Titanium-Based Alloys with High-Performance: Design and Development

*Ram Krishna*

## **Abstract**

In recent years, titanium alloys with better properties have become increasingly popular. Their composition must be precisely designed to meet these demands. Screening alloy properties such as corrosion resistance, specific strength, properties to service at high temperatures, and microstructural stability requires a fair amount of effort and money to accomplish. By taking titanium-based alloys as an example, this chapter reviews the use of high-performance alloy design and development approach for industrial applications, in order to simplify the selection of titanium alloy compositions. The different high-throughput alloy design methods have been used by researchers to calculate diffusion coefficients of multiple elements using a thermodynamic database of atomic mobility. A composition with comprehensively optimal properties is selected by applying a rigorous screening criterion and then evaluating it in an experimental setting in order to come up with an optimal composition. Comparing this strategy with the data-driven material design methods that have been developed in recent times, few methods are more accurate and efficient, mainly because the diffusion pairs, the atomic mobility databases, and the refined physical models work together to make this strategy the most accurate and efficient. This approach could help develop high-performance titanium alloys, to overcome challenges of developing titanium alloys.

**Keywords:** titanium alloys, alloy design, high-throughput methods, microstructural stability, high-performance alloy

## **1. Introduction**

Titanium alloys are used extensively for the manufacture of components used in automobile, chemical, aerospace, power generation, and biomedical applications that are subjected to complex operating conditions. It remains a challenge, however, to develop an alloy that has the desired combination of properties at an affordable cost. As an engineering material, they are useful in the manufacturing of turbine engines and aircraft components. The reason for this is that titanium alloys, among others, exhibit excellent strength at high-temperature applications, better creep resistance properties, good high-temperature microstructural stability, and resistance to corrosion and oxidation. The use of titanium alloys is not suitable for all parts of turbine

engines due to phase equilibria and microstructural stability and these conventional titanium alloys cannot withstand operating temperatures greater than 600°C [1].

There is a range of advanced titanium alloys, which may provide titanium alloys with a higher temperature capability, such as TiAl, Ti3Al, and Ti2AlNb, as well as titanium/titanium aluminides [2]. This is a vital issue in terms of ensuring that these engineering materials maintain a high-temperature structural capability in order to be able to meet the increasing requirements for high thrust-to-weight ratios and energy efficiency that are currently being developed for turbine engine applications. These materials must be able to meet the increasing needs of high thrust-to-weight ratios and energy efficiency [3].

Titanium alloys have become more popular over the past few decades due to their ability to alloy with a wide variety of elements, such as Mo, Al, Ta, Zr, V, Zr, Mn, Fe, Ni, Co, Cr, Cu, and Nb [4]. There is no doubt that alloying elements play a crucial role in stabilizing either the low-temperature or high-temperature phases in titanium alloys [3, 5]. Depending upon the alloying elements in varying proportions in an alloy, they often result in the formation of low-temperature and high-temperature phases in the titanium alloy system. These phases included in the Ti-based alloy systems are defined as α, β, near-α, near-β, etc. Therefore, titanium alloys have been able to achieve excellent properties as a result of their chemical compositions [6]. There are three types of base titanium alloys mainly identified as α, β, and α + β alloys according to their phase stability.

Despite this, research has shown that there is a correlation between the overall properties of the alloy and the level of impurities it contains. Impurities adversely affect the plasticity of the alloy. There is a plastic deformation associated with titanium alloys when hydrogen, carbon, oxygen, and nitrogen combine with them [7]. As there are so many possible compositions and it is not possible to screen them in a practical manner using a random combination of these elements, there seems to be a compelling need for new approaches that will enable us to make efficient choices of compositions for the manufacturing of titanium alloys, and other advanced alloys of high performance. The interaction between the elemental composition of titanium alloys, their manufacturing techniques of them, and their microstructure of the alloys must be taken into account when determining the properties of titanium alloys and when they are being designed to get the best results. It is therefore imperative that one utilizes the correlation between the microstructure, properties, and performance of titanium alloys in order to gain a more comprehensive understanding of them. As a result of this correlation, it is expected that it will provide an opportunity to develop novel designs as a result of this correlation [8].

It has been demonstrated in this study that it will be feasible to use a new approach that will give correlations between the evolution of microstructure and properties of alloy systems and the interdiffusion properties of their compositional elements to arrive at a new approach to the problem. Evidence exists that indicates that titanium alloys obtain their strength as a result of the strengthening of the solid solution and grain refinement.

## **2. High-throughput materials characterization techniques**

### **2.1 Mapping spatial data using statistical techniques**

In order to characterize an alloy, the intrinsic heterogeneity of the material is utilized as a basis for a statistical spatial mapping technique that allows high

## *Titanium-Based Alloys with High-Performance: Design and Development DOI: http://dx.doi.org/10.5772/intechopen.108748*

throughput. There are tens of thousands of material microarrays that are being used to obtain different compositions, structures, and properties of a material through the cross-scale characterization of the material. It is necessary to formulate a statistical spatial-mapping model between the two sets of parameters based on the original material.

 With the aid of high-throughput computational studies, it is possible to create databases by screening lattice units, which determine the properties of screened material with the help of high-throughput computations, and then creating a database of materials once the materials have been screened, and after the materials have been screened, the database can be created. There are numerous types of materials-design optimization strategies that have been developed to guide new materials discovery, process optimization, and material modification, as shown in **Figure 1** .

 A Ti-alloy can have slightly different compositions, structures, as well as properties at different points of its structure, and the arrangement of these small differences determines the overall quality of the alloy as a whole, which is determined by the combination of these small differences. A wide range of rapid characterization techniques can be used to gather data from the macroscopic to the microscopic scale for the purpose of high-throughput statistical spatial mapping on a micron level. In order to meet practical sample sizes, fast and reasonable turnaround times are required. This is in order for composition, structure, and property datasets to be gathered at each of the locations. A database can be constructed based on the data entered into it, which contains spatial mapping lattices, based on precisely placed positional coordinates and references to point-to-point correspondence, in order to construct a spatial mapping map. The spatial mapping datasets are selected from a database within the required target intervals based on requirements for material research and development.

 In order to determine a suitable design that is more likely to meet the targeted requirements based on a statistical analysis of the data, a statistical analysis can be performed in order to determine the appropriate design. A number of studies have

 **Figure 1.**

 *It is a statistical spatial mapping technique based on the heterogeneity of the materials to generate the maps [ 9 ].* 

demonstrated that optimizing process parameters allows the assembly of these genetic units at the mesoscale to be verified, and quantitative correlations have been established between the micro-, meso-, and macroscales, as well as between practical samples and across the spectrum composition, structure, and properties. Recent years have seen the use of high-throughput statistical spatial mapping techniques to characterize a variety of material systems, including a wide variety of titanium alloys [10].

An alloy is a material that is heterogeneous, multielementary, and complex in structure. As a result, the structural composition, properties, and structure of an alloy may differ slightly at different points within it, and it is the amalgamation of these variances that determines the global functioning of the alloy. As a building block, a unit cell arrangement is used to provide insight into the properties of a material. The arrangement of unit cells is therefore critical for understanding the material and determining its properties at nanoscale. It is therefore possible to establish a correlation between the microscales, mesoscales, macroscales, and across-scale spans, as well as the compositional information, the structural information, and the properties of those spans, so as to enable the creation of novel materials and the amendment of current materials efficiently and economically. Rapid measurement of compositional information, structural information, and properties related to application at multilocations are performed in order to obtain practical sample sizes based on the available data. With the use of accurate positional coordinates, as well as point-to-point communication, a database is created to represent spatial mappings. Spatial mapping datasets are selected based on the target intervals as part of the design requirement of developing new materials. Based on a variety of factors, a variety of statistical analyses may be used to determine which design is best suited to meet the targeted requirements. Several criteria can be used to determine which design is best suited to meet the intended requirements, such as metrics and models that can be used to determine the frequency of occurrence within the range of parameters, the correlation ratio between parameters, and the statistical elimination of outliers. Many researchers used the process and used advanced microscopy and spectroscopy for data acquisition and statistical distribution analysis [11, 12].

### **2.2 Diffusion-multiple approach**

In order to produce sizeable, multicomponents compositional deviations in alloy system samples through thermal interdiffusion, Zhao developed the diffusion multiple technologies, based on diffusion couples that generate compositional variations in bulk samples through diffusion [13]. Various experimental and analytical tools can be used to analyze diffusion multiples to extract the dependence of structure and properties on components. The application of the novel procedure significantly enhanced both the competence of elemental compositional design as well as the screening of appropriate heat treatment practices in comparison with the traditional methods that use a single alloy model to analyze the advancement rules of properties and microstructural information.

The infusion of a variety of alloying elements into titanium alloys can be investigated to determine how they affect their structure and properties by using combination of different diffusional multiple elements. It is, therefore, important and necessary to use diffusion multiple methods to study titanium alloys in order to achieve the best results [14].

It is possible to investigate kinetics, phase diagrams, and compositional-structuralproperties relationships of alloy systems by using the diffusional multiple approaches,

### *Titanium-Based Alloys with High-Performance: Design and Development DOI: http://dx.doi.org/10.5772/intechopen.108748*

which uses the formation of compositional gradients and phase developed by longterm annealing of the alloy [15]. In order to determine diffusion coefficients and phase diagrams, conventional diffusion pairs and diffusion triples have been used for more than three decades. It has previously been demonstrated that it is possible to determine a number of composition-structure–property relationships by performing localized microscale property measurements on single-phase compositions [16].

Many systems have demonstrated the ability of the diffusion-multiple approach to be used as a tool for determining very complex phase diagrams, and this has been demonstrated for many different systems. In order to compare the phase diagrams of simple and very complex ternary systems, diffusion multiples have been used in place of equilibrated alloys to determine the phase diagrams [17]. In light of the results of this experiment, it can be concluded that the phase diagrams which have been determined from diffusion multiples are of very high accuracy [18].

As a result, a diffusion multiple analysis is a method that can be used to analyze diffusion data and to predict the microstructure and properties of alloy compositions using diffusion data for a variety of alloy compositions using the diffusion data as an input.

## **2.3 Computational thermodynamics using CALPHAD**

CALPHAD is a computational thermodynamics program that can be used to compute and develop phase diagrams. It is commonly used for designing and developing new alloy systems [19]. In order to achieve the desired properties and consequently potential applications, it is important to examine the phase structure and phase equilibrium of the alloy systems. In the strategy of novel and advanced alloy systems, one of the advantages of using CALPHAD over entropy alone is its ability to analyze phase formation using free energy rather than entropy alone. This allows CALPHAD to better understand the function of system enthalpy, as well as the function of the entropy in the design of the alloy system [20].

A phase diagram provides detailed information on microstructural phase information as a function of its compositional information, its temperature information, and its information related to pressure. As such, it serves as a guide when designing and developing new Ti alloys. CALPHAD has been proven to be an effective tool for estimating the phases present in titanium alloy systems based on extensive research using titanium alloys. In spite of this, this method still did not yield enough screened titanium alloys which were able to produce the phases required in the temperature range of interest [21].

As a result of this, the thermodynamic databases in CALPHAD continue to grow as more and more experiments are conducted. Therefore, the accuracy of CALPHAD is also expected to increase as the number of experimental data continues to increase. It is expected that high-throughput CALPHAD simulations will be able to provide more accurate and reliable results for creating and optimizing titanium alloy compositions based on desired alloy properties because of more reliable and accurate simulation studies.

### **2.4 Machine learning and statistical methodology approach**

This approach applies machine learning to a variety of different approaches, such as deep learning. This is a subset of machine learning, which is a subset of what we refer to as artificial intelligence. This is often referred to as artificial intelligence.

This refers to a computer system's capability of learning from the inputs it receives. This allows it to improve itself in order to find out how to do things better in the future. In the field of alloy design and development, artificial intelligence, including deep learning and machine learning algorithms, has emerged as a possible computational solution to overcome the challenges of alloy design and development, as well as control the costs and speed of processes in alloy design and development through artificial intelligence [14].

Recent years have seen a rise in interest in alloy development research centered on this approach. With the use of deep learning and machine learning algorithms, it will be feasible to rapidly transform large quantities of experimental data into usable feature information. With the aid of learning algorithms, it is possible to develop computer models that quickly generate judgment results based on input data. In addition, it is possible to addition, it is conceivable to extract information from alloy systems based on the prior knowledge of the system controllers.

Using high-throughput experiments and algorithms based on machine learning, Zhu et al. [14] have developed a titanium alloy using high-throughput experimental techniques. In the field of titanium-based alloy systems, the artificial neural network techniques, as traditional machine learning methods, have been successfully used for a variety of functions, such as predicting properties such as flow stresses, evolutions of microstructures, mechanical properties and parameters that affect during the processing of titanium alloys [22–24].

It has been discovered that when one technology is combined with machine learning, the screening of alloys becomes more efficient. A machine learning algorithm can be used to envisage the microstructure of an alloy and the anticipated results can be equated with those of the experimental results. Zhu et al. [14] reported the findings of a applied diffusion multiple in combination with machine learning algorithms to formulate a new Ti-based alloy system (Ti-3Al-2Nb-1.2 V-1Zr-1Sn-4Cr-4Mo). After the solution was heat-treated at 750°C for 6 hours and the material was aged at 550°C for 6 hours. In the study, researchers reported that better strength and plasticity could be obtained. The evidence suggests that the globular primary α phases elongated during deformation, while the secondary acicular α phases resist dislocation sliding, therefore, providing both a high degree of plasticity and strength for the alloys, that are subjected to the deformation [14].

## **3. Effect of alloying elements on properties**

Various alloying elements are present in titanium alloys. The role of these alloying elements is to strengthen them either in low-temperature phase or in high-temperature phase, depending on the alloying elements. The alloying elements, in varying proportions, stabilize the close-packed hexagonal alpha (a) phase at low temperatures and the body-centered cubic beta (b) phase at high temperatures. It is their contents that determine the morphology and distribution of these phases. It is known that the alpha phase is a solid solution-strengthened phase that is stabilized by aluminum. This increases tensile and creeps strength. Tin is used in conjunction with aluminum to provide strength without embrittlement. Up to 5% of zirconium increases strength at low to intermediate temperatures. As the oxygen content in the titanium alloy increases, the ductility, toughness, and high-temperature strength of the titanium alloy decreases [25].

*Titanium-Based Alloys with High-Performance: Design and Development DOI: http://dx.doi.org/10.5772/intechopen.108748*

The high-temperature beta phase is stabilized by molybdenum, which increases the short-term strength at high temperatures. As well as being a beta-phase stabilizer, Niobium is also added to improve the stability of the surface at high temperatures. At all temperatures, silicon increases the strength as well as the creep resistance of titanium alloys. The other trace elements, such as chromium, cobalt, and nickel, are not beneficial for creep, and their contents are restricted to less than 0.01 percent [26].

Several studies have been conducted on the biocompatibility of titanium alloys for applications such as biomedical implants containing molybdenum, tantalum, and niobium, and on these bases, the developed alloys are Ti-Mo-Zr-Ta, Ti-12Mo-5Ta, and Ti-Nb-Zr-Mo [27, 28].

### **3.1 Case study on Ti-based alloy**

It is important to realize that Ti-based alloy systems possess high-temperature mechanical properties, making them a very important group of structural materials that can be used in a wide range of strategic applications. These two-phase alloys are used for advanced engineering purposes, incorporating third alloying elements to enhance the ductility and strength, and maintain the properties at elevated temperatures. This is a two-phase lamellar structure consisting of alternate layers of tetragonal (L10) and hexagonal (D019) phases that consist of titanium and alloying elements in alternate layers. It is important to note that the optimum volume fraction for lamellar structure leads to an exceptional level of ductility that is virtually nonexistent in pure alloys. As a result of the process of plastic working, as well as the heat treatment, the microstructure of these alloys can be significantly altered in order to achieve a finely tuned mechanical property as well as fatigue behavior depending on the application [29]. There are a variety of mechanical properties depending on the morphology and the distribution of phases. As a result, the mechanical properties of titanium alloys with two phases are strongly influenced by the morphology of each phase. Many factors can affect the strength of an alloy with a lamellar microstructure; however, the thickness and diameter of the lamellae have the greatest impact [30]. In order to improve the mechanical properties of different alloys, the volume fractions, distribution, and morphology of the different phases play a critical role. Ti, Al, Cr, and Nb make up the elemental composition of the Ti-base alloy with nominal chemical compositions of Ti-40Al-2Cr-2Nb-0.4Y-0.2Zr, which has been used in this particular case. After one-hour heat treatment at 1350°C, the samples are furnace cooled to room temperature.

The engineering stress and strain and true stress–strain diagrams are shown in **Figure 2**. In fact, the true stress and strain values are very high because a smaller cross-sectional area is being used, whose section decreases continuously during elongation. True stress values indicate that, unlike engineering stress–strain values, material becomes stronger as strain is increased, compared to engineering stress– strain values. An alloy's mechanical properties can be greatly affected by the size of the colonies of crystallographically oriented lamellae within the alloy since it is a measure of the effective length of the slip that affects the alloy's mechanical properties. In spite of this, the transition to 'basket weave' microstructures will mean it will be even more challenging to determine the size of colonies as they emerge. Therefore, in order to illustrate the effect of microstructure refinement on mechanical properties, the thickness of lamellae was also taken into account as a quantitative parameter to illustrate the effect [31].

 **Figure 2.**

 *A co mparison of (a) engineering stress-strain curve and (b) true stress-strain curve of as-forged and solution heat-treated Ti-based alloys.* 

 Strain hardening is the property of materials that exhibits this property. As part of the forming process, stain hardening (work hardening) plays an important role. Observing the plot, it was evident that stress rises without showing a drop in yield, as indicated in the figure. It can therefore be concluded from the shape of the true stress–strain curve that a material is prone to fracture before it is prone to yield, based on the shape of the curve.

 The optical micrographs of as-forged and solution heat-treated Ti-alloy are shown in **Figure 3** . The lamellae structure of the annealed sample at 1350°C can be seen as having a random orientation due to the annealing process. This lamellae structure consists of alternate layers of the alloys γ- and α-phases. The solution heat-treated Ti-alloys has shown better property than forged alloys. This is due to the fact that the load-transferring capacity of lamellae is greater than that of duplex grains and near grains. A colony size of 80–100 μm was found to be the maximum size of the lamellae in the colony.

 It can be seen in **Figure 4** that Ti-alloy has an even microstructure in an as-forged condition, which consists of equiaxed grains of γ and α phases and alternate plates of α and γ phases. Depending on the sample's history, the morphology of the grains differs from one sample to another. In the present case, the dislocations are thermal in origin. Several second-phase particles larger than 500 nm are usually found on the grain/interphase boundaries. It is worth noting that there is a wide variation in the

 **Figure 3.**  *Optical micrographs of (a) as-forged, and (b) solution-treated Ti-base alloy.* 

## *Titanium-Based Alloys with High-Performance: Design and Development DOI: http://dx.doi.org/10.5772/intechopen.108748*

grain size in this multiphase microstructure. A few of the grains have a size of less than a micrometer, and there are a few others that are larger.

**Figure 5** shows a dark field TEM micrograph of an alpha grain that is disordered. While the formation of α2 from α is taking place, there are a number of finely ordered domains that are being formed, which are more apparent at the outset.

 The deformation mechanism is also identified in the Ti-alloy. There is a high density of dislocations in the γ-phase, while there are very few dislocations in the α<sup>2</sup> phase. There is a great deal of difficulty in deforming the ordered alpha by dislocation slip. It is caused by the slitting of the dislocations, which causes them to become

### **Figure 4.**

 *Transmission electron microscope (TEM) micrographs of as-forged Ti-alloy in bright field mode showing equiaxed grains of γ and α phases and alternate plates of α and γ phases. The morphology of the phases depends on their history, or at least on the stage of their origin in the evolutionary process, which determines their morphology.* 

### **Figure 5.**

 *A dark field TEM micrograph showing the transformation of α → α2 in disordered α grains showing the transformation in a blown out image. There are a number of the finely ordered domain during the formation of one α to another α2 , which are more evident at first.* 

super-dislocations. Super dislocations, as it is well known, require a greater amount of energy in order to move forward.

There is a deformation of the ordered α2 phase by twinning. The disordered α phases are a high-temperature phase, which when cooled to room temperature decomposes to the ordered α2 + γ at room temperature. A crucial aspect that we have collected in our research is the deformation characteristics of the constituent phases, which has provided us with invaluable information. A comparison should be made between the mechanical properties and microstructures of the selected alloy with a few other alloys that have been identified, and similar tests should be performed on those alloys that have been identified. By performing this comparative study, it will be easier to identify which sample is the best of those that have been tested. In addition, these alloys will also be able to provide an idea as to how to further improve the alloy design in Ti-base alloy systems by adding alloying elements or choosing the process of heat treatment, etc. which will result in better alloys.

## **4. Conclusions**

It is evident from this article that the screening of alloy properties as well as microstructural stability is a substantial undertaking that requires a considerable amount of time and effort. In order to simplify the selection of titanium alloy compositions, a highthroughput-based alloy design approach is used. Different high-throughput methods have been used to calculate diffusion coefficients for a number of different elements, using a database of atomic mobility as a basis for calculating diffusion coefficients.

As a result of applying a rigorous screening criterion and evaluating it in an experimental setting in order to come up with the optimal composition, an optimal composition is selected that has comprehensively optimal properties. As compared to the data-driven materials design methods that have been used in recent years, few methods are more accurate and efficient, mainly because diffusion pairs, atomic mobility databases, and refined physical models work together so as to make this strategy the most accurate and efficient.

This approach is believed to be able to enable the development of high-performance titanium alloys regardless of the composition of the alloy, which is believed to be beneficial in overcoming the challenges that are associated with the development of novel titanium alloys for applications in high-temperature structural applications.

## **Author details**

Ram Krishna Department of Metallurgical and Materials Engineering, National Institute of Technology, Jamshedpur, India

\*Address all correspondence to: krishnamme@gmail.com

© 2022 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

*Titanium-Based Alloys with High-Performance: Design and Development DOI: http://dx.doi.org/10.5772/intechopen.108748*

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## **Chapter 2**

## Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material

*Bipin Kedia and Ilangovan Balasundar*

## **Abstract**

Titanium alloys subjected to suitable thermomechanical processing (TMP) schedules can exhibit superplasticity. Most studies on superplasticity of titanium alloys are directed to sheet materials while studies on bulk materials are rather limited. Bulk Superplastic materials require lower load for forging aeroengine components. It further facilitates forming using non-conventional processes such as superplastic roll forming (SPRF). Multi axial forging (MAF), is employed here to achieve bulk superplasticity by imparting large strain without any concomitant change in external dimension. A comparison between uniaxial and MAF with respect to strain, strain path, initial microstructure and heat treatment was carried out to ascertain the microstructure refinement in Ti-6Al-4V alloy. A fine-grained structure was obtained after 3 cycles of MAF followed by static recrystallization at 850°C. Grain boundary sliding was observed in identified processing domain along with strain rate sensitivity (SRS) of 0.46 and maximum elongation of 815%. Validation of established ther¬momechanical sequence on a scaled-up work piece exhibited 640% elongation in domain (T = 820°C, ε ̇= 3 x 10-4/s) which indicated that the established TMP scheme can be used on a reliable and repeatable basis to achieve superplasticity in bulk material.

**Keywords:** titanium, Ti-6Al-4V, superplasticity, multi-axial forging, severe plastic deformation

## **1. Introduction**

Titanium alloys are used extensively in the aerospace industry owing to their high specific strength, good static and dynamic properties, corrosion resistance, etc. [1, 2]. As an engine stator and rotor, they are used extensively in the compressor section as rings, discs and blades. A typical aeroengine compressor (low and high pressure) disc varies from 15 to 100 cm in diameter. These compressor discs are generally manufactured using either conventional or advanced forging techniques such as isothermal or near isothermal forging. Near α and α+β titanium alloys that are used in the compressor region of aeroengine exhibit a martensitic structure with coarse grains in the as-cast and homogenised condition [2]. By subjecting the material to thermomechanical processing (TMP), the lamellar structure can be converted into a fine-grained equiaxed structure. Numerous studies have been carried out to identify TMP parameters to break lamellar structure, and various theories have been put forward to describe the continuous dynamic recrystallisation (CDRX) or globularisation of α lamellae [1, 2]. It has been suggested that the formation of α/α interface (due to sub grain or shear band formation) in contact with β platelet gives rise to surface-driven penetration of the β phase into the lamellae leading to the breaking of lamellae into equiaxed structure. Thin lamellae structure has been reported to show better globularisation kinetics [3] due to easy penetration of β phase into α lamellae. It has been reported that high strain (ε ≥ 3) is necessary to obtain a completely globularised structure [4] which could exhibit superplasticity.

Numerous studies have been carried out towards achieving superplasticity in various titanium alloy sheets [5–7]. However, work reported on realising superplasticity in the bulk material is very few. Multi-axial forging (MAF) was used to achieve submicron grain size in titanium alloy Ti-6Al-4V by Zherebstov et al*.* [8]. It was reported by Salishev et al. [4] that strain path change inherent in MAF process aids in increasing the globularisation kinetics or CDRX as multiple slip systems are activated during the strain path change. Poths et al. [9] subjected Ti-6Al-4V to monotonic and cyclic torsion in order to understand the effect of strain path on the globularisation kinetics of α lamellae. Based on the study and in contrast to the findings of Salishev et al. [4], Poths et al. [9] reported a decrease in kinetics of α globularisation and attributed the same to the change in strain path. Since contrasting results have been reported in literature and α lamellae globularisation is essential to achieve superplasticity in titanium alloys, it is imperative to carry out a systematic investigation on the influence of various factors that affect the globularisation kinetics of α lamellae. Further, as the objective is to achieve superplasticity in the bulk material, it is required to achieve this without modifying the external dimension of the material that facilitates subsequent secondary processing to produce the desired product/component.

A systematic study on the effect of strain, strain path, deformation temperature and starting or initial microstructure of Ti-6Al-4V on the globularisation or CRDX kinetics in a work horse titanium alloy Ti-6Al-4V is presented here. Further, a suitable thermomechanical process scheme that can maximise globularisation in the material on a reliable and repeatable basis is presented along with the temperature-strain rate regime under which the globularised material exhibits superplasticity.

## **2. Material**

Ti-6Al-4V is an α+β alloy designed to provide moderately high strength, good fatigue strength and reasonable fracture toughness up to a temperature of 350°C. For aeroengine applications, the alloy is produced by vacuum arc remelting followed by thermomechanical processing in order to improve the structural integrity of the material. For the current study, triple vacuum arc remelted (VAR) titanium alloy Ti-6Al-4V ingot subjected to primary processing in the β and α+β field followed by mill annealing at 700°C for 1 h was procured from M/s M/s Mishra Dhatu Nigam, Hyderabad, India. The β transus of the 15 cm cylindrical mill annealed bars was reported to be 995 ± 5°C, and the same was reconfirmed through heat treatment experiments. Two billets of 15 cm diameter and 5 cm length were cut from the as-received mill annealed material. The billets extracted were coated with glass coating to prevent oxidation

*Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

**Figure 1.**

*Typical microstructure of Ti-6Al-4V subjected to heat treatment at 1015°C for 60 min followed by (a) water quenching – martensitic structure and (b) air cooling – lamellar structure.*

during high-temperature exposure and were subjected to heat treatment above the β transus at a temperature of 1015°C. The billets were held at this temperature for a period of 60 minutes to achieve thermal homogeneity. One billet was then removed from the furnace and quenched in water while the second billet was cooled in air to obtain martensitic and lamellar microstructures respectively. Typical optical microstructures obtained after water quenching and air cooling are shown in **Figure 1a** and **b** respectively. The prior β grain size was estimated to be 601 + 86 μm and 684 + 64 μm for the water quenched and air-cooled material, respectively.

## **3. Effect of strain, strain path and microstructure**

To evaluate effect of strain, strain path and initial microstructure on the globularisation kinetics of titanium alloy Ti-6Al-4V, both isothermal hot compression (uniaxial monotonic) and multi-axial (non-monotonic) compression experiments were carried out.

### **3.1 Isothermal hot compression**

Cylindrical compression test samples with a constant height-to-diameter ratio of 1.5 were extracted from the water-quenched and air-cooled material using a wire-cut electro discharge machine (EDM) and lathe machine. The edges of the samples were chamfered to avoid fold formation during initial stages of deformation. The prepared samples were coated with glass (Deltaglaze 347) which acts as an oxidation resistor and lubricant. The coated samples were then heated to the desired test temperature and held at that temperature for 30 min in order to achieve thermal homogeneity. The cylindrical samples were then subjected to isothermal hot compression or uniaxial monotonic deformation (height reduction). A deformation of 25%, 40%, 58% and 78% was imparted to the samples at a constant true strain rate of 10−3/s using a computer-controlled 200kN formability and workability testing machine custom built by M/s BISS, Bangalore. The deformation imparted corresponds to an equivalent strain of 0.29, 0.58, 0.87 and 1.51, respectively. After deformation, all the samples were quenched in water and were cut parallel to the deformation direction, hot-mounted and subjected

to metallographic investigation following standard procedures. For stereological measurements of α fraction globularised, automatic and semi-automatic procedures were used and the α lamellae with an aspect ratio of <2.0 was considered to be globularised as reported in various literatures [4, 9]. Standard statistical measures such as relative accuracy (RA) and 95% confidence level [10, 11] were used to ensure reliability of stereology measurements. The microstructure of Ti-6Al-4V with an initial lamellar structure subjected to increasing amount of strain through isothermal hot compression, i.e. uniaxial monotonic deformation is shown in **Figure 2a**–**d**. Deformation of α lamellae and increased globularisation of α lamellae with increasing strain can be readily observed from **Figure 2**. Maximum globularisation was observed in the sample subjected to 78% deformation (**Figure 2d**) that corresponds to a strain of 1.51.

Conversion of lamellar structure of α phase into a globular morphology during deformation is considered to be a recrystallisation process, namely CDRX as against the discontinuous dynamic recrystallisation (DDRX) which has a distinctive nucleation and growth stage [12–14]. Globularisation of α lamellae present in the colony and at the grain boundary has been reported to take place by either sub grain or shear band formation. It was proposed by Margolin and Cohen [15] that subgrains form within the α lamellae during deformation followed by penetration of β phase into the α/α boundary with a simultaneous rotation of boundaries against each other resulting in coarsening of the recrystallised α when compared with the lamellae thickness from

### **Figure 2.**

*Ti-6Al-4V with martensitic microstructure subjected to (a) 25 (b) 40 (c) 58 and (d) 78 percentage reduction or deformation at 900°C with a strain rate of 10−3/s.*

## *Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

which it originated. Weiss et al. [3] reported formation of shear bands as main reason for CDRX of α lamellae. Irrespective of the mechanism, the formation of α/α interface in contact with β induces surface tension driven penetration of β phase resulting in globularisation. Balasundar [16] reported that both these mechanisms, namely sub grain and shear band formation operate in titanium alloys depending on the orientation of α lamellae and the processing conditions.

Based on the orientation of α lamellae with respect to the compression direction, bending and/or kinking of α lamellae can also be readily observed at different locations of the deformed sample. HCP α lamellae that have their c-axis aligned to the deformation direction require very high load or stress to deform [17] because such lamellae have a low Schmid factor and shear force along the slip direction. Though the globularisation fraction of α lamellae increases with increasing deformation, isolated regions where the α lamellae are still intact can be observed in the material even after imparting a stain of 151%. Though globularisation of grain boundary α lamellae could be observed after a stain of 58%, the prior β grain boundaries could be readily observed in the deformed material.

Microstructure of Ti-6Al-4V with an initial martensitic structure subjected to increasing amount of uniaxial monotonic deformation through isothermal hot compression is shown in **Figure 3a**–**d**. The influence of strain on the material with martensitic starting microstructure was observed to be similar to that in lamellar microstructure described above. However, the prior β grains were found to be destroyed completely by deformation, and it was not possible to identify them in the material. A quantitative discussion on the α fraction globularised is presented in Section 3.3.

### **Figure 3.**

*Ti-6Al-4V with lamellar microstructure subjected to (a) 25 (b) 40 (c) 58 and (d) 78 percentage reduction or deformation at 900°C with a strain rate of 10−3/s.*

## **3.2 Multi-axial deformation**

Conventional deformation process, such as forging, rolling, extrusion, etc., alters the dimension of the material that is being deformed. Severe plastic deformation is a metal working technique in which very high plastic strain can be imparted to the material without any concomitant change in the geometry or dimension. As the work piece material geometry is not altered, it provides an opportunity to deform the material repeatedly till the desired amount of strain is imparted. This large plastic strain results in the formation of ultrafine grain structure in the material. A large number of SPD processes are in vogue, MAF is one suchtechnique which is easier to implement as no special die or tooling is required for deforming the material. In MAF, the material is deformed in cyclic way along all the three orthogonal directions. Each MAF cycle consist of three deformation steps, i.e. imparting equal amount of deformation along the three orthogonal directions (X, Y, Z) as shown in **Figure 4**.

The strain path is altered when the sample is rotated by 90° during MAF. This change in strain path is expected to assist in refining the grain size. After three processing or deformation steps, i.e. after a cycle of MAF, the work piece reverts back to its original dimension. Since the dimension of the work piece remains the same, it is possible to carry out MAF cycles multiple times and thereby impart a large amount of strain to the material as desired or till the failure of material. The amount of strain imparted for a given percentage of reduction and the number of MAF cycles can be calculated as per the following relation:

$$
\varepsilon\_{eq} = \mathbf{3N} \ln(\mathbf{1} - \mathbf{R}) \tag{1}
$$

where R is the amount of reduction imparted per direction which is generally assumed to be the same along all the three orthogonal directions, and N is the number of MAF cycles.

*Typical multi-axial forging cycle for imparting 40% deformation along all the three directions (all dimensions in mm).*

## *Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

Two numbers of cubic specimens with dimension of 2.5 x 2.5 x 2.5 cm3 were prepared from the water-quenched and air-cooled Ti-6Al-4V material respectively for carrying out MAF experiments at 900°C with a constant true strain rate of 10−3/s. The first sample from the water-quenched and air-cooled material was subjected to 25% reduction along all the three directions while the second sample was imparted with 40% reduction which corresponds to a cumulative strain of 87% for the sample subjected to 25% reduction and 153% for the 40% reduction sample. While carrying of MAF experiment, the sample was first deformed to the required amount in a particular direction (e.g. X direction), the furnace was opened, and then the sample was rotated and positioned for deforming along the next direction (e.g. Y direction). The furnace was closed and sample was reheated to the desired temperature and held at this temperature for 30 min before imparting the desired amount of deformation along this direction (e.g. Y direction). The procedure was repeated for deforming along the third direction (e.g. Z direction). After completion of desired number of MAF cycle, i.e. after imparting desired amount of reduction along all the three directions as shown in **Figure 4**, the samples were water-quenched to free the microstructure of the material.

Microstructure of Ti-6Al-4V with an initial lamellar structure subjected to a deformation of 25% and 40% along all the three orthogonal direction using MAF process is shown in **Figure 5a** and **b** respectively. No major microstructural change is observed in the material subject to 25% when compared with the initial starting microstructure expect for coarsening of α lamellae and few isolated regions of α globularisation within the grain and at grain boundaries. Partial globularisation of α can be observed at all the three faces of the same subjected to 40% deformation through MAF. Though isolated globularisation of grain boundary α was observed, the prior β grains were found to be intact and distinct.

The microstructures of Ti-6Al-4V with martensitic structure subjected to 25% and 40% reduction using MAF process are shown in **Figure 6a** and **b** respectively. Similar to that of the air-cooled structure, martensitic structure also shows increasing

### **Figure 5.**

*Microstructure of Ti-6Al-4V with lamellar microstructure subjected to (a) 25% and (b) 40% deformation along all the three orthogonal directions through MAF at 900°C with a strain rate of 10−3/s.*

*Titanium Alloys - Recent Progress in Design, Processing, Characterization, and Applications*

### **Figure 6.**

*Microstructure of Ti-6Al-4V with martensitic microstructure subjected to (a) 25 and (b) 40% deformation per direction through MAF process at 900°C with a strain rate of 10−3/s.*

globularisation with increasing amount of deformation. However, the breaking of the prior β grain boundaries and globularisation can be observed to be more in martensitic structure when compared with the lamellar structure.

### **3.3 Globularisation fraction**

The globular α volume fraction was estimated in Ti-6Al-4V with martensitic and lamellar starting structure subjected to uniaxial (monotonic, i.e. no change in strain path) and multi axial (non-monotonic, i.e. changing strain path as the sample is rotated) deformation and is shown in **Figure 7**. It can be seen that, irrespective of the deformation process whether it is monotonic (uniaxial) or non-monotonic (multi-axial), the volume fraction of globularised α increases with increasing strain. The rate of increase in globularised α volume fraction is observed to be high for Ti-6Al-4V with martensitic structure when compared with the material with lamellar structure. This observation here concurs with the report of Shell et al. [18] where they have attributed the increased kinetics of globularisation to α lamellae thickness. When the α lamellae is thin, it is relatively easier for the β phase to penetrate the lamellae and transverse boundary when compared with thick lamellae. It can be noted that for an equivalent strain of 0.87 which corresponds to 58% reduction through uniaxial or monotonic deformation and 25% deformation along all the three directions through multi-axial or nonmonotonic deformation, altering the strain path through multi-axial deformation results in reduced globularisation. However, with increasing equivalent strain from 0.87 to 1.51, i.e. increasing the deformation from 25 to 40% per direction in multi-axial deformation, the fraction of α phase globularised is quite comparable to that obtained through uniaxial deformation (with 78% reduction) as shown in **Figure 7**.

It was reported by Banerjee et al. [19] that when Ti alloys are subjected to thermomechanical processing in the α/β regime, an equiaxed microstructure can *Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

**Figure 7.** *Effect of strain and strain path on globularisation of α lamellae in Ti-6Al-4V.*

be obtained only if the amount of deformation is greater than 30%. As stated earlier, the CRDX or globularisation of α lamellae depends on the formation of α/α interface by shearing, subgrain, buckling and kinking, etc. and the penetration of β phase. As the formation of α/α interface is a result of two contending processes of dislocation accumulation and annihilation, it is proposed that altering strain path before attaining a limiting critical strain leads to large annihilation of dislocations and a possible reduction or disappearance of the substructure. In the material subjected to 25% deformation through MAF, the strain path is changed before reaching the critical amount of deformation (~30%), and the fraction of globularisation is less due to large annihilation of dislocations. With increasing strain from 87 to 151%, adequate strain is available for accumulation of dislocations and formation of stable substructure. Hence, altering the strain path does not influence the fraction of globularisation. From **Figure 7**, it can be readily inferred that 40% deformation through non-monotonic MAF does not cause significant reduction in the globularisation kinetics when compared with uniaxial deformation. Concomitantly, no significant improvement has been observed as reported by Salischev et al. [4].

On the basis of above study, it can be readily inferred that a martensitic structure exhibits better globularisation kinetics in comparison to lamellar structure. Altering the strain path through non-monotonic multi-axial deformation below a critical limiting strain results in reduced globularisation, whereas beyond the critical strain, the results are quite comparable. The critical deformation limit has been identified to be ≥40% through the current investigation. Though required globularisation can be achieved by uniaxial deformation, it results in altering the dimension of the work piece, whereas multi-axial deformation process which does not cause any

concomitant change in the external dimension [20]. Further, unlike uniaxial deformation which produces a dead metal zone at the fiction hill, no such un-deformed region is observed in multi-axial deformation; hence, multi-axial deformation process with 40% reduction per direction is best suited to globularise the α lamellae and refine the microstructure to obtain superplasticity.

## **4. Thermomechanical processing scheme to achieve ultrafine grain structure**

To obtain grain refinement and achieve superplasticity, Ti-6Al-4V with a martensitic starting microstructure was subjected to three cycles of multi-axial deformation with 40% reduction per direction by progressively decreasing the deformation temperature at each cycle from 850, 800 and 750°C, i.e. first cycle of MAF is carried out at 850°C, second cycle of MAF on the sample is carried out at 800°C and so on. Microstructures of Ti-6Al-4V obtained after two and three cycles of multiaxial deformation are shown in **Figure 8a** and **b** respectively. It can be seen that three cycles of multi-axial deformation have resulted in complete globularisation of α lamellae.

The mechanism of globularisation or CDRX of α lamellae depends on formation of α/α boundaries which is followed by penetration of β phase into α lamellae to separate the boundary [15]. Formation of stable substructure (α/α boundaries) or low-angle grain boundaries (LAGB) depends on interaction and multiplication of dislocation in the lamellar structure. When the HCP α lamellae c-axis is oriented parallel to the deformation direction, high shear stresses are required for the operation of basal and prism slip which is the major source of slip in HCP structure [17]. Such unfavourable orientated lamellae exhibit no major deformation leading to low dislocation density and less α/α boundaries which reduces the globularisation kinetics. However, rotation of specimen by 90o during subsequent steps of deformation during multi-axial deformation leads to activation of prism slip system in such lamellae due to an increase in value of Schmid factor. Activation of prism slip system initiates the process of slip, and with increasing

**Figure 8.** *Microstructure of Ti-6Al-4V after (a) two cycles and (b) three cycles of MAF.*

## *Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

strain and grain rotation, the basal slip system also starts operating leading to formation of substructure, which eventually results in dynamic recrystallised or globularised equiaxed α grain. A comparison of the microstructures shown in **Figure 6b**, **8a** and **b** reveals that with increasing the number of MAF cycle, globularisation increases. It can also be seen that the isolated α lamellae visible in Ti-6Al-4V specimen after two cycles of multi-axial deformation are converted into equiaxed grains after three cycles. Further, prior β grain boundaries were completely destroyed, and they are no longer distinguishable after three MAF cycles.

To gain further insight on the microstructure evolution, the material subjected to three cycles of multi-axial deformation was subjected to electron back scattered diffraction (EBSD) characterisation. The low-angle (LAGB; red line) and highangle (HAGB; black line) boundary present in the material was estimated to be 24% and 76% respectively. The observed fraction of HAGB and LAGB is consistent with reported literature which states the presence of around 25–30% of LAGB in heavily deformed structures [20]. The average size of HAGB was estimated to be 1.65 ± 0.6 μm. Further, the inverse pole figure (IPF map) shown in **Figure 9(b)** reveals a random texture with no preferred orientation.

Superplasticity occurs by grain boundary sliding accommodated by diffusion at grain boundaries and lattice [21, 22]. A material with high fraction of HAGB will have higher degree of disorder. As diffusion occurs down the potential gradient, large fraction of HAGB will lead to higher gradient which increases the diffusion rate. It has been reported in the literature that the presence of high fraction of HAGB with maximum fraction lying in the range of 30–60° is necessary for ease of grain boundary sliding [4]. Further, a fairly uniform grain structure with grains of similar size is essential for superplastic forming. The presence of large grains and small grains (mixed or bimodal grains) in same microstructure has a negative effect on superplasticity. It has been reported that kinetics of diffusion is very slow around the large grain in comparison to the kinetics around smaller grain structure [23]. In the absence of slow diffusion around large grains, chances of cavitation are higher around the larger grains leading to early failure without appreciable tensile elongation. It is also well known that higher volume fraction

### **Figure 9.**

*Microstructure of Ti-6Al-4V after three cycles of MAF(a) band contrast image highlighting HAGB (black line) & LAGB (red line) (b)I PF colour map obtained through EBSD.*

### **Figure 10.**

*Microstructure of Ti-6Al-4V after three cycles of MAF and heat treatment at 850°C for 2 h followed by air cooling (a) band contrast image highlighting HAGB (black line) & LAGB (red line) (b) IPF colour map (c) misorientation plot obtained through EBSD.*

of second phase improves superplastic property due to easier α/β grain boundary sliding [24]. Further, in titanium alloys, the diffusivity in BCC β grain is about of two orders of magnitude higher than α alloys [1, 2, 20]. Higher diffusivity of β phase improves superplastic property of titanium alloy. However, at the same time, β phase grows much faster than α phase, and hence, α phase helps in pinning the grain boundary of β and does not allow the grains to grow. Ideally, presence of both phases is essential for easier grain sliding and achieving optimal superplastic property [25].

As the material after three cycles of multi-axial deformation exhibits nearly 24% low-angle boundaries and predominantly α phase, it is essential to convert this microstructure into a one that can exhibit superplasticity. To improve microstructure and to obtain uniform grain size distribution with increased volume fraction of β phase and higher fraction of high-angle grain boundary, annealing of the deformed material was carried out at 850°C for 2 h followed by air cooling. A higher temperature was used for annealing to obtain the desired α and β proportions in the material. From the EBSD band contrast image shown in **Figure 10a**, the β phase present in the heat treated material was estimated to be ~15%. The fraction of HAGB in the heat-treated material was estimated to be 95% with an average grain size of 2.53 ± 0.65 μm. The IPF map shown in **Figure 10b** clearly indicates the random orientation of the grains. A random texture is important for a structure to exhibit superplastic behaviour since cavitation may occur along the transverse direction during deformation due to strain incompatibility [21]. The misorientation profile of the grains shows significant improvement in HAB with a high fraction of grains boundary in the range of 30–60°. The resultant microstructure satisfies the condition of superplasticity as a major mode of deformation.

## **5. Superplastic domain identification and validation**

In order to identify the temperature-strain rate domain under which the material exhibits superplasticity, the SRS of the material subjected to three cycles of MAF and heat treatment was evaluated by carrying out isothermal hot compression tests over a range of temperature and constant true strain rate. Using the flow curves, the SRS of the material was estimated using standard relations [26] and plotted as a function of temperature and strain rate as shown in **Figure 11**. It can be seen

*Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

### **Figure 11.**

*(a) Strain rate sensitivity map at ε = 0.5 (b) initial tensile specimen (gage length = 6 mm, diameter = 4 mm) (c) sample after tensile deformation at 820°C with a strain rate of 10−3/s strain rate and (d) sample after tensile deformation at 820°C with a strain rate of 3×10−4/s strain rate.*

that the material exhibits a maximum SRS of 0.46 between 810°C and 825°C for a constant strain rate of 3×10−4/s.

In order to validate the domain identified, cylindrical tensile samples of 0.6 cm gage length and 0.4 cm diameter (**Figure 11b**) were prepared from the material subjected to three cycles of MAF and heat treatment. Tensile tests were carried out at 820°C with a strain rate of 10−3 and 3×10−4/s. A maximum elongation of 815% was obtained in the sample tested with a strain rate of 3×10−4/s (**Figure 11d**). As the strain rate is increased to 10−3/s, the % elongation obtained decreases from 815 to 447% as shown in **Figure 11c**. During tensile deformation, necking starts in the weak part of the structure and further deformation gets concentrated at this necked region. A tri-axial state-of-stress exists in the neck region and strain rate in this region does not follow the strain rate of the specimen. The strain rate in the region depends on the rate of decrease of the area which is given by [26]:

$$\frac{dA}{dt} = \left(\frac{P}{C}\right)^{\frac{1}{m}} \left(\frac{1}{A^{(1-m)/m}}\right) \tag{2}$$

As the rate of decrease in area is inversely proportional to strain rate sensitivity (m), a higher m value leads to slower strain rate in the necked region. So a higher elongation is observed during tensile deformation.

In order to further substantiate the findings, EBSD characterisation of the material subjected to compression test (after three cycles of multi-axial deformation and heat treatment) at 800°C with a strain rate of 3×10−4/s was carried out to evaluate the grain size, boundary fractions, etc. From the band contrast image shown in **Figure 12**, the fraction of HAB and the average grain size were estimated to be 83% and 2.64 ± 0.89 μm respectively. Comparing the microstructure

**Figure 12.**

*Band contrast image of fine grained Ti-6Al-4V after hot compression at 800°C and strain rate of 0.0003 S−1 highlighted with HAGBs (black), LAGBs (red).*

obtained in the material after deformation in the superplasticity domain and that obtained after three cycles of multi-axial deformation and heat treatment (**Figure 10a**) clearly indicates no major change in grain size. Therefore, from the observations on % elongation and microstructural features, it can be confirmed that the heat-treated material after subjecting to three cycles of multi-axial deformation exhibits superplasticity between 800 and 840°C when deformed with a strain rate of 3×10−4/s

## **6. Scaling-up**

In order to ensure viability of the identified thermomechanical scheme for industrial scale processing, a large-size billet of 15 cm × 12 cm × 9 cm was prepared and subjected to MAF using 2000MT hydraulic forge press. **Figure 13** compares the size of scaled up Ti-6Al-4V billet with that of smaller-size specimen. The microstructure evolution in the large billet after each cycle of MAF is shown in **Figure 14**. It can be clearly seen that the globularisation fraction increases with increasing cycle and a completely globularised structure is achieved after three cycles of MAF. Postdeformation heat treatment for the material was carried out at 850 and 900°C for 2 h followed by air cooling, and the resulting microstructures are shown in **Figure 15**. Heat treatment led to increase in the volume fraction of beta phase with slight coarsening in the grain size.

In order to validate super plasticity, standard cylindrical tensile samples of 2.0 cm gage length and 0.4 cm diameter were prepared and subjected to tensile testing. Tensile tests were carried out at 820°C with a Strain rate of 3 x 10−4 and 10−3/s. Tensile test was carried out for as deformed (three cycles of MAF) specimen also. It can be seen from **Figure 16** that a maximum elongation of 640% has been obtained for the material.

The result is comparable with the one achieved during tensile test of smaller-size MAF processed specimen. The obtained results have been also compared with available literature as shown in **Table 1**. It can be seen that obtained m values and elongation values are comparable with the data available in literature. Elongation obtained during tensile test of specimen extracted from smaller-size MAF specimen and large-size MAF billets clearly indicates superplastic behaviour of three cycles of MAF *Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

**Figure 13.**

*Typical microstructure of large size Ti-6Al-4V after (a,d) first MAF cycle (b,e) second MAF cycle (c, f) third MAF cycle.*

Ti-6Al-4V alloy. Hence, it can be concluded that the TMP scheme established using smaller sample is repeatable and is validated by repeating the process using 2000 T hydraulic Forge press.

#### **Figure 14.**

 *Microstructure of Ti-6Al-4V obtained after third MAF cycle followed by heat treatment at (a) 850°C/2 h/ac and (b) 900°C/2 h/ac.* 

#### **Figure 15.**

 *Comparison of scaled-up Ti-6Al-4V billet and small size specimen.* 

### **Figure 16.**

 *Tensile elongation of Ti-6Al-4V subjected to three cycles of MAF followed by heat treatment at 850°C/2 h/ac and (a) tested at έ = 0.0003/s, T = 820°C (b) tested at έ = 0.001/s, T = 820°C (c) Tensile elongation of Ti-6Al-4V subjected to three cycles of MAF tested at έ = 0.001/s, T = 820°C.* 


*Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

### **Table 1.**

*Comparison of achieved superplastic property with available literature.*

## **7. Conclusions**

From the study it can be concluded that three cycles of MAF with 40% reduction per direction (cumulative effective strain of ~4.6) leads to complete globularisation of martensitic structure with an average α grain size of 1.65 ± 0.6 μm. Annealing of the deformed material at 850°C increases HAB fraction and β phase volume fraction but with a marginal increase in grain size (2.53 ± 0.65 μm). The MAF + annealed material exhibits a maximum SRS of 0.46 when deformed between 810°C and 825°C with constant strain rate of 3 x 10−4/s. A maximum tensile elongation of 815% was obtained with strain rate of 3 x 10−4/s at 820°C. TMP designed was implemented on a large-size work piece under near isothermal condition, and the process was found to be reliable and repeatable to obtain superplasticity in bulk Ti-6Al-4V in large-size specimen for aeroengine applications.

## **Acknowledgements**

The funding provided by Defence Research and Development Organisation is acknowledged.

## **Conflict of interest**

The authors declare no conflict of interest.

## **Author details**

Bipin Kedia and Ilangovan Balasundar\* Near Net Shape Group, Directorate of Advanced Materials and Manufacturing Processes, Defence Metallurgical Research Laboratory, Hyderabad, India

\*Address all correspondence to: i-balasundar.dmrl@gov.in

© 2022 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

*Titanium Alloys: Thermomechanical Process Design to Achieve Superplasticity in Bulk Material DOI: http://dx.doi.org/10.5772/intechopen.108463*

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Section 2
