz *et al.*, 2005). They wanted to evaluate the dry sliding wear performance on Ti-6Al-4V by providing hard surface. They observed that hardness and surface roughness increases with the increase of oxidation time and these hard oxide layers show significant improvement in dry sliding wear resistance. Another team of researchers investigated the effect of oxidation time and temperature to developed well adherent rutile based oxide surface in order to improve wear resistance (Biswas *et al.*, 2009). It is also noted that rutile structure provides higher hardness compared to anatase structure. They observed that hardness is proportional to oxidation time as well as temperature. However, higher temperature shows significant increase of hardness compared to higher oxidation time.

An appropriate articulating implant should possess the modulus of elasticity close to the bone. Otherwise, this could lead to stress shielding effect which is a loss of bone density. A

explored further to study the chemical composition of the deeper oxide layer surface and found that rutile structure dominates in oxide layer of TiZr based alloy (López *et al.*, 2002). Another group of workers investigated the oxide structure to develop thicker oxide layer for improving the corrosion resistance as well as biocompatibility (Morant *et al.*, 2003). They found that oxide layer was compact and uniform with the granular structure in TiNb based alloy and longitudinal groove structure in TiZr based alloy. It was investigated the corrosion- wear responses by increasing hardness through thermal oxidation of CP-Ti and Ti-6Al-4V (Dearnley *et al.*, 2004). They found that corrosion-wear resistance improved by

Some investigators study the effect of oxidation time on surface roughness of oxide layer through this method to improve surface structure for better corrosion resistance (Gutiérrez *et al.*, 2006). It is noted that higher surface roughness will provide better cell adhesion. They observed that surface roughness increases with the increase of oxidation time. A group of researchers investigated in depth chemical composition of oxide layer to understand the diffusion of elements in the substrate during oxidation (Gutiérrez *et al.*, 2008). Their motivation was to produce thick oxide layer for reducing ion release as well as corrosion protection. They observed that TiZr based alloy showed thicker oxide layer than TiNb based alloy but less homogenous. Another group of researchers studied extensively to optimize the oxidation temperature and time (Kumar *et al.*, 2010b). Their objectives were to produce well adherence oxide layer with rutile structure for improving corrosion resistance and biocompatibility. They observed that best corrosion resistance achieved by oxidation at 650 oC for 24 hr and the hardness increased threefold at 650 oC for 48 hr compare to bare metal. Excessive wear rate is another issue that limit the usage of titanium alloy in various articulation applications. To address the wear resistance issue, another group of researchers investigated the effect of oxidation and temperature on hardness of the oxide layer formed through thermal oxidation on CP-Ti (Yan *et al.*, 2004). They found that thickness of the oxide layer increases with increasing temperature or time and hardness also increases accordingly. Another team of researchers studied the effect of temperature on adhesion and hardness of the oxide layer through thermal oxidation (Rastkar *et al.*, 2005). Their aim was to improve sliding wear resistance by providing hard surface on TiAl based alloy. They observed that higher temperature oxide layer is non-adherent where lower temperature produced adherent oxide layer and also hardness increases with the oxidation temperature increment. Other group of workers also investigated the effect of oxidation time on oxide layer produced through this method (Guleryuz *et al.*, 2005). They wanted to evaluate the dry sliding wear performance on Ti-6Al-4V by providing hard surface. They observed that hardness and surface roughness increases with the increase of oxidation time and these hard oxide layers show significant improvement in dry sliding wear resistance. Another team of researchers investigated the effect of oxidation time and temperature to developed well adherent rutile based oxide surface in order to improve wear resistance (Biswas *et al.*, 2009). It is also noted that rutile structure provides higher hardness compared to anatase structure. They observed that hardness is proportional to oxidation time as well as temperature. However, higher temperature shows significant increase of hardness compared to higher

An appropriate articulating implant should possess the modulus of elasticity close to the bone. Otherwise, this could lead to stress shielding effect which is a loss of bone density. A

oxidation where surface of oxidized Ti-6Al-4V is harder than CP-Ti.

oxidation time.

team of researchers made an effort to address the stress shielding issue of TiNb and TiZr based implant alloys (Munuera *et al.*, 2007). They studied the surface structural properties by evaluating the nanoscale elastic properties of oxide layers at various oxidation times. They found that most cases the Young moduli of the oxide layer are lower than 65 GPA and in some cases it is almost near to bone i.e. 20 GPA. In other study, the nanomechanical properties of oxide scale (hardness and Young modulus) was also investigated (Cáceres *et al.*, 2008). They noticed that TiZr based alloy shows increment in hardness and Young modulus after thermal oxidation. However, TiNb shows reduction in hardness and Young modulus at prolonged oxidation duration which is near to bone. Several researchers also investigated the effective oxidation parameters to produce rutile structure through thermal oxidation. As mentioned earlier, rutile structure is more preferable compared to other structures due to better resistant to bacterial attack and also having higher hardness. A group of researchers carried out thermal oxidation process on Ti–6Al–7Nb, Ti–13Nb–13Zr, and Ti–15Zr–4Nb at 750 °C for 24 hours (López *et al.*, 2003). They reported that TiNbZr based alloys present a thicker scale with rutile structure. Other group of researchers studied the effect of producing rutile structure on AISI 316L coated with titanium (Krishna *et al.*, 2005). They found that the presence of rutile structure improves the hardness and corrosion resistance. Another group of workers investigated the effect of thermal oxidation temperature on the Commercial Pure Titanium (CP-Ti) (Kumar *et al.*, 2010a). They reported that rutile structure can be obtained at 800 oC after continuous heating for 24 hours. Another group of researchers carried out experiments to investigate the effects of different pickling times as well as temperature on the adhesion strength of oxide layer formed on the Ti-6Al-4V after oxidation process (S. Izman *et al.*, 2011a). It was revealed that the thickness of oxide layer increases with pickling time but the adhesion strengths become lower. It was also found that the adhesion strength of oxide layer formed on Ti substrate surface increases with the increase of temperature while the thickness of the oxide layer decreased within 40 oC pickling temperature. Izman *et al* took an attempt to evaluate the effect of thermal oxidation temperature on surface morphology and structure of the Ti13Nb13Zr biomedical material (S. Izman *et al.*, 2011c). It is noted that all thermally oxidized samples exhibit the presence of oxides without spallation regardless of the thermal oxidation temperatures. Surface morphology of oxidized substrates changes from smooth to nodular particles-like shape when the oxidation temperature increases from low to high. Rutile structure dominants the surface area when the substrate is thermally oxidized at 850 °C. In summary, thermal oxidation is a simple and low cost method to produce protective oxide layer with rutile structure on titanium alloys. Studies show that the Young modulus of rutile structure is near to that of the bone (less than 65 GPa) and has antibacterial effect, better corrosion and wear resistance. Despite these encouraging properties, limited works have been reported on the adhesion strength of oxide layer formed on the titanium substrate.

#### **4.3 Anodic oxidation**

Anodic oxidation is an electrochemical reaction which is a combined phenomenon of diffusion between oxygen and metal ion. In this technique, the metal ion is driven by an electric field. This phenomenon leads to oxide layer formation on the surface of anode (Liu *et al.*, 2004). Anodic oxidation can be used for producing different types of protective oxide layer on different metals. Common electrolytes used in the process are various diluted acids such as H2SO4, H3PO4, acetic acid, etc. The main advantage of anodic oxidation compared to

Surface Modification Techniques for Biomedical Grade of

temperature, anode potential and electrolyte composition.

**4.4 Micro-arc oxidation (MAO)** 

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 213

This also indicates that prior to the formation of apatite on the surfaces, a titanium oxide with three-dimensional micro-porous structure may be essential. It is also noted that surfaces can be bioactive by containing Ca and/or P which leads to osteoinduction of new bones. Wojciech (2011) investigated the effective anodic voltage for producing better corrosion resistance bioactive oxide layer containing Ca and P on TiZr based alloy through anodic oxidation method. He found that lower anodic voltage produced highest corrosion resistance. However, higher anodic voltage provides bioactive oxide layer rich in Ca and P on TiZr based alloy which also increase the corrosion resistance. In summary, anodic oxidation is a simple and effective method of surface modification for providing better bioactive surface of titanium alloys which also homogenous and highly crystalline. However, bioactive apatite formation on titanium alloy through this method required posttreatment such as hydrothermal heat treatment. The oxide film produced by anodic oxidation method shows various properties such as better biocompatibility, corrosion resistance, osteoconductive, etc. These properties rely on the microstructure and composition of the materials as well as anodic oxidation parameters, such as current,

Another name of micro-arc oxidation is anodic spark oxidation or plasma electrolytic oxidation (PEO). Micro arc oxidation is an electrochemical surface modification process for producing oxide coatings on metals such as Al, Ti, Mg, Ta, W, Zn, and Zr and their alloys (Liu *et al.*, 2010). According to Yerokhin *et al.* (1999), MAO can be defined as a complex plasma-enhanced physico-chemical process which involved micro-arc discharge, diffusion and plasma chemical reactions. Basically, it is a new type of anodic oxidation technique, but the difference between MAO and the conventional anodic oxidation is it employs higher potentials to discharges and the resulting plasma modifies the structure of oxide layer. This process can be used to grow crystalline oxide coating with thickness range from ten to hundreds µm. The coating thickness depends on process parameters such as current density, process time, electrolyte temperature, applied voltage, electrolyte composition, alloy composition (Dunleavy *et al.*, 2009). A large number of short-lived sparks (micro arc discharges) produced in MAO process is a result of localized electrical breakdown of the growing coating. These discharges play the key role in the coating growth mechanism as they deposit 'craters' on the free surface of the growing coating. In MAO process, the anode is immersed in electrolyte which is an aqueous solution. The anode is made from valve metals. Valve metals usually refer to Ti, Al, Mg, Ta, W, Zn, and Zr due to their usage as a cathode to emit electron in electronic valve. They are also known as 'thermionic valve' materials in early days. However, disputes on the right definition of these terms have been remained among researchers since Al is not a suitable material for high temperature resistance to emit electron. In MAO, an unequal alternating voltage between the anode and cathode initiates an electrical discharge at the anode. The typical voltage range for anode and cathode is from 150 to 1000V and from 0 to 100V respectively. Temperature and local pressure in the discharge channels are among the parameters that affect the MAO coating qualities such as high strength, well adhesion, high micro-hardness, and wear resistance (Liu *et al.*, 2010). Since MAO can provide high hardness and a continuous barrier, this coating is suitable for protection against wear, corrosion or heat as well as electrical insulation (Curran and Clyne, 2005). General characteristics of these coatings are porous,

other oxidation methods is their ability to form bioactive oxide film on surface of titanium and its alloy. Anodic oxidation increases thickness of the oxide layer for reducing ion release as well as improving corrosion protection. By varying the anodic oxidation parameters, such as current, process temperature, electrolyte composition, and anode potential, the oxide film properties i.e. chemical or structural can be changed.

Principal reactions cause oxidations at the anode are as follows: At the Ti/Ti oxide interfaces: Ti ↔ Ti2+ + 2e-At the Ti oxide/electrolyte interfaces: 2H2O ↔ 2O2- + 4H+ (oxygen ions react with Ti to form oxide) , 2H2O ↔ O2 (gas) + 4H+ + 4e- (O2 gas evolves or stick at anode surface). At both interfaces: Ti2+ + 2O2- ↔ TiO2 + 2e-

In anodic oxidation, a linear correlation exists between the oxide film thickness and applied voltage. If the final oxide thickness is d and the applied voltage is U, then the relationship is d = αU. α is a constant and its typical range is 1.5–3 nmV-1. Ishizawa and Ogino *et al.* is the pioneer in developing Ca and P contained oxide layer through anodizing titanium in βglycerophosphate sodium and calcium acetate contained electrolyte (Ishizawa and Ogino, 1995). They further proceeded exploring and able to transform it into hydroxyapatite by applying hydrothermal treatment. The results showed that the electrolyte possessed some impurities (e.g. sodium). These impurities decreased oxide layer's strength. A group of researchers reported that desirable cellular behaviour such as cell growth, cell attachment, etc. can be obtained from the thin HA layer on the surface of CP-Ti which was produced by anodization and subsequently followed by hydrothermal treatment (Takebe *et al.*, 2000). It is observed that cellular attachment and spreading are affected by this thin HA layer on the CP-Ti surface. It is also revealed a thin HA layer on titanium surface shows more osteoconductive behaviour to cell attachment as compared to bare CP-Ti. Other group of workers investigated new electrolytes consists of calcium glycerophosphate and calcium acetate for producing anodic oxide films that consist of Ca and P on titanium implants (Zhu *et al.*, 2001). The anodic oxide film of titanium obtained using this method is highly crystalline with porous structure and rich in Ca and P. The recommended optimum conditions are: (i) 350 V as final voltage, (ii) 70 A m-2 as current density, and (iii) concentrations of the calcium glycerophosphate (0.02M) and calcium acetate (0.15 M). Ca and P ratio near to 1.67 was achieved using this recommended condition. Positive biological response also observed from the properties of that anodic oxide layer surface. Yang *et al.* reported that using anodic oxidation in H2SO4 solution united with consequent heat treatment is an efficient method for obtaining titanium alloy with bioactive surface (Yang *et al.*, 2004). They also observed that the porous structure Titania of anatase and/or rutile phase covered on the surface after anodic oxidation. It was interesting to observe that apatite can be formed on titanium alloy by anodic oxidation in simulated body fluid. The initial time for apatite formation was inversely proportional to the quantity of rutile or anatase phase (Liu *et al.*, 2004). Apatite cannot be formed without spark discharge on the surface although anatase was produced on anodically oxidized titanium. Hence, a combination of anodic oxidation with heat treatment is required for the apatite formation on titanium in SBF without spark discharge treatment. Heat treatment induces apatite formation in SBF since the amount of anatase and/or rutile increases by the heat treatment.

This also indicates that prior to the formation of apatite on the surfaces, a titanium oxide with three-dimensional micro-porous structure may be essential. It is also noted that surfaces can be bioactive by containing Ca and/or P which leads to osteoinduction of new bones. Wojciech (2011) investigated the effective anodic voltage for producing better corrosion resistance bioactive oxide layer containing Ca and P on TiZr based alloy through anodic oxidation method. He found that lower anodic voltage produced highest corrosion resistance. However, higher anodic voltage provides bioactive oxide layer rich in Ca and P on TiZr based alloy which also increase the corrosion resistance. In summary, anodic oxidation is a simple and effective method of surface modification for providing better bioactive surface of titanium alloys which also homogenous and highly crystalline. However, bioactive apatite formation on titanium alloy through this method required posttreatment such as hydrothermal heat treatment. The oxide film produced by anodic oxidation method shows various properties such as better biocompatibility, corrosion resistance, osteoconductive, etc. These properties rely on the microstructure and composition of the materials as well as anodic oxidation parameters, such as current, temperature, anode potential and electrolyte composition.

#### **4.4 Micro-arc oxidation (MAO)**

212 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

other oxidation methods is their ability to form bioactive oxide film on surface of titanium and its alloy. Anodic oxidation increases thickness of the oxide layer for reducing ion release as well as improving corrosion protection. By varying the anodic oxidation parameters, such as current, process temperature, electrolyte composition, and anode potential, the oxide film

(O2 gas evolves or stick at anode surface).

In anodic oxidation, a linear correlation exists between the oxide film thickness and applied voltage. If the final oxide thickness is d and the applied voltage is U, then the relationship is d = αU. α is a constant and its typical range is 1.5–3 nmV-1. Ishizawa and Ogino *et al.* is the pioneer in developing Ca and P contained oxide layer through anodizing titanium in βglycerophosphate sodium and calcium acetate contained electrolyte (Ishizawa and Ogino, 1995). They further proceeded exploring and able to transform it into hydroxyapatite by applying hydrothermal treatment. The results showed that the electrolyte possessed some impurities (e.g. sodium). These impurities decreased oxide layer's strength. A group of researchers reported that desirable cellular behaviour such as cell growth, cell attachment, etc. can be obtained from the thin HA layer on the surface of CP-Ti which was produced by anodization and subsequently followed by hydrothermal treatment (Takebe *et al.*, 2000). It is observed that cellular attachment and spreading are affected by this thin HA layer on the CP-Ti surface. It is also revealed a thin HA layer on titanium surface shows more osteoconductive behaviour to cell attachment as compared to bare CP-Ti. Other group of workers investigated new electrolytes consists of calcium glycerophosphate and calcium acetate for producing anodic oxide films that consist of Ca and P on titanium implants (Zhu *et al.*, 2001). The anodic oxide film of titanium obtained using this method is highly crystalline with porous structure and rich in Ca and P. The recommended optimum conditions are: (i) 350 V as final voltage, (ii) 70 A m-2 as current density, and (iii) concentrations of the calcium glycerophosphate (0.02M) and calcium acetate (0.15 M). Ca and P ratio near to 1.67 was achieved using this recommended condition. Positive biological response also observed from the properties of that anodic oxide layer surface. Yang *et al.* reported that using anodic oxidation in H2SO4 solution united with consequent heat treatment is an efficient method for obtaining titanium alloy with bioactive surface (Yang *et al.*, 2004). They also observed that the porous structure Titania of anatase and/or rutile phase covered on the surface after anodic oxidation. It was interesting to observe that apatite can be formed on titanium alloy by anodic oxidation in simulated body fluid. The initial time for apatite formation was inversely proportional to the quantity of rutile or anatase phase (Liu *et al.*, 2004). Apatite cannot be formed without spark discharge on the surface although anatase was produced on anodically oxidized titanium. Hence, a combination of anodic oxidation with heat treatment is required for the apatite formation on titanium in SBF without spark discharge treatment. Heat treatment induces apatite formation in SBF since the amount of anatase and/or rutile increases by the heat treatment.

properties i.e. chemical or structural can be changed.

At the Ti/Ti oxide interfaces:

2H2O ↔ O2 (gas) + 4H+ + 4e-

At the Ti oxide/electrolyte interfaces:

Ti ↔ Ti2+ + 2e-

At both interfaces: Ti2+ + 2O2- ↔ TiO2 + 2e-

Principal reactions cause oxidations at the anode are as follows:

2H2O ↔ 2O2- + 4H+ (oxygen ions react with Ti to form oxide) ,

Another name of micro-arc oxidation is anodic spark oxidation or plasma electrolytic oxidation (PEO). Micro arc oxidation is an electrochemical surface modification process for producing oxide coatings on metals such as Al, Ti, Mg, Ta, W, Zn, and Zr and their alloys (Liu *et al.*, 2010). According to Yerokhin *et al.* (1999), MAO can be defined as a complex plasma-enhanced physico-chemical process which involved micro-arc discharge, diffusion and plasma chemical reactions. Basically, it is a new type of anodic oxidation technique, but the difference between MAO and the conventional anodic oxidation is it employs higher potentials to discharges and the resulting plasma modifies the structure of oxide layer. This process can be used to grow crystalline oxide coating with thickness range from ten to hundreds µm. The coating thickness depends on process parameters such as current density, process time, electrolyte temperature, applied voltage, electrolyte composition, alloy composition (Dunleavy *et al.*, 2009). A large number of short-lived sparks (micro arc discharges) produced in MAO process is a result of localized electrical breakdown of the growing coating. These discharges play the key role in the coating growth mechanism as they deposit 'craters' on the free surface of the growing coating. In MAO process, the anode is immersed in electrolyte which is an aqueous solution. The anode is made from valve metals. Valve metals usually refer to Ti, Al, Mg, Ta, W, Zn, and Zr due to their usage as a cathode to emit electron in electronic valve. They are also known as 'thermionic valve' materials in early days. However, disputes on the right definition of these terms have been remained among researchers since Al is not a suitable material for high temperature resistance to emit electron. In MAO, an unequal alternating voltage between the anode and cathode initiates an electrical discharge at the anode. The typical voltage range for anode and cathode is from 150 to 1000V and from 0 to 100V respectively. Temperature and local pressure in the discharge channels are among the parameters that affect the MAO coating qualities such as high strength, well adhesion, high micro-hardness, and wear resistance (Liu *et al.*, 2010). Since MAO can provide high hardness and a continuous barrier, this coating is suitable for protection against wear, corrosion or heat as well as electrical insulation (Curran and Clyne, 2005). General characteristics of these coatings are porous,

Surface Modification Techniques for Biomedical Grade of

**5.1 Basic mechanism of carburization** 

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 215

the carburized layers due to variation of carbon concentration in the surface region (Saleh *et al.*, 2010). The discussion of this chapter starts with the basic mechanism of carburization

Carburization is a process widely used method to harden the surface and enhance the properties of components that made from metal. Carburizing consists of absorption and diffusion of carbon into solid metal alloys by heating at high temperature. Historically, the carburizing process is generally done at elevated temperatures with a carbon medium that can supply adequate quantity of atomic carbon for absorption and diffusion into the steel (Luo *et al.*, 2009). The carbon medium that use for carburizing process can be solid (charcoal), molten salt (cyanide), a gaseous or plasma medium (Prabhudev, 1998). There are three methods of carburizing process, i.e. solid carburizing, liquid carburizing, and gas carburizing. All these three methods have their own compounds medium that is used for the carbon supply during the process. In solid carburizing process, carburizing compound such as charcoal or graphite powder is used for its medium. In the liquid carburizing method, molten cyanide is used for carbon enrichment. Lastly, for the gas carburizing

followed by three popular carburization methods, i.e. thermal, gas and laser melting.

method, hydrocarbon gas or plasma is used as the source of the carburizing medium.

results increase in hardness of the substrate materials surface.

**5.2 Thermal carburization** 

During carburizing, the atomic carbon is liberated from carbonaceous medium due to decomposition of carbon monoxide into carbon dioxide and atomic carbon as given below:

Then, the carbon atom from carburizing medium is transferred to the surface of the metal. These metal surfaces will absorb the carbon and diffuse deep into it. Thus, this phenomena

Thermal carburization process is considered the earliest carburization technique and it is a kind of solid carburization. Generally solid particle such as charcoal, graphite powder, etc is used as a carbon source to surround titanium substrate during carburization process. Titanium can easily react with oxygen in ambient environment and form a thin passive layer of TiO2 on the outer surface with thickness range of 3 to 7 nm (Liu *et al.*, 2004) . This passive layer becomes a barrier for carbon atom diffusion into the titanium surface. Since titanium is highly affinity to oxygen, an inert or vacuum environment is preferable for conducting carburization process. Argon gas is commonly used as a medium to remove oxygen in tube or muffle furnace heating chamber from pre-oxidizing the titanium substrate surface. The quality of carburized layer largely depends on the carburizing temperature, soaking time, source of carbon (type and particle size) and the absence of oxygen level in the chamber. The carburizing parameters may have significant effects on the thickness, adhesion, density and chemical composition of carburized layer formed on the titanium substrate. Studies on titanium carbide powder synthesis by carbothermal method in argon environment requires high temperature in the range of 1700–2100 °C (Weimer, 1997) and long reaction time (10–24 h) (Gotoh *et al*., 2001). Other workers tried to synthesize TiC powder at lower temperatures and shorter time with success. For instance, Lee *et al.* studied the chemical kinetics at various

2CO CO2 + Cat (6)

firm adhesion to substrates and the pores are homogenously distributed on the coating's surface with nanostructure grains (Kim *et al.*, 2002). Due to superior corrosion resistance, thermal stability, photocatalytic activity, wear resistance and CO sensing properties makes MAO coatings as a popular research area (Shin *et al.*, 2006, Jin *et al.*, 2008). MAO has been popular in the biomedical community since Ishizawa *et al.* pioneered the technique to biomedical titanium implants (Ishizawa and Ogino, 1995). Biomimetic deposition of apatite is possible on Ca and P-containing MAO coatings (Song et al., 2004). Zhao *et al.* found that the MAO coatings benefit osteoblast adhesion (Zhao *et al.*, 2007). They compared the adhesion performance of MAO coatings on various modified smooth and rough surfaces. Other researchers investigated the effect of variations in the electrolyte compositions to produce different kinds of nanostructured composite coatings under this method (Kim et al., 2007, Yao *et al.*, 2008) . Cimenoglu *et al.* investigated the MAO coating on Ti6Al7Nb and found that oxide layer shows grainy appearance rather than porous and contained calcium titanate precipitates, HA and rutile structure (Cimenoglu *et al.*, 2011). In summary, MAO is a potential method for producing porous nanostructured coatings on Ti and its alloys which promote best osteoblast cell adhesion. This technique has been spreading into the field of orthopaedic and dental implant materials.

#### **5. Carburization of titanium alloy**

Poor tribological properties limit the usefulness of titanium alloy in many engineering applications (Bloyce *et al.*, 1994). Moreover, not all titanium and its alloys can meet all of the clinical requirements. In order to improve the biological, chemical, and mechanical properties, surface modification is often performed (Huang *et al.*, 2006, Kumar *et al.*, 2010b). Till now various surface modification techniques by thermo-chemical process have been studied and applied for improving wear resistance of titanium alloys. These are carburizing, nitriding and oxidation (Biswas *et al.*, 2009, Tsuji et al., 2009b, Savaloni et al., 2010). Among them, carburization technique is one of the methods that can be used to form hard ceramic coating on titanium alloys. The main objective of carburization is to provide hard surface on titanium and its alloys for increasing wear resistance in articulation application since titanium carbide is one of the potential biocompatible carbide layers (Bharathy *et al.*, 2010). It is also one of the cost-effective surface modification methods to deliberately generate a carbide layer on titanium alloy. Many researchers reported that the carbide layer enables to increase hardness, wear resistance and corrosion resistance to titanium and its alloy (Kim *et al.*, 2003). Sintered solid titanium carbide is a very important non-oxide ceramics that widely used in the fields of wear resistance tools and materials due to its high melting point (3170 oC), low density, high hardness (2500 ~3000HV), superior chemical and thermal stability, and outstanding wear resistance (Courant *et al.*, 2005). Apart from sintering, titanium carbide layer can be created by other surface modification methods, such as plasma carburizing process, thermal carburization or high-temperature synthesis, carburization by laser melting, gas-solid reaction or gas carburization and sol-gel process (Lee, 1997, Yin *et al.*, 2005, Cochepin *et al.*, 2007, Luo *et al.*, 2011). Among these methods, thermal carburization process is considered as the simplest and the most cost effective. Typically, one of the main obstacles for TiC coating is the high affinity of titanium to oxygen which leads to form TiO2 easily on the surface. To overcome this issue , vacuum carburization or inert gas environment is introduced to remove O2 contents in carburization chamber (Wu *et al.*, 1997). Another common problem related to carburization is non uniform hardness profile across the carburized layers due to variation of carbon concentration in the surface region (Saleh *et al.*, 2010). The discussion of this chapter starts with the basic mechanism of carburization followed by three popular carburization methods, i.e. thermal, gas and laser melting.

#### **5.1 Basic mechanism of carburization**

214 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

firm adhesion to substrates and the pores are homogenously distributed on the coating's surface with nanostructure grains (Kim *et al.*, 2002). Due to superior corrosion resistance, thermal stability, photocatalytic activity, wear resistance and CO sensing properties makes MAO coatings as a popular research area (Shin *et al.*, 2006, Jin *et al.*, 2008). MAO has been popular in the biomedical community since Ishizawa *et al.* pioneered the technique to biomedical titanium implants (Ishizawa and Ogino, 1995). Biomimetic deposition of apatite is possible on Ca and P-containing MAO coatings (Song et al., 2004). Zhao *et al.* found that the MAO coatings benefit osteoblast adhesion (Zhao *et al.*, 2007). They compared the adhesion performance of MAO coatings on various modified smooth and rough surfaces. Other researchers investigated the effect of variations in the electrolyte compositions to produce different kinds of nanostructured composite coatings under this method (Kim et al., 2007, Yao *et al.*, 2008) . Cimenoglu *et al.* investigated the MAO coating on Ti6Al7Nb and found that oxide layer shows grainy appearance rather than porous and contained calcium titanate precipitates, HA and rutile structure (Cimenoglu *et al.*, 2011). In summary, MAO is a potential method for producing porous nanostructured coatings on Ti and its alloys which promote best osteoblast cell adhesion. This technique has been spreading into the field of

Poor tribological properties limit the usefulness of titanium alloy in many engineering applications (Bloyce *et al.*, 1994). Moreover, not all titanium and its alloys can meet all of the clinical requirements. In order to improve the biological, chemical, and mechanical properties, surface modification is often performed (Huang *et al.*, 2006, Kumar *et al.*, 2010b). Till now various surface modification techniques by thermo-chemical process have been studied and applied for improving wear resistance of titanium alloys. These are carburizing, nitriding and oxidation (Biswas *et al.*, 2009, Tsuji et al., 2009b, Savaloni et al., 2010). Among them, carburization technique is one of the methods that can be used to form hard ceramic coating on titanium alloys. The main objective of carburization is to provide hard surface on titanium and its alloys for increasing wear resistance in articulation application since titanium carbide is one of the potential biocompatible carbide layers (Bharathy *et al.*, 2010). It is also one of the cost-effective surface modification methods to deliberately generate a carbide layer on titanium alloy. Many researchers reported that the carbide layer enables to increase hardness, wear resistance and corrosion resistance to titanium and its alloy (Kim *et al.*, 2003). Sintered solid titanium carbide is a very important non-oxide ceramics that widely used in the fields of wear resistance tools and materials due to its high melting point (3170 oC), low density, high hardness (2500 ~3000HV), superior chemical and thermal stability, and outstanding wear resistance (Courant *et al.*, 2005). Apart from sintering, titanium carbide layer can be created by other surface modification methods, such as plasma carburizing process, thermal carburization or high-temperature synthesis, carburization by laser melting, gas-solid reaction or gas carburization and sol-gel process (Lee, 1997, Yin *et al.*, 2005, Cochepin *et al.*, 2007, Luo *et al.*, 2011). Among these methods, thermal carburization process is considered as the simplest and the most cost effective. Typically, one of the main obstacles for TiC coating is the high affinity of titanium to oxygen which leads to form TiO2 easily on the surface. To overcome this issue , vacuum carburization or inert gas environment is introduced to remove O2 contents in carburization chamber (Wu *et al.*, 1997). Another common problem related to carburization is non uniform hardness profile across

orthopaedic and dental implant materials.

**5. Carburization of titanium alloy** 

Carburization is a process widely used method to harden the surface and enhance the properties of components that made from metal. Carburizing consists of absorption and diffusion of carbon into solid metal alloys by heating at high temperature. Historically, the carburizing process is generally done at elevated temperatures with a carbon medium that can supply adequate quantity of atomic carbon for absorption and diffusion into the steel (Luo *et al.*, 2009). The carbon medium that use for carburizing process can be solid (charcoal), molten salt (cyanide), a gaseous or plasma medium (Prabhudev, 1998). There are three methods of carburizing process, i.e. solid carburizing, liquid carburizing, and gas carburizing. All these three methods have their own compounds medium that is used for the carbon supply during the process. In solid carburizing process, carburizing compound such as charcoal or graphite powder is used for its medium. In the liquid carburizing method, molten cyanide is used for carbon enrichment. Lastly, for the gas carburizing method, hydrocarbon gas or plasma is used as the source of the carburizing medium.

During carburizing, the atomic carbon is liberated from carbonaceous medium due to decomposition of carbon monoxide into carbon dioxide and atomic carbon as given below:

$$\text{2CO} \quad \text{\rightarrow CO}\_2 + \text{C}\_{\text{at}} \tag{6}$$

Then, the carbon atom from carburizing medium is transferred to the surface of the metal. These metal surfaces will absorb the carbon and diffuse deep into it. Thus, this phenomena results increase in hardness of the substrate materials surface.

#### **5.2 Thermal carburization**

Thermal carburization process is considered the earliest carburization technique and it is a kind of solid carburization. Generally solid particle such as charcoal, graphite powder, etc is used as a carbon source to surround titanium substrate during carburization process. Titanium can easily react with oxygen in ambient environment and form a thin passive layer of TiO2 on the outer surface with thickness range of 3 to 7 nm (Liu *et al.*, 2004) . This passive layer becomes a barrier for carbon atom diffusion into the titanium surface. Since titanium is highly affinity to oxygen, an inert or vacuum environment is preferable for conducting carburization process. Argon gas is commonly used as a medium to remove oxygen in tube or muffle furnace heating chamber from pre-oxidizing the titanium substrate surface. The quality of carburized layer largely depends on the carburizing temperature, soaking time, source of carbon (type and particle size) and the absence of oxygen level in the chamber. The carburizing parameters may have significant effects on the thickness, adhesion, density and chemical composition of carburized layer formed on the titanium substrate. Studies on titanium carbide powder synthesis by carbothermal method in argon environment requires high temperature in the range of 1700–2100 °C (Weimer, 1997) and long reaction time (10–24 h) (Gotoh *et al*., 2001). Other workers tried to synthesize TiC powder at lower temperatures and shorter time with success. For instance, Lee *et al.* studied the chemical kinetics at various

Surface Modification Techniques for Biomedical Grade of

plasticity and fatigue strength in the titanium substrate.

**5.4 Carburization by laser melting** 

alloyed zone or carburized layer.

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 217

They revealed that fine and homogeneous dispersion of hard carbide particles such as TiC and V4C3 found in the carburized layer able to improve wear resistance as well as fatigue life for more than two folds. Tsuji *et al.* carburized Ti6Al4V at 600 oC in a Ar gas conditioned furnace using CH4-H2 plasma for 1 hr for improving hardness from 400 HV to 600 HV (Tsuji *et al.*, 2009a). They also investigated the effects of combining plasma-carburizing and deep rolling on the notched surface microstructure and morphology, micro-hardness and notch fatigue life of Ti-6Al-4V alloy specimen in a laboratory at an ambient temperature. They reported that the notch root area of plasma-carburized specimen's surface roughness has been significantly improved by deep-rolling. This method effectively introduces compressive residual stress and work hardening in the substrate. Plasma-carburization with subsequent deep-rolling largely enhances the notch fatigue strength of specimen in comparison with untreated specimen. The developed compressive residual stress and work hardening zone influence the initial crack growth rate of deep-rolled carburized specimen. The thickness of this zone is approximately 350 µm depth from the surface. However, the crack rapidly propagates toward the inside after it passes through this zone. They concluded that plasma-carburizing process combined with deep-rolling effectively improves the notch fatigue properties of Ti-6Al-4V alloy. Another researcher made an effort to investigate the plasticity effect on titanium alloy after being treated under gas carburization. Luo *et al*. carburized Ti6Al4V at 1050 oC in a vacuum furnace using C2H2 gas for 4 hrs for improving the of hardness from 350 HV to 778 HV (Luo *et al.*, 2011). TiC or also called titanium cermets were successfully formed on the surface. It was reported that the plasticity of the titanium cermets was slightly lower (10.86%) than original titanium bare material. This indicates that the carburized titanium has significantly improved in fracture toughness as compared to typical ceramics material. They concluded that carburization is a way to produce titanium cermets efficiently which consists of hard surface, high toughness and plasticity. All these properties make titanium carbide as a potential candidate for artificial articulation material. In summary, the primary objective of gas carburizing is to produce carburized layer on the substrate in order to increase wear resistance property of titanium alloys. However, improvement in hardness introduces other issues such as reduction in

Laser carburizing technique is developed from laser surface hardening of steel. In a simple way, laser carburization can be defined as a process of using laser as a source of high energy to perform carburization. There are various types of laser carburizing methods where the categories are based on laser source, such as Neodymium Yttrium Lithium Fluoride (Nd:YLF), Neodymium Yttrium Aluminium Garnet (Nd:YAG), Titanium Sapphire (Ti:Sapphire), CO2 laser, etc. Laser carburizing process involves carbon diffusion into the metal substrate using laser irradiation. The typical source of carbon is graphite powder. Other type of powder such as TiC is also being used in laser melting technique to form carburized layer on titanium based materials. Fig. 3 shows a schematic diagram of laser melting working principle. This process involved heating of specimen through continuous or pulse wave laser irradiation, rapid melting, intermixing or diffusion of carbon particle, and rapid solidification of the pre-deposited alloying elements on substrate to form an

temperatures (1100 to 1400 oC) for synthesising TiC from CP-Ti alloy and graphite powders (Lee and Thadhani, 1997). They found that Ti with compacted graphite powder shows highly activated state of reactants which reduce activation energies by 4-6 times, undergo a solid state diffusion reaction. They also concluded that increasing temperature will increase the rate of heat released. This released heat generates localized melting of unreacted Ti and initiate a combustion reaction. It has been reported that the carburizing rate of titanium dioxide, TiO2 into TiC can be accelerated by using the finest and homogenous carbon powder (Maitre *et al.*, 2000). Sen *et al* produced fine and homogeneous TiC powders by carbothermal reduction of titania/charcoal in a vacuum furnace at different reaction temperatures from 1100 °C to 1550 °C (Sen *et al.*, 2010). They observed that reaction temperature increases, uniform crystal grain arises with the liberation of much CO and higher temperature (at 1550 oC) produced large amount of TiC. They also noticed that as reaction temperature increased, formation of the compounds was in sequenced as Ti4O7, Ti3O5, Ti2O3, TiCxO1−x and TiC. Hardly found researchers study TiC formation on titanium solid substrate. Izman *et al* initiated the study to investigate the effects of different carburizing times on the adhesion strength of carbide layer formed on the Ti-6Al-7Nb (S. Izman *et al.*, 2011b). Prior to carburization process, all samples were treated to remove residual stress and oxide scales by annealing and pickling processes respectively. Hard wood charcoal powder was used as a medium. The carburizing process was carried out under normal atmospheric condition. They found that a mixture of oxide and carbide layers formed on the substrate and the thickness of these layers increases with carburizing time. It was also revealed that the longer carburizing time provides the strongest adhesion strength and TiC as the dominant layer. Porous structure of TiC was observed and this structure is believed able to facilitate the osteoblast cell growth on implant. In summary, thermal carburization is a simple and cost effective method to produced TiC for increasing the wear resistance properties of titanium and its alloys. However, the technique has not been explored rigorously this far. Issues regarding carbide grain growth, carbon particle agglomeration, non-uniform carbide particle shapes and large amounts of unreacted TiO2 and carbon in the substrate are still under on-going research.

#### **5.3 Gas carburization**

The main difference between thermal and gas carburization process is the carbon source medium. Instead of solid, hydrocarbon gas is used as a carbon source and carburization process takes place either under gaseous or plasma condition. This process is typically performed using plasma or flowing hydrocarbon gas over the Ti and Ti alloy substrate at high temperature in a inert gas or vacuum furnace. Gas carburizing also have various categories such as hydrocarbon gas carburizing (using methane or ethane), plasma carburizing, etc. The advantage of gas carburizing over solid carburising is faster processing time but this method is costlier compared to solid carburization (Robert *et al.*, 1994). Due to high affinity to oxygen, plasma carburizing method has difficulties in carburizing the Ti alloys because thin protective titanium oxide film easily forms on its surface which cause in obstruction of the carbon diffusion (Okamoto *et al.*, 2001). Kim *et al* carburized Ti6Al4V at 900 oC and 250MPa pressure using CH4–Ar-H2 plasma for 6 hrs to increase the wear resistance (Kim *et al.*, 2003). Hardness of titanium alloy was improved significantly from 400HV to 1600HV with the carburized layer thickness of about 150 µm along the surface.

temperatures (1100 to 1400 oC) for synthesising TiC from CP-Ti alloy and graphite powders (Lee and Thadhani, 1997). They found that Ti with compacted graphite powder shows highly activated state of reactants which reduce activation energies by 4-6 times, undergo a solid state diffusion reaction. They also concluded that increasing temperature will increase the rate of heat released. This released heat generates localized melting of unreacted Ti and initiate a combustion reaction. It has been reported that the carburizing rate of titanium dioxide, TiO2 into TiC can be accelerated by using the finest and homogenous carbon powder (Maitre *et al.*, 2000). Sen *et al* produced fine and homogeneous TiC powders by carbothermal reduction of titania/charcoal in a vacuum furnace at different reaction temperatures from 1100 °C to 1550 °C (Sen *et al.*, 2010). They observed that reaction temperature increases, uniform crystal grain arises with the liberation of much CO and higher temperature (at 1550 oC) produced large amount of TiC. They also noticed that as reaction temperature increased, formation of the compounds was in sequenced as Ti4O7, Ti3O5, Ti2O3, TiCxO1−x and TiC. Hardly found researchers study TiC formation on titanium solid substrate. Izman *et al* initiated the study to investigate the effects of different carburizing times on the adhesion strength of carbide layer formed on the Ti-6Al-7Nb (S. Izman *et al.*, 2011b). Prior to carburization process, all samples were treated to remove residual stress and oxide scales by annealing and pickling processes respectively. Hard wood charcoal powder was used as a medium. The carburizing process was carried out under normal atmospheric condition. They found that a mixture of oxide and carbide layers formed on the substrate and the thickness of these layers increases with carburizing time. It was also revealed that the longer carburizing time provides the strongest adhesion strength and TiC as the dominant layer. Porous structure of TiC was observed and this structure is believed able to facilitate the osteoblast cell growth on implant. In summary, thermal carburization is a simple and cost effective method to produced TiC for increasing the wear resistance properties of titanium and its alloys. However, the technique has not been explored rigorously this far. Issues regarding carbide grain growth, carbon particle agglomeration, non-uniform carbide particle shapes and large amounts of unreacted TiO2

and carbon in the substrate are still under on-going research.

The main difference between thermal and gas carburization process is the carbon source medium. Instead of solid, hydrocarbon gas is used as a carbon source and carburization process takes place either under gaseous or plasma condition. This process is typically performed using plasma or flowing hydrocarbon gas over the Ti and Ti alloy substrate at high temperature in a inert gas or vacuum furnace. Gas carburizing also have various categories such as hydrocarbon gas carburizing (using methane or ethane), plasma carburizing, etc. The advantage of gas carburizing over solid carburising is faster processing time but this method is costlier compared to solid carburization (Robert *et al.*, 1994). Due to high affinity to oxygen, plasma carburizing method has difficulties in carburizing the Ti alloys because thin protective titanium oxide film easily forms on its surface which cause in obstruction of the carbon diffusion (Okamoto *et al.*, 2001). Kim *et al* carburized Ti6Al4V at 900 oC and 250MPa pressure using CH4–Ar-H2 plasma for 6 hrs to increase the wear resistance (Kim *et al.*, 2003). Hardness of titanium alloy was improved significantly from 400HV to 1600HV with the carburized layer thickness of about 150 µm along the surface.

**5.3 Gas carburization** 

They revealed that fine and homogeneous dispersion of hard carbide particles such as TiC and V4C3 found in the carburized layer able to improve wear resistance as well as fatigue life for more than two folds. Tsuji *et al.* carburized Ti6Al4V at 600 oC in a Ar gas conditioned furnace using CH4-H2 plasma for 1 hr for improving hardness from 400 HV to 600 HV (Tsuji *et al.*, 2009a). They also investigated the effects of combining plasma-carburizing and deep rolling on the notched surface microstructure and morphology, micro-hardness and notch fatigue life of Ti-6Al-4V alloy specimen in a laboratory at an ambient temperature. They reported that the notch root area of plasma-carburized specimen's surface roughness has been significantly improved by deep-rolling. This method effectively introduces compressive residual stress and work hardening in the substrate. Plasma-carburization with subsequent deep-rolling largely enhances the notch fatigue strength of specimen in comparison with untreated specimen. The developed compressive residual stress and work hardening zone influence the initial crack growth rate of deep-rolled carburized specimen. The thickness of this zone is approximately 350 µm depth from the surface. However, the crack rapidly propagates toward the inside after it passes through this zone. They concluded that plasma-carburizing process combined with deep-rolling effectively improves the notch fatigue properties of Ti-6Al-4V alloy. Another researcher made an effort to investigate the plasticity effect on titanium alloy after being treated under gas carburization. Luo *et al*. carburized Ti6Al4V at 1050 oC in a vacuum furnace using C2H2 gas for 4 hrs for improving the of hardness from 350 HV to 778 HV (Luo *et al.*, 2011). TiC or also called titanium cermets were successfully formed on the surface. It was reported that the plasticity of the titanium cermets was slightly lower (10.86%) than original titanium bare material. This indicates that the carburized titanium has significantly improved in fracture toughness as compared to typical ceramics material. They concluded that carburization is a way to produce titanium cermets efficiently which consists of hard surface, high toughness and plasticity. All these properties make titanium carbide as a potential candidate for artificial articulation material. In summary, the primary objective of gas carburizing is to produce carburized layer on the substrate in order to increase wear resistance property of titanium alloys. However, improvement in hardness introduces other issues such as reduction in plasticity and fatigue strength in the titanium substrate.

#### **5.4 Carburization by laser melting**

Laser carburizing technique is developed from laser surface hardening of steel. In a simple way, laser carburization can be defined as a process of using laser as a source of high energy to perform carburization. There are various types of laser carburizing methods where the categories are based on laser source, such as Neodymium Yttrium Lithium Fluoride (Nd:YLF), Neodymium Yttrium Aluminium Garnet (Nd:YAG), Titanium Sapphire (Ti:Sapphire), CO2 laser, etc. Laser carburizing process involves carbon diffusion into the metal substrate using laser irradiation. The typical source of carbon is graphite powder. Other type of powder such as TiC is also being used in laser melting technique to form carburized layer on titanium based materials. Fig. 3 shows a schematic diagram of laser melting working principle. This process involved heating of specimen through continuous or pulse wave laser irradiation, rapid melting, intermixing or diffusion of carbon particle, and rapid solidification of the pre-deposited alloying elements on substrate to form an alloyed zone or carburized layer.

Surface Modification Techniques for Biomedical Grade of

two carburization techniques.

**6. Ion implantation and deposition** 

affected by the ion implantation process.

implantation system

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 219

microstructure as well as hardness of the substrate surface. It is noted that the microhardness of the surface increased 3-5 times higher than the base metal substrate when increasing the pulse duration. It is also observed that the microhardness of microstructure reduced by decreasing the irradiated energy per unit length of the material where irradiated energy can be reduced by increasing the process travel speed (Hamedi *et al.*, 2011). In a summary, laser carburization is a potential route of strong surface hardening method with a short process time to increase the wear resistance property of the titanium alloys without affecting its bulk properties. This method provides hardest carbide layer compared to other

Ion implantation or ion beam processing is a procedure in which ions of a material are accelerated in an electric field and bombarded into the solid substrate surface. Various ions such as oxygen, nitrogen, carbon, etc. can be implanted on any substrate material for a coating purpose to modify the substrate surface. When carbon is implanted on substrate material then the effect of the surface modification is similar to carburization. Similarly, this method also can be applicable for nitridation as well as oxidation. Two common types of ion implantation process are (i) Conventional beam line ion implantation and (ii) Plasma immersion ion implantation (PIII) method. The basic difference between the beam line ion implantation and plasma immersion ion implantation method is the target function. In beam line ion implantation, the target is totally isolated from the ion beam generation. In contrast, the target is an active part of the ion generation through bias voltage in PIII system (Savaloni *et al.*, 2010). Fig. 4 shows the two typical types of ion implantation systems. The ion implantation phenomena started with the acceleration of ions and it directed towards a substrate (titanium in the present case) which is called target. The energy of the ions is usually in the range of several kilo electronvolt to few mega electronvolt. This level of energy could cause significant changes in the surface by the ions penetration. However, the energy of ions is selected carefully to avoid deep penetration inside the substrate. Therefore, the surface modifications are limited to the near-surface region and a depth of 1 µm from the surface is normal (Rautray *et al.*, 2011). In other words, bulk material properties will not be

(a) (b)

Fig. 4. Schematic diagram of (a) beam line ion implantation system and (b) PIII ion

Fig. 3. Schematic diagram of typical pulsed laser carburization set up.

Investigations on laser carburizing technique were extended from steel to α-Titanium (Fouilland *et al.*, 1997), commercial pure titanium (Courant et al., 2005) , and biomedical grade titanium (Sampedro *et al.*, 2011). Laser melting carburization produces thick coating ranged between one and several hundred micrometers depending on the irradiation conditions. Other carburizing methods are more suitable for producing thin film coating. Another advantage of this method compared to other techniques is that it's capability of coating complex substrate geometry and shape such as notches or grooves where through other methods very difficult to reach these inaccessible areas. Wide heat affected zone is a general issue for thermal or plasma method heating which leading to shape distortion. On the other hand, laser carburizing method is free from these disadvantages since an accurate focused heating on the work piece can be controlled easily. Other commonly controlled laser processing parameters are laser power (W), scanning speed (mm/min), pulse/deposition time (ms), laser frequency (Hz) and overlapping factor (%). The effects of these variables are investigated in terms of changes to the hardness, compositions, heat affected zone, pores, cracks and microstructure of the carburized zone. For instance, a group of researchers investigated the effect of processing time on the TiC microstructure formation on titanium alloy using Nd-YAG laser (Courant *et al.*, 2005). They observed that the time ratio has a significant effect on the carburized microstructure. A lower time ratio caused an increase in pulse power leading to form a thick layer of melted zone with rich in carbon but free from graphite formation. In contrast, higher time ratio produces large amount of graphite formation in the melted zone which can act as a solid lubrication. This phenomenon shows the potential to reduce abrasive wear rate and hence increase the tribological performance of articulation implants. One group of researchers compared the effect of process parameters (laser power and scanning speed) on solidification of TiC microstructure using two different laser sources on Ti-6Al-4V substrate (Saleh *et al.*, 2010). They found that TiC appears either in the form of dendrites or as particles located inside the grains and at the grain boundaries. This resulted significant increment in microhardness of the surface after carburizing process. They concluded that both Nd–YAG and the CO2 lasers able to produce macroscopically homogeneous microstructures of carburized layers. However, the former laser produces deeper carburized layer compared to the later. Recently, another group of workers studied pulse wave laser method (Nd-YAG laser) to form TiC layer on CP-Ti. They investigated the effect of process parameters (irradiated energy per length and pulse duration) on the microstructure as well as hardness of the substrate surface. It is noted that the microhardness of the surface increased 3-5 times higher than the base metal substrate when increasing the pulse duration. It is also observed that the microhardness of microstructure reduced by decreasing the irradiated energy per unit length of the material where irradiated energy can be reduced by increasing the process travel speed (Hamedi *et al.*, 2011). In a summary, laser carburization is a potential route of strong surface hardening method with a short process time to increase the wear resistance property of the titanium alloys without affecting its bulk properties. This method provides hardest carbide layer compared to other two carburization techniques.

#### **6. Ion implantation and deposition**

218 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

Melting bath

Scan direction

Condensation

Investigations on laser carburizing technique were extended from steel to α-Titanium (Fouilland *et al.*, 1997), commercial pure titanium (Courant et al., 2005) , and biomedical grade titanium (Sampedro *et al.*, 2011). Laser melting carburization produces thick coating ranged between one and several hundred micrometers depending on the irradiation conditions. Other carburizing methods are more suitable for producing thin film coating. Another advantage of this method compared to other techniques is that it's capability of coating complex substrate geometry and shape such as notches or grooves where through other methods very difficult to reach these inaccessible areas. Wide heat affected zone is a general issue for thermal or plasma method heating which leading to shape distortion. On the other hand, laser carburizing method is free from these disadvantages since an accurate focused heating on the work piece can be controlled easily. Other commonly controlled laser processing parameters are laser power (W), scanning speed (mm/min), pulse/deposition time (ms), laser frequency (Hz) and overlapping factor (%). The effects of these variables are investigated in terms of changes to the hardness, compositions, heat affected zone, pores, cracks and microstructure of the carburized zone. For instance, a group of researchers investigated the effect of processing time on the TiC microstructure formation on titanium alloy using Nd-YAG laser (Courant *et al.*, 2005). They observed that the time ratio has a significant effect on the carburized microstructure. A lower time ratio caused an increase in pulse power leading to form a thick layer of melted zone with rich in carbon but free from graphite formation. In contrast, higher time ratio produces large amount of graphite formation in the melted zone which can act as a solid lubrication. This phenomenon shows the potential to reduce abrasive wear rate and hence increase the tribological performance of articulation implants. One group of researchers compared the effect of process parameters (laser power and scanning speed) on solidification of TiC microstructure using two different laser sources on Ti-6Al-4V substrate (Saleh *et al.*, 2010). They found that TiC appears either in the form of dendrites or as particles located inside the grains and at the grain boundaries. This resulted significant increment in microhardness of the surface after carburizing process. They concluded that both Nd–YAG and the CO2 lasers able to produce macroscopically homogeneous microstructures of carburized layers. However, the former laser produces deeper carburized layer compared to the later. Recently, another group of workers studied pulse wave laser method (Nd-YAG laser) to form TiC layer on CP-Ti. They investigated the effect of process parameters (irradiated energy per length and pulse duration) on the

Graphite powder Carburized layer

Fig. 3. Schematic diagram of typical pulsed laser carburization set up.

Argon flow

Pulsed Laser beam

Titanium

Ion implantation or ion beam processing is a procedure in which ions of a material are accelerated in an electric field and bombarded into the solid substrate surface. Various ions such as oxygen, nitrogen, carbon, etc. can be implanted on any substrate material for a coating purpose to modify the substrate surface. When carbon is implanted on substrate material then the effect of the surface modification is similar to carburization. Similarly, this method also can be applicable for nitridation as well as oxidation. Two common types of ion implantation process are (i) Conventional beam line ion implantation and (ii) Plasma immersion ion implantation (PIII) method. The basic difference between the beam line ion implantation and plasma immersion ion implantation method is the target function. In beam line ion implantation, the target is totally isolated from the ion beam generation. In contrast, the target is an active part of the ion generation through bias voltage in PIII system (Savaloni *et al.*, 2010). Fig. 4 shows the two typical types of ion implantation systems. The ion implantation phenomena started with the acceleration of ions and it directed towards a substrate (titanium in the present case) which is called target. The energy of the ions is usually in the range of several kilo electronvolt to few mega electronvolt. This level of energy could cause significant changes in the surface by the ions penetration. However, the energy of ions is selected carefully to avoid deep penetration inside the substrate. Therefore, the surface modifications are limited to the near-surface region and a depth of 1 µm from the surface is normal (Rautray *et al.*, 2011). In other words, bulk material properties will not be affected by the ion implantation process.

Fig. 4. Schematic diagram of (a) beam line ion implantation system and (b) PIII ion implantation system

Surface Modification Techniques for Biomedical Grade of

a future prospective research area.

detail while the discussions on other methods are in brief.

issues simultaneously, i.e. tribological and clinical.

**7. Conclusions** 

Titanium Alloys: Oxidation, Carburization and Ion Implantation Processes 221

tribological properties such as hardness as well as corrosion resistance. However, in the recent development, it shows that the research interests in this method have been expanded to include implanting both metal and non-metallic ion simultaneously on titanium alloy. The main driving force of introducing this dual implantation method is to address the clinical and tribological issues concurrently. For example, Ca and Mg ion implanted into titanium alloy for increasing the bone integration (Kang *et al.*, 2011). Ag and N ion have been used to have dual effects on titanium alloy (Li *et al*., 2011). Ag provides antibacterial effect and nitride layer (TiN) formed on the titanium surface increases wear and corrosion resistance. The ion implantation sequence in this dual method also has impact on the deposited particle size and distribution. In a summary, ion implantation method is more suitable for wear and corrosion resistance application. However, recent research trend on ion implantation shows the focus is not only on tribological issues but also on the effect in clinical aspects. Therefore, various metallic ions implantation on titanium alloy appear to be

Various surface modification methods used for improving properties of biomedical grade titanium and its alloys are discussed in this chapter. There are at least six (6) different methods available in the current practice. These are mechanical, chemical, physical, sol-gel, carburization and ion implantation. Oxidation and carburization methods are discussed in

Oxidation method modifies the titanium surface into various types of oxides. The main objective is to produce porous oxide structure for promoting cell growth and cell attachment. There are cases where corrosion and wear resistance are also improved by applying this technique. The recent trend shows that the oxide layer formed on the titanium

Carburization is mainly used to improve wear resistance by increasing titanium surface hardness via thermal, gas and laser melting methods. Hardness of titanium carbide layer formed through these methods varies from 1.5 to 5 times as compared to bare material. Higher hardness of carbide layer assists to increase wear and corrosion resistance of implant surface. Ion implantation method provides better wear and corrosion resistance than other thermal surface modification techniques. In the recent trend, ion implantation technique is found to

Generally, it is observed that the overall trends of surface modification methods seem to shift from the use of conventional source (chemical, induction heater and gas) to the application of advanced technology (electrolyte based, laser, plasma and ion). This could be due to the low efficiency of conventional methods that require longer time and huge amount of energy. The works on surface modifications also appear to expand from focusing on tribological issues such as wear resistance, corrosion resistance and hardness of modified layer to clinical issues such as cell growth, cell attachment and antibacterial effects. These developments demand newer technologies in the future for providing solutions of dual

substrate serves as a basis for growing hydroxyapatite layer to increase bioactivity.

provide dual effects concurrently such as wear resistance and antibacterial effect.

For instance, in carbon ion implantation process, the implanted carbon ions are limited either to form titanium carbides or carbon atoms with C–C bonds near the surface. This may result in improvement of the mechanical properties as well as biocompatibility of titanium alloys. However, studies on the issue regarding corrosion resistance of TiC formation by ion implantation carburization are still underway. It is reported that a very high carbon ion dose of implantation (1018 cm-2) will reduce surface hardness of the titanium substrate (Viviente *et al.*, 1999). The reaction between excess carbon and titanium produces mixed layer of graphite (C–C bonds) and TiC which cause reduction in hardness. In other study, a moderate dose of ion implantation from 5 x 1015 to 1x1017 cm-2 able to create nanocrystalline titanium carbide (TiC) layer which hardness of more than two folds on Ti–6Al–4V alloy substrate (Liu *et al*., 2004). Liu et al. also reported that the tribological properties of titanium alloys are significantly improved at ion implant doses of over 4 x 1017 cm-2, producing friction coefficients of 0.2–0.3. Ion implantation method is free from some disadvantages of plasma process such as thick coating, different phases of mixed crystalline and low crystallinity which leads to delamination problem (Rautray *et al*., 2011). Effects of ion implantation process on wear resistance also have been studied by various researchers. Williams *et al.* investigated carbon ion implantation effect on the wear resistance of Ti–6Al– 4V alloys in a corrosive environment with the composition of 0.9% NaCl or 0.9% NaCl + 10% serum (Williams, 1985). Two-stage carbon ion implantation: 2.5x 1016 cm-2 at 35 kV followed by 1.6 x 1017 cm-2 at 50 kV were carried out for the test. They revealed that ion implanted sample shows reduction in corrosion current by a factor of 100 compared to that untreated samples. A group of researchers investigated Ti–6Al–4V alloy's corrosion resistance after 80 kV, 3 x 1017 cm-2 carbon ion implantation (Zhang *et al*., 1991). They carried out examinations using electrochemical methods in two media: 0.5 M H2SO4 and (HCl + NaCl) solution (pH = 0.1) at 25 oC. In both solutions, ion implanted samples show higher corrosion potential (Ecor) than unimplanted samples. They also reported that the increment in the surface corrosion resistance was due to a durable solid passive layer formation. Other group of researchers experimented various carbon doses on the titanium alloy for evaluating corrosion resistance of TiC formation at energy of 100 keV. in 0.9% NaCl solution at a temperature of 37 °C (Krupa *et al*., 1999). They revealed that the corrosion resistance of titanium alloy improved significantly by producing a continuous solid nanocrystalline TiC layer when applying 1×1017 C+ cm−2 of carbon dose or more. Another group of workers studied the formation of TiC on titanium alloy using PIII method by varying the deposition times (Baba *et al*., 2007). They concluded that the formation of TiC through ion implantation on titanium alloy depends on amount of carbon ion implantation which is proportional to ion implantation process time. Corrosion resistance on biomedical grade titanium alloy can also be improved by nitrogen ion implantation. It was reported that increasing of N+ flux will influence the corrosion potential, corrosion current and passive current. These changes lead to initial increase in the corrosion resistance of the titanium alloy (Savaloni *et al*., 2010). Other group of researchers investigated the effect of process temperature and implantation time on the corrosion properties of Ti-6Al-4V. It was found that prolonged implantation times do not contribute to a major changes in corrosion resistance where process temperature does (Silva *et al*., 2010). They also reported that the best corrosion resistance achieved at 760 oC with 2 hr processing time. Previous studies on PIII method basically focused on single non-metallic ion implantation to improve tribological properties such as hardness as well as corrosion resistance. However, in the recent development, it shows that the research interests in this method have been expanded to include implanting both metal and non-metallic ion simultaneously on titanium alloy. The main driving force of introducing this dual implantation method is to address the clinical and tribological issues concurrently. For example, Ca and Mg ion implanted into titanium alloy for increasing the bone integration (Kang *et al.*, 2011). Ag and N ion have been used to have dual effects on titanium alloy (Li *et al*., 2011). Ag provides antibacterial effect and nitride layer (TiN) formed on the titanium surface increases wear and corrosion resistance. The ion implantation sequence in this dual method also has impact on the deposited particle size and distribution. In a summary, ion implantation method is more suitable for wear and corrosion resistance application. However, recent research trend on ion implantation shows the focus is not only on tribological issues but also on the effect in clinical aspects. Therefore, various metallic ions implantation on titanium alloy appear to be a future prospective research area.

### **7. Conclusions**

220 Titanium Alloys – Towards Achieving Enhanced Properties for Diversified Applications

For instance, in carbon ion implantation process, the implanted carbon ions are limited either to form titanium carbides or carbon atoms with C–C bonds near the surface. This may result in improvement of the mechanical properties as well as biocompatibility of titanium alloys. However, studies on the issue regarding corrosion resistance of TiC formation by ion implantation carburization are still underway. It is reported that a very high carbon ion dose of implantation (1018 cm-2) will reduce surface hardness of the titanium substrate (Viviente *et al.*, 1999). The reaction between excess carbon and titanium produces mixed layer of graphite (C–C bonds) and TiC which cause reduction in hardness. In other study, a moderate dose of ion implantation from 5 x 1015 to 1x1017 cm-2 able to create nanocrystalline titanium carbide (TiC) layer which hardness of more than two folds on Ti–6Al–4V alloy substrate (Liu *et al*., 2004). Liu et al. also reported that the tribological properties of titanium alloys are significantly improved at ion implant doses of over 4 x 1017 cm-2, producing friction coefficients of 0.2–0.3. Ion implantation method is free from some disadvantages of plasma process such as thick coating, different phases of mixed crystalline and low crystallinity which leads to delamination problem (Rautray *et al*., 2011). Effects of ion implantation process on wear resistance also have been studied by various researchers. Williams *et al.* investigated carbon ion implantation effect on the wear resistance of Ti–6Al– 4V alloys in a corrosive environment with the composition of 0.9% NaCl or 0.9% NaCl + 10% serum (Williams, 1985). Two-stage carbon ion implantation: 2.5x 1016 cm-2 at 35 kV followed by 1.6 x 1017 cm-2 at 50 kV were carried out for the test. They revealed that ion implanted sample shows reduction in corrosion current by a factor of 100 compared to that untreated samples. A group of researchers investigated Ti–6Al–4V alloy's corrosion resistance after 80 kV, 3 x 1017 cm-2 carbon ion implantation (Zhang *et al*., 1991). They carried out examinations using electrochemical methods in two media: 0.5 M H2SO4 and (HCl + NaCl) solution (pH = 0.1) at 25 oC. In both solutions, ion implanted samples show higher corrosion potential (Ecor) than unimplanted samples. They also reported that the increment in the surface corrosion resistance was due to a durable solid passive layer formation. Other group of researchers experimented various carbon doses on the titanium alloy for evaluating corrosion resistance of TiC formation at energy of 100 keV. in 0.9% NaCl solution at a temperature of 37 °C (Krupa *et al*., 1999). They revealed that the corrosion resistance of titanium alloy improved significantly by producing a continuous solid nanocrystalline TiC layer when applying 1×1017 C+ cm−2 of carbon dose or more. Another group of workers studied the formation of TiC on titanium alloy using PIII method by varying the deposition times (Baba *et al*., 2007). They concluded that the formation of TiC through ion implantation on titanium alloy depends on amount of carbon ion implantation which is proportional to ion implantation process time. Corrosion resistance on biomedical grade titanium alloy can also be improved by nitrogen ion implantation. It was reported that increasing of N+ flux will influence the corrosion potential, corrosion current and passive current. These changes lead to initial increase in the corrosion resistance of the titanium alloy (Savaloni *et al*., 2010). Other group of researchers investigated the effect of process temperature and implantation time on the corrosion properties of Ti-6Al-4V. It was found that prolonged implantation times do not contribute to a major changes in corrosion resistance where process temperature does (Silva *et al*., 2010). They also reported that the best corrosion resistance achieved at 760 oC with 2 hr processing time. Previous studies on PIII method basically focused on single non-metallic ion implantation to improve

Various surface modification methods used for improving properties of biomedical grade titanium and its alloys are discussed in this chapter. There are at least six (6) different methods available in the current practice. These are mechanical, chemical, physical, sol-gel, carburization and ion implantation. Oxidation and carburization methods are discussed in detail while the discussions on other methods are in brief.

Oxidation method modifies the titanium surface into various types of oxides. The main objective is to produce porous oxide structure for promoting cell growth and cell attachment. There are cases where corrosion and wear resistance are also improved by applying this technique. The recent trend shows that the oxide layer formed on the titanium substrate serves as a basis for growing hydroxyapatite layer to increase bioactivity.

Carburization is mainly used to improve wear resistance by increasing titanium surface hardness via thermal, gas and laser melting methods. Hardness of titanium carbide layer formed through these methods varies from 1.5 to 5 times as compared to bare material. Higher hardness of carbide layer assists to increase wear and corrosion resistance of implant surface.

Ion implantation method provides better wear and corrosion resistance than other thermal surface modification techniques. In the recent trend, ion implantation technique is found to provide dual effects concurrently such as wear resistance and antibacterial effect.

Generally, it is observed that the overall trends of surface modification methods seem to shift from the use of conventional source (chemical, induction heater and gas) to the application of advanced technology (electrolyte based, laser, plasma and ion). This could be due to the low efficiency of conventional methods that require longer time and huge amount of energy. The works on surface modifications also appear to expand from focusing on tribological issues such as wear resistance, corrosion resistance and hardness of modified layer to clinical issues such as cell growth, cell attachment and antibacterial effects. These developments demand newer technologies in the future for providing solutions of dual issues simultaneously, i.e. tribological and clinical.

Surface Modification Techniques for Biomedical Grade of

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#### **8. Acknowledgements**

Authors would like to express highest gratitude to Ministry of Higher Education (MOHE)and Ministry of Science, Technology and Innovation (MOSTI) for providing grant to conduct this study via vote numbers Q.J130000.7124.02H60, 78611 and 79374. Authors also would like to thank the Faculty of Mechanical Engineering, UTM for providing their facilities to carry out this study.

#### **9. Nomenclature**


Nd:YAG Neodymium Yttrium Aluminium Garnet


#### **10. References**


Authors would like to express highest gratitude to Ministry of Higher Education (MOHE)and Ministry of Science, Technology and Innovation (MOSTI) for providing grant to conduct this study via vote numbers Q.J130000.7124.02H60, 78611 and 79374. Authors also would like to thank the Faculty of Mechanical Engineering, UTM for providing their

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BCP Biphasic Calcium Phosphates CVD Chemical Vapour Deposition

HVOF High Velocity Oxygen Fuel spraying

Nd:YAG Neodymium Yttrium Aluminium Garnet Nd:YLF Neodymium Yttrium Lithium Fluoride PIII Plasma Immersion Ion Implantation

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CP Commercial Pure DGUN Detonation Gun

HA Hydroxyapatite LSA Laser Surface Alloying MAO Micro Arc Oxidation

SBF Simulated Body Fluid

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### *Edited by A.K.M. Nurul Amin*

The first section of the book includes the following topics: fusion-based additive manufacturing (AM) processes of titanium alloys and their numerical modelling, mechanism of α-case formation mechanism during investment casting of titanium, genesis of gas-containing defects in cast titanium products. Second section includes topics on behavior of the (α + β) titanium alloys under extreme pressure and temperature conditions, hot and super plasticity of titanium (α + β) alloys and some machinability aspects of titanium alloys in drilling. Finally, the third section includes topics on different surface treatment methods including nanotube-anodic layer formation on two phase titanium alloys in phosphoric acid for biomedical applications, chemico-thermal treatment of titanium alloys applying nitriding process for improving corrosion resistance of titanium alloys.

Photo by VladZymovin / iStock

Titanium Alloys - Towards Achieving Enhanced Properties for Diversified Applications

Titanium Alloys

Towards Achieving Enhanced Properties for

Diversified Applications

*Edited by A.K.M. Nurul Amin*