from a coaxial injector liquid-gas, Thesis of the central School of Paris, 1995. **Part 5**

**Materials and Fabrication** 

290 Advances in Gas Turbine Technology

Zamuner, B., 1995, Experimental and numerical Study of a spray in combustion resulting

**13** 

*India* 

*VIT University* 

**Materials for Gas Turbines – An Overview** 

Advancements made in the field of materials have contributed in a major way in building gas turbine engines with higher power ratings and efficiency levels. Improvements in design of the gas turbine engines over the years have importantly been due to development of materials with enhanced performance levels. Gas turbines have been widely utilized in aircraft engines as well as for land based applications importantly for power generation. Advancements in gas turbine materials have always played a prime role – higher the capability of the materials to withstand elevated temperature service, more the engine efficiency; materials with high elevated temperature strength to weight ratio help in weight reduction. A wide spectrum of high performance materials - special steels, titanium alloys and superalloys - is used for construction of gas turbines. Manufacture of these materials often involves advanced processing techniques. Other material groups like ceramics, composites and inter-metallics have been the focus of intense research and development; aim is to exploit the superior features of these materials for improving the performance of

The materials developed at the first instance for gas turbine engine applications had high temperature tensile strength as the prime requirement. This requirement quickly changed as operating temperatures rose. Stress rupture life and then creep properties became important. In the subsequent years of development, low cycle fatigue (LCF) life became another important parameter. Many of the components in the aero engines are subjected to fatigue- and /or creep-loading, and the choice of material is then based on the capability of

Coating technology has become an integral part of manufacture of gas turbine engine components operating at high temperatures, as this is the only way a combination of high level of mechanical properties and excellent resistance to oxidation / hot corrosion

The review brings out a detailed analysis of the advanced materials and processes that have come to stay in the production of various components in gas turbine engines. While there are thousands of components that go into a gas turbine engine, the emphasis here has been on the main components, which are critical to the performance of the engine. The review also takes stock of the R&D activity currently in progress to develop higher performance materials for gas turbine engine application. On design aspects of gas turbine engines, the reader is referred to the latest edition of the Gas Turbine Engineering Handbook (Boyce,

**1. Introduction** 

gas turbine engines.

the material to withstand such loads.

resistance could be achieved.

2006).

Nageswara Rao Muktinutalapati

## **Materials for Gas Turbines – An Overview**

Nageswara Rao Muktinutalapati *VIT University India* 

## **1. Introduction**

Advancements made in the field of materials have contributed in a major way in building gas turbine engines with higher power ratings and efficiency levels. Improvements in design of the gas turbine engines over the years have importantly been due to development of materials with enhanced performance levels. Gas turbines have been widely utilized in aircraft engines as well as for land based applications importantly for power generation. Advancements in gas turbine materials have always played a prime role – higher the capability of the materials to withstand elevated temperature service, more the engine efficiency; materials with high elevated temperature strength to weight ratio help in weight reduction. A wide spectrum of high performance materials - special steels, titanium alloys and superalloys - is used for construction of gas turbines. Manufacture of these materials often involves advanced processing techniques. Other material groups like ceramics, composites and inter-metallics have been the focus of intense research and development; aim is to exploit the superior features of these materials for improving the performance of gas turbine engines.

The materials developed at the first instance for gas turbine engine applications had high temperature tensile strength as the prime requirement. This requirement quickly changed as operating temperatures rose. Stress rupture life and then creep properties became important. In the subsequent years of development, low cycle fatigue (LCF) life became another important parameter. Many of the components in the aero engines are subjected to fatigue- and /or creep-loading, and the choice of material is then based on the capability of the material to withstand such loads.

Coating technology has become an integral part of manufacture of gas turbine engine components operating at high temperatures, as this is the only way a combination of high level of mechanical properties and excellent resistance to oxidation / hot corrosion resistance could be achieved.

The review brings out a detailed analysis of the advanced materials and processes that have come to stay in the production of various components in gas turbine engines. While there are thousands of components that go into a gas turbine engine, the emphasis here has been on the main components, which are critical to the performance of the engine. The review also takes stock of the R&D activity currently in progress to develop higher performance materials for gas turbine engine application. On design aspects of gas turbine engines, the reader is referred to the latest edition of the Gas Turbine Engineering Handbook (Boyce, 2006).

Materials for Gas Turbines – An Overview 295

Fig. 1. Relative creep capability of titanium alloys used for compressor parts in the form of a

**3. Compressor blading materials for land based gas turbines – Special steels**  Until recently, all production blades for compressors are made from 12% chromium containing martensitic stainless steel grades 403 or 403 Cb (Schilke, 2004). Corrosion of compressor blades can occur due to moisture containing salts and acids collecting on the blading. To prevent the corrosion, GE has developed patented aluminum slurry coatings for the compressor blades. The coatings are also meant to impart improved erosion resistance to the blades. During the 1980's, GE introduced a new compressor blade material, GTD-450, a precipitation hardened martensitic stainless steel for its advanced and uprated machines (Schilke, 2004). Without sacrificing stress corrosion resistance, GTD-450 offers increased tensile strength, high cycle fatigue strength and corrosion fatigue strength, compared to type 403. GTD-450 also possesses superior resistance to acidic salt environments to type 403, due

to higher concentration of chromium and presence of molybdenum (Schilke, 2004).

Ti64 Ti-6Al-4V 315 Ti811 Ti-8Al-1Mo-1V 400 Alloy 685 Ti-6Al-5Zr-0.5Mo-0.25Si 520 Alloy 829 Ti-5.5Al-3.5Sn-3Zr-1Nb-0.25Mo—0.3Si 550 Alloy 834 Ti-5.8Al-4Sn-3.5Zr-0.7Nb-0.5Mo-0.35Si-0.06C 600 Ti1100 Ti-6Al-2.8Sn-4Zr-0.4Mo-0.4Si 600

Ti6242S Ti-6Al-2Sn-4Zr-2Mo-0.2Si 540

Table 1. Titanium alloys used for compressor parts in aircraft engines – chemical

Nominal chemical composition Maximum service

temperature (oC)

Larson Miller plot (Schematic).

Ti6242 Ti-6Al-2Sn-4Zr-2Mo

composition and maximum service temperature

Grade designation

## **2. Compressor parts for aircraft engines – Titanium alloys**

Titanium, due to its high strength to weight ratio, has been a dominant material in compressor stages in aeroengines. Titanium content has increased from 3 % in 1950s to about 33% today of the aeroengine weight. Unlike predictions made for requirements of ceramic and metal matrix composites for aeroengines, predictions made for titanium alloys have come true or even surpassed. High temperature titanium alloys have found extensive application in aeroengines. Ti-6Al-4V is used for static and rotating components in gas turbine engines. Castings are used to manufacture the more complex static components. Forgings are typically used for the rotating components. For example, the alloy is used for fan disc and low pressure compressor discs and blades for the Pratt and Whitney 4084 engine. The alloy is used in the cooler compressor stages up to a maximum temperature of about 315 oC. Ti-8Al-1Mo-1V is used for fan blades in military engines (Bayer, 1996). The alloys 685 (Ti-6Al-5Zr-0.5Mo-0.25Si) and 829 (Ti-5.5Al-3.5Sn-3Zr-1Nb-0.25Mo-0.3Si) are used in many current European aeroengines such as RB2111, 535E4 in fully beta heat treated condition to maximize creep resistance (Gogia, 2005). Alloy 834 (Ti-5.8Al-4Sn-3.5Zr-0.7Nb-0.5Mo-0.35Si-0.06C), a relatively recent grade, in contrast is used in + condition, with a 5- 15% equiaxed in the microstructure to optimize both creep and fatigue strength (Gogia, 2005). The alloy was aimed at replacing the Alloys 685 and 829 preferred in European jet engines. Alloy 834 is used as a compressor disc material in the last two stages of the medium-pressure compressor, and the first four stages of the high pressure compressor in variants of the Rolls-Royce Trent series commercial jet engine. The Ti-1100 (Ti-6Al-2.8Sn-4Zr-0.4Mo-0.4Si), a competitive alloy to IMI834, is designed to be used in the heat treated condition. The alloy is under evaluation by Allison Gas Turbine Engines for higher thrust versions of their 406/GMA3007/GMA2100 family of engines, primarily for castings (Gogia, 2005). The alloy has a claimed use temperature of 600 oC. IN US, Ti6-2-4-2 (Ti-6Al-2Sn-4Zr-2Mo) is the preferred high temperature alloy for jet engine applications. A variant of this alloy, Ti6-2-4-2S is also commercially available. The 'S' denotes addition of 0.1-0.25 % Si to improve the creep resistance. It is used for rotating components such as blades, discs and rotors at temperatures up to about 540 oC (Bayer, 1996). It is used in high pressure compressors at temperatures too high for Ti-6-4, above about 315 oC, for structural applications.

Today, the maximum temperature limit for near- alloys for elevated temperature applications is about 540 oC. This temperature limitation for titanium alloys mean the hottest parts in the compressor, i.e. the discs and blades of the last compressor stages, have to be manufactured from Ni-based superalloys at nearly twice the weight. Additionally, problems arise associated with the different thermal expansion behavior and the bonding techniques of the two alloy systems. Therefore enormous efforts are underway to develop a compressor made completely of titanium. Titanium alloys are required that can be used at temperatures of 600 oC or higher. This has been the impetus for extensive research and development work in the area of elevated temperature titanium alloys.

Table 1 gives the chemical composition and the maximum service temperature of various grades of titanium alloys mentioned above. Figure 1 shows schematically the relative creep capability of these grades in the form of a Larson Miller plot. The reader is referred to some excellent reviews on use of titanium alloys in gas turbine engines (Bayer, 1998; Gogia, 2005). The technical guide on titanium published by ASM International (Donachie, 2000) also gives much information on titanium as a gas turbine material.

Titanium, due to its high strength to weight ratio, has been a dominant material in compressor stages in aeroengines. Titanium content has increased from 3 % in 1950s to about 33% today of the aeroengine weight. Unlike predictions made for requirements of ceramic and metal matrix composites for aeroengines, predictions made for titanium alloys have come true or even surpassed. High temperature titanium alloys have found extensive application in aeroengines. Ti-6Al-4V is used for static and rotating components in gas turbine engines. Castings are used to manufacture the more complex static components. Forgings are typically used for the rotating components. For example, the alloy is used for fan disc and low pressure compressor discs and blades for the Pratt and Whitney 4084 engine. The alloy is used in the cooler compressor stages up to a maximum temperature of about 315 oC. Ti-8Al-1Mo-1V is used for fan blades in military engines (Bayer, 1996). The alloys 685 (Ti-6Al-5Zr-0.5Mo-0.25Si) and 829 (Ti-5.5Al-3.5Sn-3Zr-1Nb-0.25Mo-0.3Si) are used in many current European aeroengines such as RB2111, 535E4 in fully beta heat treated condition to maximize creep resistance (Gogia, 2005). Alloy 834 (Ti-5.8Al-4Sn-3.5Zr-0.7Nb-0.5Mo-0.35Si-0.06C), a relatively recent grade, in contrast is used in + condition, with a 5- 15% equiaxed in the microstructure to optimize both creep and fatigue strength (Gogia, 2005). The alloy was aimed at replacing the Alloys 685 and 829 preferred in European jet engines. Alloy 834 is used as a compressor disc material in the last two stages of the medium-pressure compressor, and the first four stages of the high pressure compressor in variants of the Rolls-Royce Trent series commercial jet engine. The Ti-1100 (Ti-6Al-2.8Sn-4Zr-0.4Mo-0.4Si), a competitive alloy to IMI834, is designed to be used in the heat treated condition. The alloy is under evaluation by Allison Gas Turbine Engines for higher thrust versions of their 406/GMA3007/GMA2100 family of engines, primarily for castings (Gogia, 2005). The alloy has a claimed use temperature of 600 oC. IN US, Ti6-2-4-2 (Ti-6Al-2Sn-4Zr-2Mo) is the preferred high temperature alloy for jet engine applications. A variant of this alloy, Ti6-2-4-2S is also commercially available. The 'S' denotes addition of 0.1-0.25 % Si to improve the creep resistance. It is used for rotating components such as blades, discs and rotors at temperatures up to about 540 oC (Bayer, 1996). It is used in high pressure compressors at temperatures too high for Ti-6-4, above about 315 oC, for structural

Today, the maximum temperature limit for near- alloys for elevated temperature applications is about 540 oC. This temperature limitation for titanium alloys mean the hottest parts in the compressor, i.e. the discs and blades of the last compressor stages, have to be manufactured from Ni-based superalloys at nearly twice the weight. Additionally, problems arise associated with the different thermal expansion behavior and the bonding techniques of the two alloy systems. Therefore enormous efforts are underway to develop a compressor made completely of titanium. Titanium alloys are required that can be used at temperatures of 600 oC or higher. This has been the impetus for extensive research and development work

Table 1 gives the chemical composition and the maximum service temperature of various grades of titanium alloys mentioned above. Figure 1 shows schematically the relative creep capability of these grades in the form of a Larson Miller plot. The reader is referred to some excellent reviews on use of titanium alloys in gas turbine engines (Bayer, 1998; Gogia, 2005). The technical guide on titanium published by ASM International (Donachie, 2000) also gives

**2. Compressor parts for aircraft engines – Titanium alloys** 

applications.

in the area of elevated temperature titanium alloys.

much information on titanium as a gas turbine material.

Fig. 1. Relative creep capability of titanium alloys used for compressor parts in the form of a Larson Miller plot (Schematic).

## **3. Compressor blading materials for land based gas turbines – Special steels**

Until recently, all production blades for compressors are made from 12% chromium containing martensitic stainless steel grades 403 or 403 Cb (Schilke, 2004). Corrosion of compressor blades can occur due to moisture containing salts and acids collecting on the blading. To prevent the corrosion, GE has developed patented aluminum slurry coatings for the compressor blades. The coatings are also meant to impart improved erosion resistance to the blades. During the 1980's, GE introduced a new compressor blade material, GTD-450, a precipitation hardened martensitic stainless steel for its advanced and uprated machines (Schilke, 2004). Without sacrificing stress corrosion resistance, GTD-450 offers increased tensile strength, high cycle fatigue strength and corrosion fatigue strength, compared to type 403. GTD-450 also possesses superior resistance to acidic salt environments to type 403, due to higher concentration of chromium and presence of molybdenum (Schilke, 2004).


Table 1. Titanium alloys used for compressor parts in aircraft engines – chemical composition and maximum service temperature

Materials for Gas Turbines – An Overview 297

more than 25 years (Schilke, 2004). Both these alloys have been produced through the

Powder Metallurgy (PM) processing is being extensively used in production of superalloy components for gas turbines. PM processing is essentially used for Nickel-based superalloys. It is primarily used for production of high strength alloys used for disc manufacture such as IN100 or Rene95 which are difficult or impractical to forge by conventional methods. LC Astroloy, MERL 76, IN100, Rene95 and Rene88 DT are the PM superalloys where ingot metallurgy route for manufacture of turbine discs was replaced by

Superalloys such as IN-100 or Rene95 difficult or impractical to forge by conventional

Improves homogeneity / minimizes segregation, particularly in complex Ni-base alloy

 Allows closer control of microstructure and better property uniformity within a part than what is possible in cast and ingot metallurgy wrought products. Finer grain size

Consolidated powder products are often super-plastic and amenable to isothermal

It is a near net shape process; hence significantly less raw material input required and

Several engines manufactured by General Electric and Pratt and Whitney are using

718 Ni19Cr18.5Fe3Mo0.9Ti0.5Al5.1Cb 0.03C Nickel-iron-base superalloy;

Iron-base superalloy; ingot

ingot metallurgy route

Nickel-base superalloy; powder metallurgy route

Nickel-base superalloy; powder metallurgy route

Nickel-base superalloy; powder metallurgy route

Nickel-base superalloy; powder metallurgy route

Nickel-base superalloy; powder metallurgy route

Nickel-base superalloy; ingot metallurgy / powder

Low C, low B variant of

metallurgy route

Udimet 720.

metallurgy route

also reduced machining cost, than in case of conventional ingot metallurgy.

Alloy development flexibility due to elimination of macro-segregation.

conventional ingot metallurgy route.

The advantages of PM processing are listed in the following:

methods. P/M processing provides a solution

forging, reducing force requirements for forging.

Table 4 gives the details of disc superalloys for aircraft engines.

IN 100 60Ni10Cr15Co3Mo4.7Ti5.5Al0.15C 0.015B

Rene 95 61Ni14Cr8Co3.5Mo3.5W3.5Nb2.5Ti3.5Al 0.16C0.01B0.05Zr

Grade Chemical composition Remarks

4.3Ti5.1Al0.02C0.03B0.35Hf0.06Zr

57Ni16Cr15Co3Mo1.25W5Ti2.5Al0.025C0.

superalloy discs manufactured through PM route.

A286 Fe15Cr25Ni1.2Mo2Ti0.3Al0.25V 0.08C 0.006B

0.06Zr1.0V

LC Astroloy 56.5Ni15Cr 15Co5.25Mo3.5Ti4.4Al 0.06C0.03B0.06Zr

Rene88 DT 56.4Ni16cr13Co4Mo4W0.7Nb3.7Ti

0.033B0.03Zr

018B0.03Zr

Table 4. Disc superalloys for aircraft engines

2.1Al0.03C0.015B0.03Zr

Udimet 720 55Ni18Cr14.8Co3Mo1.25W5Ti2.5Al0.035C

MERL-76 54.4Ni12.4Cr18.6co3.3Mo1.4Nb

the PM route.

systems

Udimet 720LI

can be realized.

Grade Chemical composition Remarks AISI 403 Fe12Cr0.11C Martensitic stainless steel AISI 403+Nb Fe12Cr0.2Cb0.15C Martensitic stainless steel with Nb addition GTD-450 Fe15.5Cr6.3Ni0.8Mo0.03C Precipitation hardening stainless steel

Table 2 gives the chemical composition of the different steel grades used for compressor blading.

Table 2. Compressor blade materials for land based gas turbines

## **4. Combustion hardware for aircraft and industrial gas turbines (IGTs)**

Driven by the increased firing temperatures of the gas turbines and the need for improved emission control, significant development efforts have been made to advance the combustion hardware, by way of adopting sophisticated materials and processes.

The primary basis for the material changes that have been made is improvement of high temperature creep rupture strength without sacrificing the oxidation / corrosion resistance. Traditionally combustor components have been fabricated out of sheet nickel-base superalloys. Hastelloy X, a material with higher creep strength was used from 1960s to 1980s. Nimonic 263 was subsequently introduced and has still higher creep strength (Schilke, 2004). As firing temperatures further increased in the newer gas turbine models, HA-188, a cobalt base superalloy has been recently adopted for some combustion system components for improved creep rupture strength (Schilke, 2004). Coutsouradis et al. reviewed the applications of cobalt-base superalloys for combustor and other components in gas turbines (Coutsouradis et al., 1987). Nickel base superalloys 617 and 230 find wide application for combustor components (Wright & Gibbons, 2007). Table 3 gives the chemical composition of combustor materials.


Table 3. Combustor materials

In addition to designing with improved materials, combustion liners and transition pieces of advanced and uprated machines involving higher firing temperatures are given a thermal barrier coating (TBC). The coating serves to provide an insulating layer and reduces the underlying base metal temperature. Section 9 deals with the subject of TBC in detail.

## **5. Turbine disk applications**

## **5.1 Aircraft engines – Superalloys**

A286, an austenitic iron-base alloy has been used for years in aircraft engine applications (Schilke, 2004). Superalloy 718 has been used for manufacture of discs in aircraft engines for

Table 2 gives the chemical composition of the different steel grades used for compressor

AISI 403+Nb Fe12Cr0.2Cb0.15C Martensitic stainless steel with Nb addition GTD-450 Fe15.5Cr6.3Ni0.8Mo0.03C Precipitation hardening stainless steel

Driven by the increased firing temperatures of the gas turbines and the need for improved emission control, significant development efforts have been made to advance the

The primary basis for the material changes that have been made is improvement of high temperature creep rupture strength without sacrificing the oxidation / corrosion resistance. Traditionally combustor components have been fabricated out of sheet nickel-base superalloys. Hastelloy X, a material with higher creep strength was used from 1960s to 1980s. Nimonic 263 was subsequently introduced and has still higher creep strength (Schilke, 2004). As firing temperatures further increased in the newer gas turbine models, HA-188, a cobalt base superalloy has been recently adopted for some combustion system components for improved creep rupture strength (Schilke, 2004). Coutsouradis et al. reviewed the applications of cobalt-base superalloys for combustor and other components in gas turbines (Coutsouradis et al., 1987). Nickel base superalloys 617 and 230 find wide application for combustor components (Wright & Gibbons, 2007). Table 3 gives the chemical composition of combustor materials.

**4. Combustion hardware for aircraft and industrial gas turbines (IGTs)** 

combustion hardware, by way of adopting sophisticated materials and processes.

Grade Chemical composition Remarks

Table 3. Combustor materials

**5. Turbine disk applications 5.1 Aircraft engines – Superalloys** 

Hastelloy X Ni22Cr1.5Co1.9Fe0.7W9Mo0.07C0.005B Nickel-base superalloy Nimonic 263 Ni20Cr20Co0.4Fe6Mo2.1Ti0.4Al0.06C Nickel-base superalloy HA188 Co22Cr22Ni1.5Fe14W0.05C0.01B Cobalt-base superalloy 617 54Ni22Cr12.5Co8.5Mo1.2Al Nickel-base superalloy 230 55Ni22Cr5Co3Fe14W2Mo0.35Al0.10C0.015B Nickel-base superalloy;

In addition to designing with improved materials, combustion liners and transition pieces of advanced and uprated machines involving higher firing temperatures are given a thermal barrier coating (TBC). The coating serves to provide an insulating layer and reduces the

A286, an austenitic iron-base alloy has been used for years in aircraft engine applications (Schilke, 2004). Superalloy 718 has been used for manufacture of discs in aircraft engines for

underlying base metal temperature. Section 9 deals with the subject of TBC in detail.

values for Co, Fe and B are upper limits.

Grade Chemical composition Remarks

Table 2. Compressor blade materials for land based gas turbines

AISI 403 Fe12Cr0.11C Martensitic stainless steel

blading.

more than 25 years (Schilke, 2004). Both these alloys have been produced through the conventional ingot metallurgy route.

Powder Metallurgy (PM) processing is being extensively used in production of superalloy components for gas turbines. PM processing is essentially used for Nickel-based superalloys. It is primarily used for production of high strength alloys used for disc manufacture such as IN100 or Rene95 which are difficult or impractical to forge by conventional methods. LC Astroloy, MERL 76, IN100, Rene95 and Rene88 DT are the PM superalloys where ingot metallurgy route for manufacture of turbine discs was replaced by the PM route.

The advantages of PM processing are listed in the following:


Several engines manufactured by General Electric and Pratt and Whitney are using superalloy discs manufactured through PM route.


Table 4 gives the details of disc superalloys for aircraft engines.

Table 4. Disc superalloys for aircraft engines

Materials for Gas Turbines – An Overview 299

to reduce stringers and clusters of carbides, borides or carbonitrides have led to the development of the Alloy 720LI. Both these alloys have been of considerable interest to land based gas turbines. They have also been incorporated in some aircraft gas turbines (Furrer & Fecht, 1999). Table 5 gives details of special steels / superalloys used for production of discs

The reader is referred to an overview by Furrer and Fecht on nickel-based superalloys for turbine discs for land based power generation and aircraft propulsion (Furrer & Fecht,

Recognition of the material creep strength as an important consideration for the gas turbine engines, understanding generated between age hardening, creep and volume fraction and the steadily increasing operating-temperature requirements for the aircraft engines resulted in development of wrought alloys with increasing levels of aluminum plus titanium. Component forgeability problems led to this direction of development not going beyond a certain extent. The composition of the wrought alloys became restricted by the hot workability requirements. This situation led to the development of cast nickel-base alloys. Casting compositions can be tailored for good high temperature strength as there was no forgeability requirement. Further the cast components are intrinsically stronger than forgings at high temperatures, due to the coarse grain size of castings. Das recently

Buckets (rotating airfoils) must withstand severe combination of temperature, stress and environment. The stage 1 bucket is particularly loaded, and is generally the limiting component of the gas turbine. Function of the nozzles (stationary airfoils) is to direct the hot gases towards the buckets. Therefore they must be able to withstand high temperatures. However they are subjected to lower mechanical stresses than the buckets. An important design requirement for the nozzle materials is that they should possess excellent high

Cast alloy IN-713 was among the early grades established as the materials for the airfoils in the most demanding gas turbine application. Efforts to increase the volume fraction to realize higher creep strength led to the availability of alloys like IN 100 and Rene 100 for airfoils in gas turbine engines. Increased amount of refractory solid solution strengtheners such as W and Mo were added to some of the grades developed later and this led to the availability of grades like MAR-M200, MAR-M246, IN 792 and M22. Addition of 2 wt% Hf improved ductility and a new series of alloys became available with Hf addition such as

General Electric pursued own alloy development with Rene 41, Rene 77, Rene 80 and Rene 80+Hf having relatively high chromium content for improved corrosion resistance at the cost of some high temperature strength. Other similar alloys with high chromium content

Table 6 gives details of superalloy compositions of airfoils produced by conventional

for land-based gas turbines.

**6. Turbine blades and vanes – Cast superalloys** 

temperature oxidation and corrosion resistance.

MAR-M200+Hf, MAR-M246+Hf, Rene 125+Hf.

are IN738C, IN738LC, Udimet 700, Udimet 710.

equiaxed investment casting process.

**6.1.1 Aircraft engines** 

**6.1 Conventional equiaxed investment casting process** 

reviewed the advances made in nickel-based cast superalloys (Das, 2010).

1999).

## **5.2 IGTs – Steels and superalloys**

Turbine discs of most GE single shaft heavy duty gas turbines are made of 1%Cr-1.25%Mo-0.25%V steel in hardened and tempered condition (Schilke, 2004). 12%Cr steels such as M152 have higher rupture strength than Cr-Mo-V steel, in addition to outstanding fracture toughness and capacity to attain uniform and high mechanical properties in large sections. Use of A286 for IGTs started in 1965, when the technological advancements made it possible to produce large ingots of this material with required quality (Schilke, 2004).

With the advent of advanced of gas turbine engines with much higher firing temperatures and compressor ratios, it became necessary to utilize a nickel-base superalloy, alloy 706 for the rotors. The use of this material provides the necessary temperature capability required to also meet the firing temperature requirements in the future. This alloy is similar to the Alloy 718, an alloy that has been used for rotors in aircraft engines for more than 25 years. Alloy 706 contains lower concentrations of alloying elements known for their tendency to segregate. Consequently it is less segregation-prone than Alloy 718 and could be produced in large diameters unlike Alloy 718. Accordingly large sized rotors of Alloy 706 could be produced to serve large IGTs for land-based power generation (Schilke, 2004). Alloy 718, the most frequently used superalloy for aircraft gas turbines, because of its segregation tendency, could be produced, until the turn of the century, to a maximum ingot size of 500mm. Developments made with reference to remelting techniques, together with very close control on chemical composition have enabled production of ingots of Alloy 718 as large as 750 mm in diameter. This has resulted in the ability to process Alloy 718 to the large disk sizes needed in modern IGTs.

The importance of Alloy 718 and Alloy 706 can be seen from the fact that several international conferences have been devoted to developments related to these alloys (Loria, 1989, 1991, 1994, 1997, 2001, 2005; Caron et al., 2008).


Table 5. Disc materials for IGTs

Udimet 720 also evolved as an advanced wrought alloy for land based gas turbines. Reductions in Cr content to prevent sigma phase formation and in carbon and boron levels

Turbine discs of most GE single shaft heavy duty gas turbines are made of 1%Cr-1.25%Mo-0.25%V steel in hardened and tempered condition (Schilke, 2004). 12%Cr steels such as M152 have higher rupture strength than Cr-Mo-V steel, in addition to outstanding fracture toughness and capacity to attain uniform and high mechanical properties in large sections. Use of A286 for IGTs started in 1965, when the technological advancements made it possible

With the advent of advanced of gas turbine engines with much higher firing temperatures and compressor ratios, it became necessary to utilize a nickel-base superalloy, alloy 706 for the rotors. The use of this material provides the necessary temperature capability required to also meet the firing temperature requirements in the future. This alloy is similar to the Alloy 718, an alloy that has been used for rotors in aircraft engines for more than 25 years. Alloy 706 contains lower concentrations of alloying elements known for their tendency to segregate. Consequently it is less segregation-prone than Alloy 718 and could be produced in large diameters unlike Alloy 718. Accordingly large sized rotors of Alloy 706 could be produced to serve large IGTs for land-based power generation (Schilke, 2004). Alloy 718, the most frequently used superalloy for aircraft gas turbines, because of its segregation tendency, could be produced, until the turn of the century, to a maximum ingot size of 500mm. Developments made with reference to remelting techniques, together with very close control on chemical composition have enabled production of ingots of Alloy 718 as large as 750 mm in diameter. This has resulted in the ability to process Alloy 718 to the large

The importance of Alloy 718 and Alloy 706 can be seen from the fact that several international conferences have been devoted to developments related to these alloys (Loria,

Fe1Cr0.5Ni1.25Mo0.25V0.30C Medium carbon

low alloy steel

Iron-base superalloy

superalloy

superalloy

superalloy

superalloy

Grade Chemical composition Remarks

M152 Fe12Cr2.5Ni1.7Mo0.3V0.12C 12% Cr steel

706 Ni16Cr37Fe1.8Ti2.9Cb0.03C Nickel-iron-base

718 Ni19Cr18.5Fe3Mo0.9Ti0.5Al5.1Cb 0.03C Nickel-iron-base

55Ni18Cr14.8Co3Mo1.25W5Ti2.5Al0.035C0.033B0.03Zr Nickel-base

57Ni16Cr15Co3Mo1.25W5Ti2.5Al0.025C0.018B0.03Zr Nickel-base

Udimet 720 also evolved as an advanced wrought alloy for land based gas turbines. Reductions in Cr content to prevent sigma phase formation and in carbon and boron levels

to produce large ingots of this material with required quality (Schilke, 2004).

**5.2 IGTs – Steels and superalloys** 

disk sizes needed in modern IGTs.

CrMoV steel

Udimet 720

Udimet 720LI

Table 5. Disc materials for IGTs

1989, 1991, 1994, 1997, 2001, 2005; Caron et al., 2008).

A286 Fe15Cr25Ni1.2Mo2Ti0.3Al0.25V 0.08C 0.006B

to reduce stringers and clusters of carbides, borides or carbonitrides have led to the development of the Alloy 720LI. Both these alloys have been of considerable interest to land based gas turbines. They have also been incorporated in some aircraft gas turbines (Furrer & Fecht, 1999). Table 5 gives details of special steels / superalloys used for production of discs for land-based gas turbines.

The reader is referred to an overview by Furrer and Fecht on nickel-based superalloys for turbine discs for land based power generation and aircraft propulsion (Furrer & Fecht, 1999).

## **6. Turbine blades and vanes – Cast superalloys**

Recognition of the material creep strength as an important consideration for the gas turbine engines, understanding generated between age hardening, creep and volume fraction and the steadily increasing operating-temperature requirements for the aircraft engines resulted in development of wrought alloys with increasing levels of aluminum plus titanium. Component forgeability problems led to this direction of development not going beyond a certain extent. The composition of the wrought alloys became restricted by the hot workability requirements. This situation led to the development of cast nickel-base alloys. Casting compositions can be tailored for good high temperature strength as there was no forgeability requirement. Further the cast components are intrinsically stronger than forgings at high temperatures, due to the coarse grain size of castings. Das recently reviewed the advances made in nickel-based cast superalloys (Das, 2010).

Buckets (rotating airfoils) must withstand severe combination of temperature, stress and environment. The stage 1 bucket is particularly loaded, and is generally the limiting component of the gas turbine. Function of the nozzles (stationary airfoils) is to direct the hot gases towards the buckets. Therefore they must be able to withstand high temperatures. However they are subjected to lower mechanical stresses than the buckets. An important design requirement for the nozzle materials is that they should possess excellent high temperature oxidation and corrosion resistance.

#### **6.1 Conventional equiaxed investment casting process 6.1.1 Aircraft engines**

Cast alloy IN-713 was among the early grades established as the materials for the airfoils in the most demanding gas turbine application. Efforts to increase the volume fraction to realize higher creep strength led to the availability of alloys like IN 100 and Rene 100 for airfoils in gas turbine engines. Increased amount of refractory solid solution strengtheners such as W and Mo were added to some of the grades developed later and this led to the availability of grades like MAR-M200, MAR-M246, IN 792 and M22. Addition of 2 wt% Hf improved ductility and a new series of alloys became available with Hf addition such as MAR-M200+Hf, MAR-M246+Hf, Rene 125+Hf.

General Electric pursued own alloy development with Rene 41, Rene 77, Rene 80 and Rene 80+Hf having relatively high chromium content for improved corrosion resistance at the cost of some high temperature strength. Other similar alloys with high chromium content are IN738C, IN738LC, Udimet 700, Udimet 710.

Table 6 gives details of superalloy compositions of airfoils produced by conventional equiaxed investment casting process.

Materials for Gas Turbines – An Overview 301

GE engines use FSX 414, a GE-patented cobalt base alloy for all stage 1 nozzles and some later stage nozzles. Cobalt base alloys possess superior strength at very high temperatures compared to nickel base superalloys – hence the choice of cobalt base superalloy. It has a two-three fold oxidation resistance compared to X40 and X45, also cobalt based superalloys used for nozzle applications. Use of FSX 414 over C40/C45 hence enables increased firing

Later stage nozzles must also possess adequate creep strength and GE developed a nickel base superalloy GTD222 for some stage 2 and stage 3 applications. The alloy has significantly higher creep strength compared to FSX414. N155, an iron-based superalloy, has good weldability and is used for later stage nozzles of some GE engines (Schilke, 2004).

X40 Co-25Cr10Ni8W1Fe0.5C0.01B Cobalt-base superalloy X45 Co-25Cr10Ni8W1Fe0.25C0.01B Cobalt-base superalloy FSX414 Co-28Cr10Ni7W1Fe0.25C0.01B Cobalt-base superalloy N155 Fe-21Cr20Ni20Co2.5W3Mo0.20C Iron-base superalloy GTD-222 Ni-22.5Cr19Co2.0W2.3Mo1.2Ti0.8Al0.10V 0.008C1.0B Nickel-base superalloy

The major failure mechanism for gas turbine airfoils involved nucleation and growth of cavities along transverse grain boundaries. Elimination of transverse grain boundaries through directional solidification of turbine blades and vanes made an important step in temperature capability of these castings. Use of DS superalloys could improve the turbine blade metal temperature capability by about 14 oC relative to the conventionally cast

Chemical composition Remarks

generation

generation

generation

generation

59.5Ni9Cr10Co12.5W2Hf1.8Nb2Ti5Al0.05Zr0.15C0.015B First

CM247LC 61.7Ni8.1Cr9.2Co0.5Mo9.5W3.2Ta1.4Hf0.7Ti5.6Al0.01Zr0.07C0.015B First

PWA1422 59.2Ni9Cr10Co12W1.5Hf1Nb2Ti5Al0.1Zr0.14C0.015B First

DMD4 66.8Ni2.4Cr4Co5.5W6.5Re8Ta1.2Hf0.3Nb5.2Al0.07C0.01B Third

Table 9. DS nickel-base superalloys for blading applications in aircraft engines

**6.1.2.2 Nozzle materials for land based gas turbines**

temperatures for a given oxidation life (Schilke, 2004).

Table 8. Nozzle materials for IGTs

**6.2.1 Aircraft engines** 

superalloys.

Grade designation

DS MAR M-200+Hf

**6.2 Directionally solidified (DS) castings** 

Table 8 gives the details of materials used for nozzles in IGTs.

Grade Chemical composition Remarks


Table 6. Conventionally cast nickel-base superalloys for gas turbine blading applications in aircraft gas turbines

## **6.1.2 Land-based gas turbine engines**

#### **6.1.2.1 Bucket materials for land based gas turbines**

Many of the GE engines used U-500 for stage 1 buckets in mid1960's. It is being used for later stages of buckets in selected gas turbine models (Schilke, 2004). IN738 has been used as stage 1 bucket material on several GE engines during 1971-1984. In recent years it has been also used as stage 2 bucket material in some GE engines (Schilke, 2004). The alloy has an outstanding combination of elevated temperature strength and hot corrosion resistance and this makes it attractive for heavy duty gas turbine applications. Developments in processing technology have enabled production of the alloy in large ingot sizes. The alloy is used throughout the heavy duty gas turbine industry. Subsequently GE has developed the alloy GTD-111, with higher strength levels than 738, but maintaining its hot corrosion resistance. GTD-111 has replaced IN738 as bucket material in different GE engine models (Schilke, 2004).

Table7 gives details of conventionally cast superalloys for blading applications in IGTs.


Table 7. Conventionally cast nickel-base superalloys for blading applications in IGTs

IN 713 74.2Ni12.5Cr4.2Mo2Nb0.8Ti6.1Al0.1Zr0.12C0.01B IN 100 60.5Ni10Cr15Co3Mo4.7Ti5.5Al0.06Zr0.18C0.014B Rene 100 62.6Ni9.5Cr15Co3Mo4.2Ti5.5Al0.06Zr0.15C0.015B MAR-M200 59.5Ni9Cr10Co12.5W1.8Nb2Ti5Al0.05Zr0.15C0.015B

MAR-M200+Hf Ni8Cr9Co12W2Hf1Nb1.9Ti5.0Al0.03Zr0.13C0.015B

Rene 41 56Ni19Cr10.5Co9.5Mo3.2Ti1.7Al0.01Zr0.08C0.005B Rene 77 53.5Ni15Cr18.5Co5.2Mo3.5Ti4.25Al0.08C0.015B Rene 80 60.3Ni14Cr9.5Co4Mo4W5Ti3al0.03Zr0.17C0.015B

M 22 71.3Ni5.7Cr2Mo11W3Ta6.3Al0.6Zr0.13C

MAR-M246 59.8Ni9Cr10Co2.5Mo10W1.5Ta1.5Ti5.5Al0.05Zr0.14C0.015B IN 792 60.8Ni12.7Cr9Co2Mo3.9W3.9Ta4.2Ti3.2Al0.1Zr0.21C0.02B

MAR-M246+Hf Ni9Cr10Co2.5Mo10W1.5Hf1.5Ta1.5Ti5.5Al0.05Zr0.15C0.015B

Rene 80+Hf 59.8Ni14Cr9.5Co4Mo4W0.8Hf4.7Ti3Al0.01Zr0.15C0.015B

Udimet 700 59Ni14.3Cr14.5Co4.3Mo3.5Ti4.3Al0.02Zr0.08C0.015B Udimet 710 54.8Ni18Cr15Co3Mo1.5W2.5Ti5Al0.08Zr0.13C TMD-103 59.8Ni3Cr12Co2Mo6W5Re6Ta0.1Hf6Al

replaced IN738 as bucket material in different GE engine models (Schilke, 2004).

IN738 Ni16Cr8.3Co0.2Fe2.6W1.75Mo3.4Ti3.4Al0.9Cb0.10C0.001B1.75Ta

Table 7. Conventionally cast nickel-base superalloys for blading applications in IGTs

Table7 gives details of conventionally cast superalloys for blading applications in IGTs.

IN 738 61.5Ni16Cr8.5Co1.75Mo2.6W1.75Ta0.9Nb3.4Ti3.4Al0.04Zr0.11C0.01B

Table 6. Conventionally cast nickel-base superalloys for gas turbine blading applications in

Many of the GE engines used U-500 for stage 1 buckets in mid1960's. It is being used for later stages of buckets in selected gas turbine models (Schilke, 2004). IN738 has been used as stage 1 bucket material on several GE engines during 1971-1984. In recent years it has been also used as stage 2 bucket material in some GE engines (Schilke, 2004). The alloy has an outstanding combination of elevated temperature strength and hot corrosion resistance and this makes it attractive for heavy duty gas turbine applications. Developments in processing technology have enabled production of the alloy in large ingot sizes. The alloy is used throughout the heavy duty gas turbine industry. Subsequently GE has developed the alloy GTD-111, with higher strength levels than 738, but maintaining its hot corrosion resistance. GTD-111 has

Grade designation Chemical composition

aircraft gas turbines

**6.1.2 Land-based gas turbine engines** 

Grade Chemical composition

Udimet 500 Ni18.5Cr18.5Co4Mo3Ti3Al0.07C0.006B Rene 77 Ni15Cr17Co5.3Mo3.35Ti4.25Al0.07C0.02B

GTD 111 Ni14Cr9.5Co3.8W1.5Mo4.9Ti3.0Al0.10C0.01B2.8Ta

**6.1.2.1 Bucket materials for land based gas turbines** 

## **6.1.2.2 Nozzle materials for land based gas turbines**

GE engines use FSX 414, a GE-patented cobalt base alloy for all stage 1 nozzles and some later stage nozzles. Cobalt base alloys possess superior strength at very high temperatures compared to nickel base superalloys – hence the choice of cobalt base superalloy. It has a two-three fold oxidation resistance compared to X40 and X45, also cobalt based superalloys used for nozzle applications. Use of FSX 414 over C40/C45 hence enables increased firing temperatures for a given oxidation life (Schilke, 2004).

Later stage nozzles must also possess adequate creep strength and GE developed a nickel base superalloy GTD222 for some stage 2 and stage 3 applications. The alloy has significantly higher creep strength compared to FSX414. N155, an iron-based superalloy, has good weldability and is used for later stage nozzles of some GE engines (Schilke, 2004). Table 8 gives the details of materials used for nozzles in IGTs.


Table 8. Nozzle materials for IGTs

## **6.2 Directionally solidified (DS) castings**

## **6.2.1 Aircraft engines**

The major failure mechanism for gas turbine airfoils involved nucleation and growth of cavities along transverse grain boundaries. Elimination of transverse grain boundaries through directional solidification of turbine blades and vanes made an important step in temperature capability of these castings. Use of DS superalloys could improve the turbine blade metal temperature capability by about 14 oC relative to the conventionally cast superalloys.


Table 9. DS nickel-base superalloys for blading applications in aircraft engines

Materials for Gas Turbines – An Overview 303

of SC superalloy grades PWA 1484, CMSX4, Rene N5, TUT92. These grades gave about 30

Development of SC superalloys continued with the target of achieving another 30 oC improvement in metal temperature capability while maintaining the environmental resistance and freedom from appearance of deleterious phases in the microstructure. This led to emergence of grades CMSX10, Rene-6, TMS75, TMS80, MC-NG developed by Onera in France, DMS4 developed by DMRL, India, TMS-196, developed by NIMS, Japan. Detailed studies / evaluation have been carried out on these grades and they are potential candidate

Figure 3 schematically shows the improvement in stress rupture strength of superalloys, by moving over from DS (CM247) to first generation SC (CMSX2) to second generation SC (CMSX4) to third to fifth generation (CMSX10 and TMS 196) SCs, in the form of a Larson

Table 11 gives details of the superalloys used for blading applications in aircraft engines.

Table 12 gives details of SC superalloys used for rotating blade application in IGTs.

Fig. 2. Relative creep deformation of equiaxed, DS and SC superalloy castings (schematic)

Development of SC castings has also benefited to improve the efficiency of combined cycle power plants by way of increasing the engine firing temperatures. GE has been applying the SC bucket technology for last several years. SC alloys such as CMSX11B, AF56, PWA1483 containing about 12%Cr for long term environmental resistance together with additions of C, B, Hf to enhance alloy tolerance to low angle boundaries have been developed as airfoil materials. SC alloys such as CMSX 11C and SC 16 have been developed with Cr >12% to increase resistance to hot corrosion and oxidation. Long term phase stability was an important consideration in design of these alloys. Gibbons reviewed the improvements that are taking place with reference to alloys and coatings for integrated gasification combined

oC metal temperature improvement over the early SC superalloys.

alloys for future gas turbine engines with enhanced performance.

Miller plot.

**6.3.2 Land-based gas turbines** 

cycle systems (IGCC) (Gibbons, 2009)

By early 1980s, DS superalloys became available and were operating in gas turbines. DS MAR-M-200+Hf became available. Another DS grade CM247LC is the outcome of extensive efforts to optimize the chemical composition to improve carbide microstructure, grain boundary cracking resistance, to minimize the formation of deleterious secondary phases and to avoid HfO2 inclusion problem. Pratt and Whitney developed an equivalent DS grade PWA 1422.

Table 9 gives details of DS superalloy compositions for aircraft engines.

## **6.2.2 Land-based gas turbine engines**

GE has been using the DS version of DTD-111 for stage 1 buckets of different engines. It is same as DTD-111 equiaxed, except tighter control on alloy chemistry. DS version of DTD-111 is stated to possess improved creep life, improved fatigue life and higher impact strength, compared to equiaxed version (Schilke, 2004). Use of DS superalloys could improve the turbine blade metal temperature capability by about 14 oC relative to the conventionally cast superalloys. TMD-103 belongs to the recent advances in DS alloy castings for IGT airfoil castings. It has very attractive long term creep rupture strength and hot corrosion resistance. The alloy could be directionally solidified in the form of large hollow blades for 2000KW IGT. Alloy chemistry of IGT buckets/vanes differs greatly from that of aeroengine blade/vane alloys, both on account of different operating scenarios and DS processing difficulties due to the large size of IGT components. Table 10 gives details of DS superalloy compositions for airfoils in IGTs.

## **6.3 Single crystals**

In single crystal (SC) castings all grain boundaries are eliminated from the microstructure and an SC with a controlled orientation is produced in an airfoil shape. SCs required no grain boundary strengtheners such as C, B, Zr and Hf. Elimination of these elements while designing the SC compositions helped in raising the melting temperature and correspondingly the high temperature strength. Figure 2 schematically shows the improvement in creep strength of a cast superalloy by switching over from equiaxed polycrystalline investment casting to DS casting to SC casting.


Table 10. DS nickel-base superalloys for application as rotating blades in IGTs

#### **6.3.1 Aircraft engines**

The early SC superalloys included RR2000, RR2060 of Rolls Royce, PWA1480 of Pratt and Whitney, CMSX2 and CMSX3 of Cannon Muskegon and ReneN4 of GE. These SC alloys provided about 20 oC metal temperature advantage over the existing DS alloys.

Attempts to further improve the metal temperature capability of SC superalloys by way of increasing the refractory alloying elements, prominently Rhenium, led to the development

By early 1980s, DS superalloys became available and were operating in gas turbines. DS MAR-M-200+Hf became available. Another DS grade CM247LC is the outcome of extensive efforts to optimize the chemical composition to improve carbide microstructure, grain boundary cracking resistance, to minimize the formation of deleterious secondary phases and to avoid HfO2 inclusion problem. Pratt and Whitney developed an equivalent DS grade

GE has been using the DS version of DTD-111 for stage 1 buckets of different engines. It is same as DTD-111 equiaxed, except tighter control on alloy chemistry. DS version of DTD-111 is stated to possess improved creep life, improved fatigue life and higher impact strength, compared to equiaxed version (Schilke, 2004). Use of DS superalloys could improve the turbine blade metal temperature capability by about 14 oC relative to the conventionally cast superalloys. TMD-103 belongs to the recent advances in DS alloy castings for IGT airfoil castings. It has very attractive long term creep rupture strength and hot corrosion resistance. The alloy could be directionally solidified in the form of large hollow blades for 2000KW IGT. Alloy chemistry of IGT buckets/vanes differs greatly from that of aeroengine blade/vane alloys, both on account of different operating scenarios and DS processing difficulties due to the large size of IGT components. Table 10 gives details of

In single crystal (SC) castings all grain boundaries are eliminated from the microstructure and an SC with a controlled orientation is produced in an airfoil shape. SCs required no grain boundary strengtheners such as C, B, Zr and Hf. Elimination of these elements while designing the SC compositions helped in raising the melting temperature and correspondingly the high temperature strength. Figure 2 schematically shows the improvement in creep strength of a cast superalloy by switching over from equiaxed polycrystalline investment

DTD 111 Same as DTD 111 for equiaxed version, but with tighter

Table 10. DS nickel-base superalloys for application as rotating blades in IGTs

provided about 20 oC metal temperature advantage over the existing DS alloys.

TMD-103 59.8Ni3Cr12Co2Mo6W5Re6Ta0.1Hf6Al0.07C0.015B Third

The early SC superalloys included RR2000, RR2060 of Rolls Royce, PWA1480 of Pratt and Whitney, CMSX2 and CMSX3 of Cannon Muskegon and ReneN4 of GE. These SC alloys

Attempts to further improve the metal temperature capability of SC superalloys by way of increasing the refractory alloying elements, prominently Rhenium, led to the development

control on alloy chemistry

Chemical composition Remarks

generation superalloy

Table 9 gives details of DS superalloy compositions for aircraft engines.

PWA 1422.

**6.3 Single crystals** 

Grade designation

casting to DS casting to SC casting.

**6.3.1 Aircraft engines** 

**6.2.2 Land-based gas turbine engines** 

DS superalloy compositions for airfoils in IGTs.

of SC superalloy grades PWA 1484, CMSX4, Rene N5, TUT92. These grades gave about 30 oC metal temperature improvement over the early SC superalloys.

Development of SC superalloys continued with the target of achieving another 30 oC improvement in metal temperature capability while maintaining the environmental resistance and freedom from appearance of deleterious phases in the microstructure. This led to emergence of grades CMSX10, Rene-6, TMS75, TMS80, MC-NG developed by Onera in France, DMS4 developed by DMRL, India, TMS-196, developed by NIMS, Japan. Detailed studies / evaluation have been carried out on these grades and they are potential candidate alloys for future gas turbine engines with enhanced performance.

Figure 3 schematically shows the improvement in stress rupture strength of superalloys, by moving over from DS (CM247) to first generation SC (CMSX2) to second generation SC (CMSX4) to third to fifth generation (CMSX10 and TMS 196) SCs, in the form of a Larson Miller plot.

Table 11 gives details of the superalloys used for blading applications in aircraft engines.

## **6.3.2 Land-based gas turbines**

Development of SC castings has also benefited to improve the efficiency of combined cycle power plants by way of increasing the engine firing temperatures. GE has been applying the SC bucket technology for last several years. SC alloys such as CMSX11B, AF56, PWA1483 containing about 12%Cr for long term environmental resistance together with additions of C, B, Hf to enhance alloy tolerance to low angle boundaries have been developed as airfoil materials. SC alloys such as CMSX 11C and SC 16 have been developed with Cr >12% to increase resistance to hot corrosion and oxidation. Long term phase stability was an important consideration in design of these alloys. Gibbons reviewed the improvements that are taking place with reference to alloys and coatings for integrated gasification combined cycle systems (IGCC) (Gibbons, 2009)

Table 12 gives details of SC superalloys used for rotating blade application in IGTs.

Fig. 2. Relative creep deformation of equiaxed, DS and SC superalloy castings (schematic)

Materials for Gas Turbines – An Overview 305

PW1480 62.5Ni10Cr5Co4W12Ta1.5Ti5Al 1060

Second PWA1484 59.4Ni5Cr10Co2Mo6W3Re9Ta5.6Al 1120

fifth CMSX10 69.6Ni2Cr3Co0.4Mo5W6Re8Ta0.03Hf0.1Nb0.2Ti5.7Al 1135

 ReneN6 57.3Ni4.2Cr12.5Co1.4Mo6W5.4R e7.2Ta0.15Hf5.8Al0.05C0.004B 1110 TMS 75 59.9Ni3Cr12Co2Mo6W5Re6Ta0.1Hf6Al 1115

 DMS4 67Ni2.4Cr4Co5.5W6.5Re9Ta0.1Hf0.3Nb5.2Al 1140 TMS 196 59.7Ni4.6Cr5.6Co2.4Mo5.0W6.4Re5.6Ta0.1Hf5.6Al5.0Ru 1150

Table 11. SC nickel-base superalloys for blading applications in aircraft engines

CMSX11B 62.1Ni12.5Cr7Co0.5Mo5W5Ta0.04Hf0.1Nb4.2Ti3.6Al

PWA1483 60.3Ni12.2Cr9Co1.9Mo3.8W5Ta0.5Hf4.1Ti3.6Al0.07C0.008B CMSX11C 64.5Ni14.9Cr3Co0.4Mo4.5W5Ta0.04Hf0.1Nb4.2Ti3.4Al

al., 2002). When there is no presence of alkali metal contaminants and sulfur in the environment, high temperature oxidation dominates. Higher the temperature, more rapid is the oxidation. In this context, it is important to note the difference in the operating conditions of aircraft engines and heavy-duty gas turbines. Metal temperatures in the former are higher. The latter are subjected to excessive contamination through presence of

Rene N5 63.1Ni7Cr7.5Co1.5Mo5W3Re6.5Ta0.15Hf6.2Al0.05C0.004b0.01Y

TMS 80 58.2Ni2.9Cr11.6Co1.9Mo5.8W4.9Re5.8Ta0.1Hf5.8Al0.5B3.0Ir

MC-NG 70.3Ni4Cr<0.2Co1Mo5W4Re5Ta0.1Hf0.5Ti6Al4.0Ru

Chemical composition Metal

temperature capability oC

Generation Grade

Third to

designation

First RR2000 62.5Ni10Cr15Co3Mo4Ti5.5Al1V RR2060 63Ni15Cr5Co2Mo2W5Ta2Ti5Al

CMSX2 66.2Ni8Cr4.6Co0.6Mo8W6Ta1Ti5.6Al

TUT 92 68Ni10Cr1.2Mo7W0.8Re8Ta1.2Ti5.3Al

Grade designation Chemical composition

Table 12. SC superalloys for advanced IGTs

AF56 61.4Ni12Cr3Mo5Ta4.2Ti3.4Al

SC16 70.5Ni16Cr3Mo3.5Ta3.5Ti3.5Al

CMSX3 66.1Ni8Cr4.6Co0.6Mo8W6Ta0.1Hf1Ti5.6Al

Rene N4 62Ni9.8Cr7.5Co1.5Mo6W4.8Ta0.15Hf0.5Nb3.5Ti4.2Al

CMSX4 61.7Ni6.5Cr9Co0.6Mo6W3Re6.5Ta0.1Hf1Ti5.6Al

## **7. Vanes - Oxide Dispersion Strengthened (ODS) superalloys**

A limited use exists for ODS superalloys in gas turbine engines. ODS superalloys are advanced high temperature materials which can retain useful strength up to a relatively high fraction of their melting point. This advantage is due to the uniformly dispersed, stable oxide particles which act as barriers to dislocation motion. MA754 has been in production by General Electric as a vane material since 1980. Because of its high long time elevated temperature strength, it has been extensively used for aircraft gas turbine vanes.

Fig. 3. Improvements made in the stress rupture strength of superalloys through advancements in processing technology (Schematic)

The modern high-performance gas turbine engines would not have been there but for the major advances made in superalloy development over the past 50 years, as outlined above. Excellent monographs / handbooks / technical guides are available on the subject of superalloys, covering different aspects – metallurgy, processing, properties and applications (Davis, 1997; Davis, 2000; DeHaemer, 1990; Donachie & Donachie, 2002; Durand-Charre, 1998; Reed, 2006). The reader is referred to them for more detailed information.

## **8. Reaction with the operating environment**

During operation the turbine components are subjected to environmentally induced degradation. The reaction with the environment is essentially of two types – hot corrosion and high temperature oxidation.

Hot corrosion is a rapid form of attack generally associated with alkali metal contaminants sodium and potassium – reacting with sulfur in the fuel to form molten sulfates. Two distinct forms of hot corrosion have been identified – high temperature hot corrosion occurring in the temperature range 850-950 oC and low temperature hot corrosion taking place in the range 593-760 oC. Macroscopic and microscopic characteristics and mechanisms of the two forms of hot corrosion in gas turbine components have been reviewed (Eliaz et

A limited use exists for ODS superalloys in gas turbine engines. ODS superalloys are advanced high temperature materials which can retain useful strength up to a relatively high fraction of their melting point. This advantage is due to the uniformly dispersed, stable oxide particles which act as barriers to dislocation motion. MA754 has been in production by General Electric as a vane material since 1980. Because of its high long time elevated

**7. Vanes - Oxide Dispersion Strengthened (ODS) superalloys** 

temperature strength, it has been extensively used for aircraft gas turbine vanes.

Fig. 3. Improvements made in the stress rupture strength of superalloys through

1998; Reed, 2006). The reader is referred to them for more detailed information.

The modern high-performance gas turbine engines would not have been there but for the major advances made in superalloy development over the past 50 years, as outlined above. Excellent monographs / handbooks / technical guides are available on the subject of superalloys, covering different aspects – metallurgy, processing, properties and applications (Davis, 1997; Davis, 2000; DeHaemer, 1990; Donachie & Donachie, 2002; Durand-Charre,

During operation the turbine components are subjected to environmentally induced degradation. The reaction with the environment is essentially of two types – hot corrosion

Hot corrosion is a rapid form of attack generally associated with alkali metal contaminants sodium and potassium – reacting with sulfur in the fuel to form molten sulfates. Two distinct forms of hot corrosion have been identified – high temperature hot corrosion occurring in the temperature range 850-950 oC and low temperature hot corrosion taking place in the range 593-760 oC. Macroscopic and microscopic characteristics and mechanisms of the two forms of hot corrosion in gas turbine components have been reviewed (Eliaz et

advancements in processing technology (Schematic)

**8. Reaction with the operating environment** 

and high temperature oxidation.


Table 11. SC nickel-base superalloys for blading applications in aircraft engines


Table 12. SC superalloys for advanced IGTs

al., 2002). When there is no presence of alkali metal contaminants and sulfur in the environment, high temperature oxidation dominates. Higher the temperature, more rapid is the oxidation. In this context, it is important to note the difference in the operating conditions of aircraft engines and heavy-duty gas turbines. Metal temperatures in the former are higher. The latter are subjected to excessive contamination through presence of

Materials for Gas Turbines – An Overview 307

The TBCs provide enough insulation for superalloys to operate at temperatures as much as 150 oC above their customary upper limit. TBCs are ceramics, based on ZrO2 – Y2O3 and

The ceramic coatings use an underlay of a corrosion protective layer e.g., MCrAlY that provides the oxidation resistance and necessary roughness for top coat adherence. Failures occur by the thermal expansion mismatch between the ceramic & metallic layers and by environmental attack on the bondcoat. This type of coating is used in combustion cans, transition pieces, nozzle guide vanes and also blade platforms. Improved efficiency of gas

Development of even more corrosion resistant coating materials has been intensely pursued over the last few years. In particular, improvement of oxidation resistance and thermal fatigue resistance has been a focal theme for the R&D efforts in the recent years. Development work is also on advanced TBCs to tailor their structure in such a way that they withstand thermal fatigue conditions better and give a longer life (Schilke, 2005). Development of techniques to ensure uniformity of applied coating has been another important area of research. High-velocity plasma appears to be getting established as the technique for application of overlay coatings. Much stronger bond between the coating and the workpiece and much higher coating densities can be achieved using this technique.

Increase in turbine inlet temperatures, beyond what is possible with superalloys, can be conceived if ceramic materials can be used in place of superalloys in gas turbine engine. Turbines would then operate at higher temperatures, yielding higher power with smaller engine sizes. Ceramic materials are known for their capability to withstand high temperatures. In addition they are quite tolerant to contaminants such as sodium and vanadium which are present in low cost fuels and highly corrosive to the currently used nickel-base superalloys. Ceramics are also up to 40% lighter than comparable high temperature alloys. They also cost much less – their cost is around 5% the cost of superalloys. Ceramic materials based on silicon carbide and silicon nitride were identified in 1960's as potential candidates for gas turbine application. Substantial efforts have subsequently been conducted worldwide to identify and seek solutions for key challenges: improvement in properties of candidate materials, establishing a design and life prediction methodology, generating a material database, developing cost-effective fabrication of turbine components, dimensional and non-destructive inspection, and validation of the materials and designs in rig and engine testing. Enormous technical progress has been made, but ceramic-based turbine components still have not found application in gas turbine engines, because of the problem of brittleness (Richerson, 2006). There have been efforts to improve their ductility, e.g., through addition of aluminum to ceramics. Unless the problem of brittleness is overcome satisfactorily, the use of ceramics in

During the last 30 years, extensive efforts have gone into development of intermetallic alloys for application in aircraft gas turbine engines. The primary driving force was to replace

turbine engines is realized by adopting TBCs (Gurrappa & Sambasiva Rao, 2006)

**9.1 R&D efforts in progress in the area of coatings** 

**10. Advanced materials under R&D** 

gas turbines will not be practical.

**10.2 Intermetallics** 

**10.1 Ceramics** 

produced by plasma spraying.

sodium and sulfur in the operating environment and there is potential for extensive hot corrosion damage.

Alloying with elements which have a beneficial effect of cutting down the extent of hot corrosion has been adopted as an important approach to mitigate the problem. Similarly alloying with aluminum to enable the material form its own protective layer has been adopted to prevent high temperature oxidation. However these approaches do not take very far, as there are other functional requirements which also have to be taken care. Accordingly protective coatings have been developed to ward off the degradation. Most of the superalloys used for gas turbine components receive protection from specially engineered coatings.

## **9. Coatings**

Having to perform under increasing firing temperatures and excessive contamination in the operating environment, it has become difficult to design superalloys which have the necessary creep strength on one side and the required resistance to corrosion / oxidation on the other side. It has hence become inescapable to bring coatings on to the surface of the blades to provide the necessary protection to the blades. The progress in coatings for gas turbine airfoils has been reviewed (Goward,1998). The function of the coating is to act as reservoir of elements which will form very protective and adherent oxide layers, thus protecting the underlying base material from oxidation, corrosion attack and degradation. Hot corrosion is distinctly different from the pure oxidation of an aircraft environment; it can therefore be readily appreciated that coatings for heavy duty gas turbines have different capabilities, compared to coatings for aircraft engines.

There are three basic types of coatings


The diffusion coatings have been the most common type for environmental protection of superalloys. An outer aluminide layer (CoAl or NiAl) with an enhanced oxidation resistance is developed by the reaction of Al with the Ni/Co in the base metal. In recent years extremely thin layers of noble metals such as platinum have been used to enhance the oxidation resistance of aluminides. For most stage 1 buckets, GE used a platinum-aluminum diffusion coating until 1983 (Schilke, 2004). This coating offered superior corrosion resistance to straight aluminide coatings both in burner rig tests and in field trials. Their high temperature performance is however limited by oxidation behavior of the coatings.

GE has since switched over to overlay type coatings for stage 1 buckets (Schilke, 2004) . At least one of the major constituents in a diffusion coating (generally Ni) is supplied by the base metal. An overlay coating, in contrast, has all the constituents supplied by the coating itself. The advantage is that more varied corrosion resistant compositions can be applied to optimize the performance of the coating and thickness of the coating is not limited by process considerations. The coatings are generally referred to as MCrAlY, where M stands for Ni or Co or Ni+Co. Incorporation of yttrium improves corrosion resistance. The coatings are generally applied by vacuum plasma spray process. A high temperature heat treatment is performed (1040-1120 oC) to homogenize the coating and ensure its adherence to the substrate.

sodium and sulfur in the operating environment and there is potential for extensive hot

Alloying with elements which have a beneficial effect of cutting down the extent of hot corrosion has been adopted as an important approach to mitigate the problem. Similarly alloying with aluminum to enable the material form its own protective layer has been adopted to prevent high temperature oxidation. However these approaches do not take very far, as there are other functional requirements which also have to be taken care. Accordingly protective coatings have been developed to ward off the degradation. Most of the superalloys used for gas turbine components receive protection from specially engineered

Having to perform under increasing firing temperatures and excessive contamination in the operating environment, it has become difficult to design superalloys which have the necessary creep strength on one side and the required resistance to corrosion / oxidation on the other side. It has hence become inescapable to bring coatings on to the surface of the blades to provide the necessary protection to the blades. The progress in coatings for gas turbine airfoils has been reviewed (Goward,1998). The function of the coating is to act as reservoir of elements which will form very protective and adherent oxide layers, thus protecting the underlying base material from oxidation, corrosion attack and degradation. Hot corrosion is distinctly different from the pure oxidation of an aircraft environment; it can therefore be readily appreciated that coatings for heavy duty gas turbines have different

The diffusion coatings have been the most common type for environmental protection of superalloys. An outer aluminide layer (CoAl or NiAl) with an enhanced oxidation resistance is developed by the reaction of Al with the Ni/Co in the base metal. In recent years extremely thin layers of noble metals such as platinum have been used to enhance the oxidation resistance of aluminides. For most stage 1 buckets, GE used a platinum-aluminum diffusion coating until 1983 (Schilke, 2004). This coating offered superior corrosion resistance to straight aluminide coatings both in burner rig tests and in field trials. Their high temperature performance is however limited by oxidation behavior of the coatings. GE has since switched over to overlay type coatings for stage 1 buckets (Schilke, 2004) . At least one of the major constituents in a diffusion coating (generally Ni) is supplied by the base metal. An overlay coating, in contrast, has all the constituents supplied by the coating itself. The advantage is that more varied corrosion resistant compositions can be applied to optimize the performance of the coating and thickness of the coating is not limited by process considerations. The coatings are generally referred to as MCrAlY, where M stands for Ni or Co or Ni+Co. Incorporation of yttrium improves corrosion resistance. The coatings are generally applied by vacuum plasma spray process. A high temperature heat treatment is performed (1040-1120 oC) to homogenize the coating and ensure its adherence to the

capabilities, compared to coatings for aircraft engines.

There are three basic types of coatings Aluminide (diffusion) coatings

Thermal barrier coatings (TBCs)

Overlay coatings

substrate.

corrosion damage.

coatings.

**9. Coatings** 

The TBCs provide enough insulation for superalloys to operate at temperatures as much as 150 oC above their customary upper limit. TBCs are ceramics, based on ZrO2 – Y2O3 and produced by plasma spraying.

The ceramic coatings use an underlay of a corrosion protective layer e.g., MCrAlY that provides the oxidation resistance and necessary roughness for top coat adherence. Failures occur by the thermal expansion mismatch between the ceramic & metallic layers and by environmental attack on the bondcoat. This type of coating is used in combustion cans, transition pieces, nozzle guide vanes and also blade platforms. Improved efficiency of gas turbine engines is realized by adopting TBCs (Gurrappa & Sambasiva Rao, 2006)

#### **9.1 R&D efforts in progress in the area of coatings**

Development of even more corrosion resistant coating materials has been intensely pursued over the last few years. In particular, improvement of oxidation resistance and thermal fatigue resistance has been a focal theme for the R&D efforts in the recent years. Development work is also on advanced TBCs to tailor their structure in such a way that they withstand thermal fatigue conditions better and give a longer life (Schilke, 2005). Development of techniques to ensure uniformity of applied coating has been another important area of research. High-velocity plasma appears to be getting established as the technique for application of overlay coatings. Much stronger bond between the coating and the workpiece and much higher coating densities can be achieved using this technique.

#### **10. Advanced materials under R&D**

#### **10.1 Ceramics**

Increase in turbine inlet temperatures, beyond what is possible with superalloys, can be conceived if ceramic materials can be used in place of superalloys in gas turbine engine. Turbines would then operate at higher temperatures, yielding higher power with smaller engine sizes. Ceramic materials are known for their capability to withstand high temperatures. In addition they are quite tolerant to contaminants such as sodium and vanadium which are present in low cost fuels and highly corrosive to the currently used nickel-base superalloys. Ceramics are also up to 40% lighter than comparable high temperature alloys. They also cost much less – their cost is around 5% the cost of superalloys. Ceramic materials based on silicon carbide and silicon nitride were identified in 1960's as potential candidates for gas turbine application. Substantial efforts have subsequently been conducted worldwide to identify and seek solutions for key challenges: improvement in properties of candidate materials, establishing a design and life prediction methodology, generating a material database, developing cost-effective fabrication of turbine components, dimensional and non-destructive inspection, and validation of the materials and designs in rig and engine testing. Enormous technical progress has been made, but ceramic-based turbine components still have not found application in gas turbine engines, because of the problem of brittleness (Richerson, 2006). There have been efforts to improve their ductility, e.g., through addition of aluminum to ceramics. Unless the problem of brittleness is overcome satisfactorily, the use of ceramics in gas turbines will not be practical.

#### **10.2 Intermetallics**

During the last 30 years, extensive efforts have gone into development of intermetallic alloys for application in aircraft gas turbine engines. The primary driving force was to replace

Materials for Gas Turbines – An Overview 309

way. The materials show large scatter in mechanical properties, with the result that the minimum property may be so low that the weight saving becomes negligible. They show very low tolerance to defects such as casting porosities, ceramic inclusions, machining cracks etc. There are a number of manufacturing difficulties. The production cost is much higher, compared to present technologies, particularly for complex components. Further research in alloy development and processing is required, it appears, before they get into

 NiAl is better than current Ni alloys in high temperature oxidation and in creep at very high temperatures and was actively considered for turbine components in the range 1100- 1650 oC. Serious obstacles to productionisation were faced – manufacturing difficulties, high cost of production, poor properties below 1000 oC, intrinsic brittleness. Large amount of research was done to improve the mechanical properties by developing complex multiphase intermetallic structures. The manufacturing difficulties and the brittleness still plague these materials and they are unlikely to be adopted as gas turbine materials (Lasalmanie, 2006). Many other intermetallic aluminide systems have been studied, to a lesser extent, in the context of application in aircraft engines; the materials have not taken off; the problems

Substantial progress has been made with reference to development and use of polymer matrix composites in the cold section of jet engines. GE is producing its front fan blades out

Continuous fibre reinforced titanium metal matrix composites have been the subject of intense R&D activity, as they can lead to design changes from the conventional disc and dovetail arrangement to a bladed ring with a weight reduction of about 70%. A number of processes are being evaluated for production of these composites. One of the major factors limiting the use of these composites is the fibre-matrix interaction leading to the degradation of properties. There have been many studies on the interaction between the fibre and the titanium based matrix. Different types of coatings have been tried to minimize this interaction. Titanium aluminides have received interest as matrix materials, because they show much less reaction with the fibres. Cutting down the cost of production, achieving improved performance levels and amenability to mass production are the key factors for the introduction of titanium based composites into gas turbine engines. Good overviews on the subject of titanium-based composite materials are available for the reader to get more

The manufacturers of gas turbines are continually striving to increase the operating temperatures of their engines, leading to greater thermal efficiency, and reduced emission of harmful exhaust gases. These two drivers place an ever increasing burden on the materials used in, and the design of, hot gas path components. The introduction of CMCs into hot gas path components such as combustion chamber liners has long since been identified as a possible route to the achievement of increasing operating temperatures without incurring

of epoxy resin-carbon fiber composites, resulting in substantial weight savings.

information.(Gogia, 2005; Kumphert et al., 1998; Leyans & Peters, 2003)

faced with their development are similar to those enumerated above.

flying aircraft engines.

**10.3 Composites** 

**10.3.1 Polymer matrix composites** 

**10.3.2 Titanium based metal matrix composites** 

**10.3.3 Ceramic matrix composites (CMCs)** 

the penalties associated with increased cooling air use.

nickel based alloys with a density of 8-8.5 gm/cm3 with lower density materials (4-7 gm/cm3) and gain weight saving of the engine. Titanium and nickel based aluminides were the systems which received the maximum attention.

Excellent reviews are available on the subject of development of titanium based intermetallics for aero-engine applications (Gogia, 2005; Kumphert et al., 1998; Lasalmanie, 2006; Leyans & Peters, 2003)

The titanium aluminum system offers two possibilities – the Ti3Al (2) based intermetallics and the TiAl based alloys. In the 1970's the research was centered mainly around the <sup>2</sup> alloys. Although interesting results were obtained, these materials did not get into flying engines because their fracture toughness and resistance to growth of fatigue cracks was significantly inferior to high temperature titanium alloys processed through conventional methods. They offered little or no advantage with reference to temperature capability over the alloys such as IMI 834, Ti6242, Ti1100. In their present state of development, there is not enough justification for wide spread usage of Ti3Al based intermetallics into aeroengines.

There is large volume of published work on second generation TiAl alloys, developed in 1990's. Considering the specific strength and oxidation resistance they are potential candidates to replace nickel alloys in the temperature range 650-750 oC. These alloys have also interesting properties at lower temperatures – high Young's modulus and resistance to fire and good HCF properties. All the main turbine engine manufacturers including General Electric Aircraft Engines, Pratt and Whitney and Rolls Royce have successfully gone through demonstration programmes for rotating and static engine components in the compressor, combustor, turbine and nozzle.

Stronger third generation TiAl alloys have been developed by GKSS with a wider temperature range of interest – room temperature to 850 oC, making them candidate materials also for LP compressor components (Lasalmanie, 2006).

The most recent alloy family within the titanium aluminides is represented by the orthorhombic titanium based intermetallics based on Ti2AlNb. They appear to have better toughness, higher ductility, higher specific strength and lower coefficient of thermal expansion than TiAl base intermetallics. This property profile makes orthorhombic titanium aluminides attractive for compressor casings. Even compressor discs can be considered, if their damage tolerance can be improved.

Table 13 is a compilation of representative grades from different groups of titanium aluminides and their maximum service temperature.


Table 13. Representative grades of different groups of titanium aluminides

None of the potential applications talked above are expected to enter the production phase in the immediate future. There are serious hurdles, both technical and economic, on the

nickel based alloys with a density of 8-8.5 gm/cm3 with lower density materials (4-7 gm/cm3) and gain weight saving of the engine. Titanium and nickel based aluminides were

Excellent reviews are available on the subject of development of titanium based intermetallics for aero-engine applications (Gogia, 2005; Kumphert et al., 1998; Lasalmanie,

The titanium aluminum system offers two possibilities – the Ti3Al (2) based intermetallics and the TiAl based alloys. In the 1970's the research was centered mainly around the <sup>2</sup> alloys. Although interesting results were obtained, these materials did not get into flying engines because their fracture toughness and resistance to growth of fatigue cracks was significantly inferior to high temperature titanium alloys processed through conventional methods. They offered little or no advantage with reference to temperature capability over the alloys such as IMI 834, Ti6242, Ti1100. In their present state of development, there is not enough justification for wide spread usage of Ti3Al based intermetallics into aeroengines. There is large volume of published work on second generation TiAl alloys, developed in 1990's. Considering the specific strength and oxidation resistance they are potential candidates to replace nickel alloys in the temperature range 650-750 oC. These alloys have also interesting properties at lower temperatures – high Young's modulus and resistance to fire and good HCF properties. All the main turbine engine manufacturers including General Electric Aircraft Engines, Pratt and Whitney and Rolls Royce have successfully gone through demonstration programmes for rotating and static engine components in the

Stronger third generation TiAl alloys have been developed by GKSS with a wider temperature range of interest – room temperature to 850 oC, making them candidate

The most recent alloy family within the titanium aluminides is represented by the orthorhombic titanium based intermetallics based on Ti2AlNb. They appear to have better toughness, higher ductility, higher specific strength and lower coefficient of thermal expansion than TiAl base intermetallics. This property profile makes orthorhombic titanium aluminides attractive for compressor casings. Even compressor discs can be considered, if

Table 13 is a compilation of representative grades from different groups of titanium

Ti-48Al-2Cr-2Nb 800

Ti-45Al-8Nb 850

None of the potential applications talked above are expected to enter the production phase in the immediate future. There are serious hurdles, both technical and economic, on the

Maximum service temperature oC

550

Chemical composition of a typical

the systems which received the maximum attention.

compressor, combustor, turbine and nozzle.

their damage tolerance can be improved.

Group of titanium



aluminide

aluminides and their maximum service temperature.

materials also for LP compressor components (Lasalmanie, 2006).

grade(s) (At%)

Ti2AlNb based Ti-22Al-25Nb 700 Table 13. Representative grades of different groups of titanium aluminides

Ti3Al based Ti-24Al-11Nb, Ti-25Al-10Nb-3V-1Mo

2006; Leyans & Peters, 2003)

way. The materials show large scatter in mechanical properties, with the result that the minimum property may be so low that the weight saving becomes negligible. They show very low tolerance to defects such as casting porosities, ceramic inclusions, machining cracks etc. There are a number of manufacturing difficulties. The production cost is much higher, compared to present technologies, particularly for complex components. Further research in alloy development and processing is required, it appears, before they get into flying aircraft engines.

 NiAl is better than current Ni alloys in high temperature oxidation and in creep at very high temperatures and was actively considered for turbine components in the range 1100- 1650 oC. Serious obstacles to productionisation were faced – manufacturing difficulties, high cost of production, poor properties below 1000 oC, intrinsic brittleness. Large amount of research was done to improve the mechanical properties by developing complex multiphase intermetallic structures. The manufacturing difficulties and the brittleness still plague these materials and they are unlikely to be adopted as gas turbine materials (Lasalmanie, 2006).

Many other intermetallic aluminide systems have been studied, to a lesser extent, in the context of application in aircraft engines; the materials have not taken off; the problems faced with their development are similar to those enumerated above.

#### **10.3 Composites**

#### **10.3.1 Polymer matrix composites**

Substantial progress has been made with reference to development and use of polymer matrix composites in the cold section of jet engines. GE is producing its front fan blades out of epoxy resin-carbon fiber composites, resulting in substantial weight savings.

#### **10.3.2 Titanium based metal matrix composites**

Continuous fibre reinforced titanium metal matrix composites have been the subject of intense R&D activity, as they can lead to design changes from the conventional disc and dovetail arrangement to a bladed ring with a weight reduction of about 70%. A number of processes are being evaluated for production of these composites. One of the major factors limiting the use of these composites is the fibre-matrix interaction leading to the degradation of properties. There have been many studies on the interaction between the fibre and the titanium based matrix. Different types of coatings have been tried to minimize this interaction. Titanium aluminides have received interest as matrix materials, because they show much less reaction with the fibres. Cutting down the cost of production, achieving improved performance levels and amenability to mass production are the key factors for the introduction of titanium based composites into gas turbine engines. Good overviews on the subject of titanium-based composite materials are available for the reader to get more information.(Gogia, 2005; Kumphert et al., 1998; Leyans & Peters, 2003)

#### **10.3.3 Ceramic matrix composites (CMCs)**

The manufacturers of gas turbines are continually striving to increase the operating temperatures of their engines, leading to greater thermal efficiency, and reduced emission of harmful exhaust gases. These two drivers place an ever increasing burden on the materials used in, and the design of, hot gas path components. The introduction of CMCs into hot gas path components such as combustion chamber liners has long since been identified as a possible route to the achievement of increasing operating temperatures without incurring the penalties associated with increased cooling air use.

Materials for Gas Turbines – An Overview 311

Molybdenum alloys are currently used as components for ultra-high-temperature applications under protective atmosphere, taking advantage of their high melting point and, very good mechanical and creep strength. However, they suffer from severe oxidation in air above around 500 °C. Compositions based on Mo-9Si-8B have shown promise as structural materials for applications in excess of 1,100 °C in air (Heilmeier et al., 2009). The silicide and boride phases serve to provide oxidation resistance. Experiments to increase the oxidation resistance through alloying with Cr have shown promising results. Alloying with the reactive element Zr was also found to bring down the rate of oxidation. Creep resistance of the Mo-9Si-8B composition was found to be comparable or even superior to that of CMSX4 over the temperature range 1100-1200 oC; there is scope for R&D to further increase the

Platinum-based alloys possess the potential to be used at temperatures up to 1,700 °C. Despite their high prices, they are attractive for some gas turbine applications due to their exceptional resistance to oxidation, high melting points, ductility, thermal shock resistance, and thermal conductivity. They are envisaged to have potential for highly thermally loaded, but non-rotating parts in gas turbines (Alven, 2004). Work is in progress for development of platinum-based alloys with microstructures similar to those seen in commercial nickel-based alloys (Yamabe-Mitarai, 2004; Vorberg et al., 2004). Numerous research groups are currently involved in the study of these alloys. Significantly more data needs to be generated before

Turbine entry temperature has increased by ~500 oC over last 6 decades and about 150 oC of that is due to improved superalloys and introduction of DS / SC technologies for blade casting. Advanced thermal barrier ceramic coatings on platform and full airfoil have contributed to another about 100 oC of this improvement. The developments in gas turbine materials and coatings have been largely due to increasing demands placed by the aircraft sector – higher engine thrust, thrust to weight ratio and fuel efficiency – necessitating higher operating temperatures and pressures. The land based industrial gas turbine industry has placed its own demands on materials, bringing in resistance to hot corrosion as an important requirement. Several SC superalloy compositions have been developed for aircraft gas turbines on one side and land based gas turbines on the other side. Partial solutioning has been adopted in a number of SC IGT alloys to avoid incipient melting and control the extent of recrystallisation. Intense R&D is also going on development of advanced materials for gas turbine engine application – intermetallics, ceramics, composites, chromium / molybdenum / platinum based materials to improve the engine efficiency and bring down the harmful emissions. Major improvements in the coating technology have also been achieved. Present day coatings last 10-20 times longer than the coatings used in the late 90's. As much as 100% improvement is now being achieved in the blade life in the field through the process of coating. TBCs are being used in the first few stages in all advanced gas turbines. Intense R&D is underway to improve the thermal fatigue of the TBC's and thereby increase their life. This includes development of techniques for production of

the platinum-based alloys are ready for the designers of gas turbine engines.

**10.5 Molybdenum based alloys as gas turbine materials** 

creep strength.

**11. Conclusion** 

uniform and high density coatings.

**10.6 Platinum based alloys** 

**SiC**-matrix composites appear to be highly tailorable materials suitable to gas turbine application at high temperatures. Melt infiltrated (MI) SiC/SiC composites are particularly attractive for gas turbine applications because of their high thermal conductivity, excellent thermal shock resistance, creep resistance, and oxidation resistance compared to other CMCs. They are tough, although their constituents are intrinsically brittle, when the fibermatrix bonding is properly optimized through the use of a thin interphase deposited on the fibers prior to the infiltration of the matrix. They display good mechanical properties at high temperatures when prepared from stable fibers, as well as a high thermal conductivity if their residual porosity is low enough. The matrix composition can also be tailored to improve the oxidation resistance of the composites.

Replacing the superalloys by light, tough, refractory and creep resistant SiC-matrix composites will permit a significant increase of service temperature and hence an increase of the engine efficiency, a reduction of the NOx/CO emission (through an optimization of the fuel/air ratio), a simplification of the part design and a weight saving (typically, 30–50%). However, their use still raises a number of questions dealing with their durability, reliability, manufacture, design and cost. Presently, the development of SiC-matrix composites is limited mostly to non rotating parts including combustor liners, after-burner components (exhaust cone and flame holder) and exhaust nozzles (outer and inner flaps) in military aerojet engines, as well as combustor liners of large size in stationary gas turbine for electrical power/steam cogeneration. The GE Rolls-Royce Fighter Engine Team's F136 development engine for the Joint Strike Fighter (JSF) contains third-stage, low-pressure turbine vanes made by GE from CMCs.

In summary, although substantial progress has been made, significant risks and challenges still remain before these composites can be commercialized for gas turbine components. The reader is referred to an overview on design, preparation and properties of non-oxide CMCs for application in gas turbine engines (Naslain, 2004)

#### **10.4 Chromium based alloys as gas turbine materials**

The increased efficiency associated with higher operating temperatures in gas turbines has prompted designers to search for new materials that can be used at temperatures above the useful limit of nickel-based superalloys. Chromium based alloys have been considered as a possible base for alloy systems due to their high melting point, good oxidation resistance, low density (20% less than most nickel-based superalloys), and high thermal conductivity (two to four times higher than most superalloys). Considerable effort was made in the past, to explore the possibility of developing chromium-based alloys for high-temperature applications such as in jet engines. Two major disadvantages came in the way of their commercial exploitation. First, chromium alloys have a high ductile-to-brittle transition temperature. Second, chromium exhibits further embrittlement resulting from nitrogen contamination during high-temperature air exposure. Since the late 1970s, chromium alloys received very little attention. In recent years there is a revival of interest in these alloys (Gu et al., 2004). Aims of the recent researches have been (i) improvement in high temperature strength (ii) protection from nitridation / oxidation embrittlement (iii) improvement in the impact ductility at ambient temperature. There have been some encouraging findings – (i) trace additions of silver can improve room temperature ductility significantly (ii) strengthening with intermetallics can improve high temperature strength (iii) substantial progress has been made in the area of high temperature coatings (Gu et al.,2004). It is still a long way, however, before the engine manufacturers get interested in these alloys.

**SiC**-matrix composites appear to be highly tailorable materials suitable to gas turbine application at high temperatures. Melt infiltrated (MI) SiC/SiC composites are particularly attractive for gas turbine applications because of their high thermal conductivity, excellent thermal shock resistance, creep resistance, and oxidation resistance compared to other CMCs. They are tough, although their constituents are intrinsically brittle, when the fibermatrix bonding is properly optimized through the use of a thin interphase deposited on the fibers prior to the infiltration of the matrix. They display good mechanical properties at high temperatures when prepared from stable fibers, as well as a high thermal conductivity if their residual porosity is low enough. The matrix composition can also be tailored to

Replacing the superalloys by light, tough, refractory and creep resistant SiC-matrix composites will permit a significant increase of service temperature and hence an increase of the engine efficiency, a reduction of the NOx/CO emission (through an optimization of the fuel/air ratio), a simplification of the part design and a weight saving (typically, 30–50%). However, their use still raises a number of questions dealing with their durability, reliability, manufacture, design and cost. Presently, the development of SiC-matrix composites is limited mostly to non rotating parts including combustor liners, after-burner components (exhaust cone and flame holder) and exhaust nozzles (outer and inner flaps) in military aerojet engines, as well as combustor liners of large size in stationary gas turbine for electrical power/steam cogeneration. The GE Rolls-Royce Fighter Engine Team's F136 development engine for the Joint Strike Fighter (JSF) contains third-stage, low-pressure

In summary, although substantial progress has been made, significant risks and challenges still remain before these composites can be commercialized for gas turbine components. The reader is referred to an overview on design, preparation and properties of non-oxide CMCs

The increased efficiency associated with higher operating temperatures in gas turbines has prompted designers to search for new materials that can be used at temperatures above the useful limit of nickel-based superalloys. Chromium based alloys have been considered as a possible base for alloy systems due to their high melting point, good oxidation resistance, low density (20% less than most nickel-based superalloys), and high thermal conductivity (two to four times higher than most superalloys). Considerable effort was made in the past, to explore the possibility of developing chromium-based alloys for high-temperature applications such as in jet engines. Two major disadvantages came in the way of their commercial exploitation. First, chromium alloys have a high ductile-to-brittle transition temperature. Second, chromium exhibits further embrittlement resulting from nitrogen contamination during high-temperature air exposure. Since the late 1970s, chromium alloys received very little attention. In recent years there is a revival of interest in these alloys (Gu et al., 2004). Aims of the recent researches have been (i) improvement in high temperature strength (ii) protection from nitridation / oxidation embrittlement (iii) improvement in the impact ductility at ambient temperature. There have been some encouraging findings – (i) trace additions of silver can improve room temperature ductility significantly (ii) strengthening with intermetallics can improve high temperature strength (iii) substantial progress has been made in the area of high temperature coatings (Gu et al.,2004). It is still a

long way, however, before the engine manufacturers get interested in these alloys.

improve the oxidation resistance of the composites.

turbine vanes made by GE from CMCs.

for application in gas turbine engines (Naslain, 2004)

**10.4 Chromium based alloys as gas turbine materials** 

#### **10.5 Molybdenum based alloys as gas turbine materials**

Molybdenum alloys are currently used as components for ultra-high-temperature applications under protective atmosphere, taking advantage of their high melting point and, very good mechanical and creep strength. However, they suffer from severe oxidation in air above around 500 °C. Compositions based on Mo-9Si-8B have shown promise as structural materials for applications in excess of 1,100 °C in air (Heilmeier et al., 2009). The silicide and boride phases serve to provide oxidation resistance. Experiments to increase the oxidation resistance through alloying with Cr have shown promising results. Alloying with the reactive element Zr was also found to bring down the rate of oxidation. Creep resistance of the Mo-9Si-8B composition was found to be comparable or even superior to that of CMSX4 over the temperature range 1100-1200 oC; there is scope for R&D to further increase the creep strength.

#### **10.6 Platinum based alloys**

Platinum-based alloys possess the potential to be used at temperatures up to 1,700 °C. Despite their high prices, they are attractive for some gas turbine applications due to their exceptional resistance to oxidation, high melting points, ductility, thermal shock resistance, and thermal conductivity. They are envisaged to have potential for highly thermally loaded, but non-rotating parts in gas turbines (Alven, 2004). Work is in progress for development of platinum-based alloys with microstructures similar to those seen in commercial nickel-based alloys (Yamabe-Mitarai, 2004; Vorberg et al., 2004). Numerous research groups are currently involved in the study of these alloys. Significantly more data needs to be generated before the platinum-based alloys are ready for the designers of gas turbine engines.

#### **11. Conclusion**

Turbine entry temperature has increased by ~500 oC over last 6 decades and about 150 oC of that is due to improved superalloys and introduction of DS / SC technologies for blade casting. Advanced thermal barrier ceramic coatings on platform and full airfoil have contributed to another about 100 oC of this improvement. The developments in gas turbine materials and coatings have been largely due to increasing demands placed by the aircraft sector – higher engine thrust, thrust to weight ratio and fuel efficiency – necessitating higher operating temperatures and pressures. The land based industrial gas turbine industry has placed its own demands on materials, bringing in resistance to hot corrosion as an important requirement. Several SC superalloy compositions have been developed for aircraft gas turbines on one side and land based gas turbines on the other side. Partial solutioning has been adopted in a number of SC IGT alloys to avoid incipient melting and control the extent of recrystallisation. Intense R&D is also going on development of advanced materials for gas turbine engine application – intermetallics, ceramics, composites, chromium / molybdenum / platinum based materials to improve the engine efficiency and bring down the harmful emissions. Major improvements in the coating technology have also been achieved. Present day coatings last 10-20 times longer than the coatings used in the late 90's. As much as 100% improvement is now being achieved in the blade life in the field through the process of coating. TBCs are being used in the first few stages in all advanced gas turbines. Intense R&D is underway to improve the thermal fatigue of the TBC's and thereby increase their life. This includes development of techniques for production of uniform and high density coatings.

Materials for Gas Turbines – An Overview 313

Gu, Y.F.; Harada, H. & Ro,Y. (2004). Chromium and Chromium-Based Alloys: Problems and

Gurrappa, I. & Sambasiva Rao, A. (2006). Thermal Barrier Coatings for Enhanced Efficiency

Heilmeier, M.; Krüger, M.; Saage, H. & Rösler, J. Metallic Materials for Structural

Kumphert, J.; Peters, M. & Keysser, W.A. The Potential of Advanced Materials on Structural

Lasalmanie, A. (2006). Intermetallics: Why is it so Difficult to Introduce them in Gas

Leyans, C & Peters, M. (2003). Titanium and Titanium Alloys: Fundamentals and

Loria, E.D. (1989). Proceedings of Conference on Superalloy 718 – Metallurgy and

Loria, E.D. (1991). Proceedings of Conference on Superalloy 718, 625 and various

Loria, E.D. (1994). Superalloys 718, 625, and Various Derivatives, The Minerals, Metals &

Loria, E.D. (1997). Superalloys 718, 625, 706 and Various Derivatives, The Minerals, Metals &

Loria, E.D. (2001). Proceedings of the Fifth International Conference on Superalloys 718, 625, 706 and Various Derivatives, TMS, ISBN 0-87339-510-7, Warrendale, PA Loria, E.D. (2005). Proceedings of Sixth International Symposium on Superalloys 718, 625,

Naslain, R. (2004) Design, preparation and properties of non-oxide CMCs for application in

Reed, R.C. (2006). Superalloys: Fundamentals and Applications, Cambridge University

Richerson, D.W. (2006). Historical review of addressing the challenges of use of ceramic

Schilke, P.W. (2004) Advanced Gas Turbine Materials and Coatings, 18.04.2011, Available from

Vorberg, S.; Wenderoth, M.; Fischer, B.; Glatzel, U. & Völkl, R. Pt-Al-Cr-Ni Superalloys: Heat Treatment and Microstructure, JOM, Vol.56, No.9, pp40-43 ISSN 1047-4838

engines and nuclear reactors: an overview. Composites Science and Technology,

components in gas turbines, Proceedings of 2006 ASME 51st Turbo Expo, ISBN

www.gepower,com/prod\_serv/products/tech\_docs/en/downloads/ger3569g.pdf

706 and Derivatives, TMS, ISBN 978-0-87339-602-8, Warrendale, PA

Applications, TMS, ISBN 0-87339-097-0, Warrendale, PA

Derivatives, TMS, ISBN 0-87339-173X, Warrendale, PA

Materials Society, ISBN 0-87339-235-3, Warrendale, PA

Materials Society, ISBN 0-87339-376-7, Warrendale, PA

Vol.64, pp155-170, ISSN 0266-3538

0791842371, Barcelona, May 2006

Press, ISBN 9780521859042

28-33, ISSN 1047-4838

3029 ISSN 0257-8972

0005-8, Toulouse, France, May 11-15, 1998

1047-4838

ISSN 0966-9795

FRG

Possibilities for High Temperature Service. JOM, Vol.56, No.9 (September 2004), pp.

of Gas Turbine Engines, Surface & Coatings Technology, Vol.201, (2006), pp3016-

Applications Beyond Nickel-based Superalloys JOM, Vol.61, No.7, pp61-67 ISSN

design of Future Aircraft engines, Proceedings of RTO AVT Symposium on " Design, Principles and Methods for Aircraft gas Turbine engines", ISBN 92-837-

Turbine Engines?, In: Intermetalics, Vol.14, No.10-11, (October 2006), pp1123-1129,

Applications, Wiley-VCH Verlag GmbH & Co, ISBN 9783527305346, Weinheim,

## **12. Acknowledgment**

The author is grateful to the Management of VIT University for their kind consent to publish this Chapter. He is also indebted to Ms. Brunda, his wife, for all the support he received from her in preparing this manuscript.

## **13. References**


The author is grateful to the Management of VIT University for their kind consent to publish this Chapter. He is also indebted to Ms. Brunda, his wife, for all the support he received

Alven, D.A. Refractory- and Precious Metal- Based Superalloys, JOM, Vol.56, No.9, pp 27

Bayer, R.R. An Overview on the Use of Titanium in the Aerospace Industry, Materials Science and Engineering A, Vol.A213, (1996), pp103-114 ISSN 0921-5093 Boyce, M.P. (2006). Gas Turbine Engineering Handbook, (Third Edition), Gulf Professional

Coutsouradis, D.; Davin, A. & Lamberigts, M.(1987) Cobalt-based Superalloys for

Das, N. (2010). Advances in Nickel-base Cast Superalloys, Transactions of The Indian Institute of Metals, Vol.63, No.2-3, (April-June 2010), pp265-274, ISSN 0972-2615 Davis, J.R. (Ed.). (1997). Heat Resistant Materials, ASM Speciality Handbook, ASM

Davis, J.R. (2000). Nickel, Cobalt and their Alloys, ASM Speciality Handbook, ASM

DeHaemer, M.J. (1990). ASM Handbook, Volume 1, Properties and Selection: Irons, Steels,

Superalloys, ASM International, ISBN 0-87170-377-7, Materials Park, OH Donachie, M.J. (August 2000). Titanium, A Technical Guide, 2 Edition, ASM International,

Donachie, M.J.; & Donachie, S.J. (2002). Superalloys A Technical Guide, (Second Edition),

Durand-Charre, M. (1998). The Microstructure of Superalloys, CRC Press, ISBN 9056990977 Eliaz, N.; Shemesh, G. & Latanision, R.M. (2002). Hot Corrosion in Gas Turbine Components. Engineering Failure Analysis, Vol.9, pp.31-43, ISSN 1350-6307 Furrer, D. & Fecht, H. (1999). Ni-Based Superalloys for Turbine Discs. JOM, Vol.51, No.1,

Gibbons, T.B. Superalloys in Modern Power Generation Application, Materials Science and

Gogia, A.K. (2005). High Temperature Titanium Alloys, Defence Science Journal, Vol.55,

Goward, G.W. (1998). Progress in coatings for gas turbine airfoils. Surface and Coatings

and High Performance Alloys, Wrought and Powder Metallurgy (P/M)

ASM International, ISBN 978-0-87170-749-9, The Materials Society, Materials Park,

Superalloys , TMS, ISBN 0873397282, Warrendale, PA

International, ISBN 978-0-87170-596-9, Materials Park, OH

International, ISBN 0-87170-685-7, Materials Park, OH

Publishing, ISBN 0-88415-732-6, OxfordCaron, P.; Green, K.A. & Reed, R.C. (Eds.). (2008) Superalloys: Proceedings of the Eleventh International Symposium on

Application in Gas Turbines, Materials Science and Engineering, Vol.88, pp11-19,

**12. Acknowledgment** 

**13. References** 

from her in preparing this manuscript.

ISSN 1047-4838

ISSN 0921-5093

ISBN 0871706865, USA

(January 1999), pp. 14-17, ISSN 1047-4838

No.2, (April 2005), pp149-173, ISSN 0011748X

Technology, Vol.108-109, pp. 73-79, ISSN 0257-8972

Technology, Vol.25, No.2, pp129-135

Ohio


**14** 

*UK* 

Mark Whittaker *Swansea University* 

**Titanium in the Gas Turbine Engine** 

The development of the gas turbine engine over the past 60 years has been mirrored by the success of the titanium industry, with a clear symbiant relationship existing between the two industries. Immediately apparent in the early days of the evolution of the gas turbine was the need for a material which could provide the strength required for component operation, whilst at the same time providing a low enough density to allow for successful flight applications. Whilst aluminium based alloys offer an excellent strength to weight ratio, their operation is limited to temperatures below approximately 130⁰C, reducing possible applications within the gas turbine to a minimum. 300 series stainless steels offer a similar strength to most conventional titanium alloys, but come with a significant density penalty of over 50% and, whilst offering reductions in cost, do not provide significant

Titanium however, has long been viewed as having a desirable balance of properties for applications towards the front end of the gas turbine engine (i.e. fan discs/blades, compressor discs/blades, along with other smaller components). Titanium has a density of 4.5g/cm3 (which, apart from a limited number of alloys such as Ti811, does not vary significantly in alloys considered for aerospace applications) which is higher than aluminium, but lower than nickel and steel alloys. Titanium is allotropic with a HCP lattice ( phase) stable to 882⁰C, transforming to a BCC (β phase) lattice above this temperature. Alloying elements act to stabilize either of these phases (Al, Sn for example stabilize the alpha phase, whereas Mo, V, Cr stabilize the beta phase) meaning that the transformation temperature can be altered, and subsequently the proportions of each phase existing at room temperature can be varied. The morphology of these phases may however vary, dependent on the process history, with alpha phase material being classed as primary alpha (persisting during heat treatment in the phase field) or secondary alpha (structures arising from the → phase transformation). This allows for the development of a range of bimodal microstructures which provide titanium alloys with inherent strength and also allows for further refinement of properties through various heat treatment and processing regimes. For example designers requiring creep strength and good elevated temperature properties may choose to opt for alloys with more alpha stabilizers (alpha or near alpha alloys), whereas metastable beta alloys, which are heavily beta stabilized offer improved forgeability. Alphabeta alloys contain a more balanced mix of stabilizers and are widely used due their balance of properties. Ti6-4 (Ti-6Al-4V) for example has been a stalwart of the titanium industry since the 1950s due to its good weldability, relatively high strength and good fatigue

**1. Introduction** 

properties.

benefits in terms of operating temperatures.


## **Titanium in the Gas Turbine Engine**

Mark Whittaker *Swansea University UK* 

#### **1. Introduction**

314 Advances in Gas Turbine Technology

Wright, I.G. & Gibbons, T.B.(2007) Recent developments in gas turbine materials and

Yamabe-Mittarai, Y. (2004). Platinum-Group-Metals-Based Intermetallics as High-Temperature Structural Materials, JOM, Vol.56, No.9, pp ISSN 1047-4838

Hydrogen Energy, Vol.32, pp 3610-3621, ISSN 0360-3199

technology and their implications for syngas firing. International Journal of

The development of the gas turbine engine over the past 60 years has been mirrored by the success of the titanium industry, with a clear symbiant relationship existing between the two industries. Immediately apparent in the early days of the evolution of the gas turbine was the need for a material which could provide the strength required for component operation, whilst at the same time providing a low enough density to allow for successful flight applications. Whilst aluminium based alloys offer an excellent strength to weight ratio, their operation is limited to temperatures below approximately 130⁰C, reducing possible applications within the gas turbine to a minimum. 300 series stainless steels offer a similar strength to most conventional titanium alloys, but come with a significant density penalty of over 50% and, whilst offering reductions in cost, do not provide significant benefits in terms of operating temperatures.

Titanium however, has long been viewed as having a desirable balance of properties for applications towards the front end of the gas turbine engine (i.e. fan discs/blades, compressor discs/blades, along with other smaller components). Titanium has a density of 4.5g/cm3 (which, apart from a limited number of alloys such as Ti811, does not vary significantly in alloys considered for aerospace applications) which is higher than aluminium, but lower than nickel and steel alloys. Titanium is allotropic with a HCP lattice ( phase) stable to 882⁰C, transforming to a BCC (β phase) lattice above this temperature. Alloying elements act to stabilize either of these phases (Al, Sn for example stabilize the alpha phase, whereas Mo, V, Cr stabilize the beta phase) meaning that the transformation temperature can be altered, and subsequently the proportions of each phase existing at room temperature can be varied. The morphology of these phases may however vary, dependent on the process history, with alpha phase material being classed as primary alpha (persisting during heat treatment in the phase field) or secondary alpha (structures arising from the → phase transformation). This allows for the development of a range of bimodal microstructures which provide titanium alloys with inherent strength and also allows for further refinement of properties through various heat treatment and processing regimes. For example designers requiring creep strength and good elevated temperature properties may choose to opt for alloys with more alpha stabilizers (alpha or near alpha alloys), whereas metastable beta alloys, which are heavily beta stabilized offer improved forgeability. Alphabeta alloys contain a more balanced mix of stabilizers and are widely used due their balance of properties. Ti6-4 (Ti-6Al-4V) for example has been a stalwart of the titanium industry since the 1950s due to its good weldability, relatively high strength and good fatigue properties.

Titanium in the Gas Turbine Engine 317

The mechanical test data detailed in the following section was produced at Swansea University as part of a number of academic programmes funded either in part or fully by

The highly anisotropic nature of the HCP phase lattice clearly gives rise to potential variations in mechanical properties of titanium alloys. This can be demonstrated by early work on the subject (Zarkades & Larson, 1970) which illustrates the variation of Young's

conservative safety factors.

EPSRC, Rolls-Royce plc and TIMET UK.

**3.1 Crystallographic texture (Ti6-4)** 

**3. Potential developments in titanium alloys** 

modulus with crystal orientation, as shown by Figure 1.

Fig. 1. The dependence of Young's modulus on increasing angle with c-axis.

understand the effects that this 'crystallographic texture' may impart to the alloy.

It is clear that the due to the denser atomic packing along the c-axis (<0001> direction) a maximum value of the Young's modulus occurs in this direction, which shows a gradual decrease with increasing angle made with the c-axis. It is also apparent that this variation can account for differences of up to 40GPa, which significantly alters the mechanical response of the alloy. However, these changes do not only affect elastic deformation of these alloys, but also the plastic deformation, in terms of the orientation of slip planes with relatively low critical resolved shear stress (CRSS) values. As such it is necessary to

**2. Generation of test data** 

variations in the mechanical properties. These variations are rarely simple, with for example, the residual stress conditions being dependent on temperature and dislocation density and a lack of understanding may result in the requirement for overly

The continued extensive use of titanium in the gas turbine, however, is constantly under threat, with potential new technologies showing promise at the lower temperature end of the gas turbine (fan blades) and at higher temperatures (HP compressor). Composite fan blades offer designers the opportunity for further weight reductions and, although concerns may still be raised over impact resistance, these materials offer great incentives to replace Ti6-4 as a fan blade material. Furthermore, in order to operate gas turbine engines more efficiently, it is necessary to continue to raise the temperature capability of the engine so that fuel can be burnt at temperatures closer to the stoichiometric value and also to achieve a higher compression ratio of the gas through the compressor stages. The second of these issues impacts significantly on the materials utilised within the compressor, since a higher pressure ratio will result in an increase in temperature where high temperature titanium alloys, such as Ti834 are utilized as disc materials. Ti834 was developed for high temperature applications and shows exceptional properties up to approximately 630°C. However, as operating temperatures at the disc rim increase, it clearly becomes necessary to consider viable alternatives, since Ti834 is limited by environmental degradation in the form of alpha case formation. Unfortunately however, no other titanium alloys show appropriate properties at temperatures higher than this, and designers are forced to consider other alloy systems. Nickel alloys are most commonly utilised as replacement alloys, although a significant density penalty is attached, with typical intermediate temperature alloys such as IN718 having a density of 8.19g/cm3. Clearly on this basis designers will opt to utilise high temperature titanium alloys as far as possible to avoid these increases in weight. With pressure being applied to the use of titanium in the gas turbine from materials offering either lower density or a higher temperature capability, particularly in the case of alloys, it is clear that it is necessary for the titanium industry to further develop the class of alloys for in service applications. However, it would be unrealistic to assume that advancements in terms of mechanical properties can be maintained through the development of new, improved alloys as has been the case in the past. Whilst alloy refinement and development may still contribute to these advances, it is clear that alternative, innovative approaches are necessary to extract further improvements in performance. The current work seeks to demonstrate three areas which may offer opportunities for advancement:-


The continued extensive use of titanium in the gas turbine, however, is constantly under threat, with potential new technologies showing promise at the lower temperature end of the gas turbine (fan blades) and at higher temperatures (HP compressor). Composite fan blades offer designers the opportunity for further weight reductions and, although concerns may still be raised over impact resistance, these materials offer great incentives to replace Ti6-4 as a fan blade material. Furthermore, in order to operate gas turbine engines more efficiently, it is necessary to continue to raise the temperature capability of the engine so that fuel can be burnt at temperatures closer to the stoichiometric value and also to achieve a higher compression ratio of the gas through the compressor stages. The second of these issues impacts significantly on the materials utilised within the compressor, since a higher pressure ratio will result in an increase in temperature where high temperature titanium alloys, such as Ti834 are utilized as disc materials. Ti834 was developed for high temperature applications and shows exceptional properties up to approximately 630°C. However, as operating temperatures at the disc rim increase, it clearly becomes necessary to consider viable alternatives, since Ti834 is limited by environmental degradation in the form of alpha case formation. Unfortunately however, no other titanium alloys show appropriate properties at temperatures higher than this, and designers are forced to consider other alloy systems. Nickel alloys are most commonly utilised as replacement alloys, although a significant density penalty is attached, with typical intermediate temperature alloys such as IN718 having a density of 8.19g/cm3. Clearly on this basis designers will opt to utilise high temperature titanium alloys as far as possible to avoid these increases in weight. With pressure being applied to the use of titanium in the gas turbine from materials offering either lower density or a higher temperature capability, particularly in the case of alloys, it is clear that it is necessary for the titanium industry to further develop the class of alloys for in service applications. However, it would be unrealistic to assume that advancements in terms of mechanical properties can be maintained through the development of new, improved alloys as has been the case in the past. Whilst alloy refinement and development may still contribute to these advances, it is clear that alternative, innovative approaches are necessary to extract further improvements in performance. The current work seeks to demonstrate three areas which may offer

i. The harnessing of crystallographic texture. The orientation of the HCP alpha phase lattice can result in significant anisotropy in mechanical properties. Whilst original research into titanium alloys sought to minimise this anisotropy, more recently it has been acknowledged that it should be possible to match the best properties from the

ii. The development of appropriate fatigue lifing methods for alloys operating in conditions where additional failure mechanisms such as creep and environmental damage interact with fatigue. Whereas the traditional view of engineers was to ensure that materials were operating at temperatures below where significant creep deformation would occur, improved lifing techniques enable accurate predictions to be

iii. An improved understanding of the effect of processing on alloys which may prove extremely sensitive to issues such as residual stress. It is acknowledged that processing of components is likely to produce a less than perfect surface condition, and that the effects of surface roughness, residual stress and 'damage' to the material may cause

textured material with the most demanding loading directions.

made in these temperature regimes, allowing for safe operation.

opportunities for advancement:-

variations in the mechanical properties. These variations are rarely simple, with for example, the residual stress conditions being dependent on temperature and dislocation density and a lack of understanding may result in the requirement for overly conservative safety factors.

## **2. Generation of test data**

The mechanical test data detailed in the following section was produced at Swansea University as part of a number of academic programmes funded either in part or fully by EPSRC, Rolls-Royce plc and TIMET UK.

#### **3. Potential developments in titanium alloys**

#### **3.1 Crystallographic texture (Ti6-4)**

The highly anisotropic nature of the HCP phase lattice clearly gives rise to potential variations in mechanical properties of titanium alloys. This can be demonstrated by early work on the subject (Zarkades & Larson, 1970) which illustrates the variation of Young's modulus with crystal orientation, as shown by Figure 1.

Fig. 1. The dependence of Young's modulus on increasing angle with c-axis.

It is clear that the due to the denser atomic packing along the c-axis (<0001> direction) a maximum value of the Young's modulus occurs in this direction, which shows a gradual decrease with increasing angle made with the c-axis. It is also apparent that this variation can account for differences of up to 40GPa, which significantly alters the mechanical response of the alloy. However, these changes do not only affect elastic deformation of these alloys, but also the plastic deformation, in terms of the orientation of slip planes with relatively low critical resolved shear stress (CRSS) values. As such it is necessary to understand the effects that this 'crystallographic texture' may impart to the alloy.

Titanium in the Gas Turbine Engine 319

Figure 3 illustrates the type of effect that may be seen in textured alloys. Under strain control testing at R=0 (20°C, 1-1-1-1 second trapezoid cycle) it can be seen that specimens taken parallel to the rolling direction of the material show longer fatigue lives than counterparts taken from the transverse direction of the plate (TD), i.e. at 90° to the rolling direction. Initially this result may be considered counterintuitive; the increased density of basal planes perpendicular to the TD direction results in a higher modulus in the TD specimens, and also higher yield stress and UTS values. However, in order to understand this effect, it is necessary to consider the behaviour of the material under strain control, and

Figure 4 illustrates the evolution of the maximum and minimum stress values achieved during a strain control test with a peak strain of 1.4%. It can be seen that there are significant differences between the RD and TD specimens, with considerably more stress relaxation in the early part of the test in the RD specimen. This is related to the availability of prismatic planes (which have a low critical resolved shear stress in titanium at room temperature) for slip in the early part of the test (see Figure 2). This leads to the TD specimen operating at a higher maximum stress. In combination with this, the TD specimen also has a higher modulus, and consequently a lower minimum stress during the test. The TD specimen therefore operates over a larger stress range for a given strain range and shows a reduced fatigue life. When considered only on a stabilised stress range basis, previous work has shown that no difference in fatigue life occurs between the RD and TD specimens in this

Fig. 2. Microstructure and texture of Titanium 6-4 UD plate.

**3.1.1 Plain specimen behaviour** 

the effect of the material texture.

material.

Clearly the anisotropy of the HCP lattice within a single grain demonstrates only the most extreme case. In terms of macroscopic properties, the averaging of these orientation effects over many thousands of grains will reduce the impact of crystallographic texture. However, as previously demonstrated (Lutjering, 1998) various forms of material processing tend to produce alignment of grain orientations, and macroscopic textures may occur which produce anisotropy in the mechanical properties of the alloy. As stated, these variations will affect not only the elastic deformation of the material, but also the plastic deformation through slip behaviour.

Fatigue accounts for approximately 80% of all in service failures and is an essential design criterion in any engineering application. LCF can be defined as failures in the region 103-105 cycles, and typically involves bulk plastic deformation within the material, whereas HCF failures (>105 cycles) usually exhibit only localized plastic deformation. Both however are affected by crystallographic texture, and it is important that designers understand the potential variations of this mechanical anisotropy. Whilst research in the 1970s sought to minimise the variations to produce a more isotropic product, more recently attention has turned to texture as a means of extracting improved properties from the alloy.

Many studies (Bowen, 1977; Peters et. al., 1984; Evans et. al., 2005a) have focussed on the type of variations in monotonic, axial fatigue and torsion-fatigue that can be expected in strongly textured alloys. However, for many design applications these conditions are overly simplistic. The presence of stress raising geometric features results in a non-linear stress state in components, which requires evaluation. In the study undertaken at Swansea, consideration was given to these features, which are represented in a laboratory environment by notched specimens. The geometrical discontinuities give rise to biaxial or even triaxial stress states which have not yet been thoroughly evaluated in textured materials. Furthermore, previous work has tended to focus on fully reversed (R=-1) or zeromaximum (R=0) loading conditions. The current work sought to further understanding by considering R ratios from R=-1 to R=0.8.

The titanium alloy Ti6-4 has long been regarded as the workhorse of the titanium industry and accounts for approximately 60% of total titanium production (Boyer, 1996). It is an alloy which is used in all forms, including forgings, bar, castings, foil, plate and sheet. In the current study however, only the plate form is considered. Uni-directional (UD) rolling of the alloy within the temperature field usually leads to the formation of a transverse or basal/transverse texture. The current material is no exception with a bimodal microstructure with an average primary alpha grain size of 15m (some deviation from equiaxality was shown in the form of slight elongation towards the rolling direction), and a basal/transverse texture with a x3 random intensity as shown in Figure 2, recorded over an area of 0.5 x 0.5 mm. The crystallographic texture of the material is completely described by the three orthogonal pole figures shown in Figure 2.

A number of previous studies have also focussed on the behaviour of the material under load controlled conditions. However, a number of potential applications for textured alloys more closely simulate strain controlled deformation behaviour, mainly because of the constraint of surrounding material. It is also argued that material at the notch root in tested specimens (and therefore at stress raising features in components) undergoes an essentially strain control type of deformation (Evans, 1998). It is therefore critical that the effect of texture under this mode of deformation is accurately characterised.

Clearly the anisotropy of the HCP lattice within a single grain demonstrates only the most extreme case. In terms of macroscopic properties, the averaging of these orientation effects over many thousands of grains will reduce the impact of crystallographic texture. However, as previously demonstrated (Lutjering, 1998) various forms of material processing tend to produce alignment of grain orientations, and macroscopic textures may occur which produce anisotropy in the mechanical properties of the alloy. As stated, these variations will affect not only the elastic deformation of the material, but also the plastic deformation

Fatigue accounts for approximately 80% of all in service failures and is an essential design criterion in any engineering application. LCF can be defined as failures in the region 103-105 cycles, and typically involves bulk plastic deformation within the material, whereas HCF failures (>105 cycles) usually exhibit only localized plastic deformation. Both however are affected by crystallographic texture, and it is important that designers understand the potential variations of this mechanical anisotropy. Whilst research in the 1970s sought to minimise the variations to produce a more isotropic product, more recently attention has

Many studies (Bowen, 1977; Peters et. al., 1984; Evans et. al., 2005a) have focussed on the type of variations in monotonic, axial fatigue and torsion-fatigue that can be expected in strongly textured alloys. However, for many design applications these conditions are overly simplistic. The presence of stress raising geometric features results in a non-linear stress state in components, which requires evaluation. In the study undertaken at Swansea, consideration was given to these features, which are represented in a laboratory environment by notched specimens. The geometrical discontinuities give rise to biaxial or even triaxial stress states which have not yet been thoroughly evaluated in textured materials. Furthermore, previous work has tended to focus on fully reversed (R=-1) or zeromaximum (R=0) loading conditions. The current work sought to further understanding by

The titanium alloy Ti6-4 has long been regarded as the workhorse of the titanium industry and accounts for approximately 60% of total titanium production (Boyer, 1996). It is an alloy which is used in all forms, including forgings, bar, castings, foil, plate and sheet. In the current study however, only the plate form is considered. Uni-directional (UD) rolling of the alloy within the temperature field usually leads to the formation of a transverse or basal/transverse texture. The current material is no exception with a bimodal microstructure with an average primary alpha grain size of 15m (some deviation from equiaxality was shown in the form of slight elongation towards the rolling direction), and a basal/transverse texture with a x3 random intensity as shown in Figure 2, recorded over an area of 0.5 x 0.5 mm. The crystallographic texture of the material is completely described by

A number of previous studies have also focussed on the behaviour of the material under load controlled conditions. However, a number of potential applications for textured alloys more closely simulate strain controlled deformation behaviour, mainly because of the constraint of surrounding material. It is also argued that material at the notch root in tested specimens (and therefore at stress raising features in components) undergoes an essentially strain control type of deformation (Evans, 1998). It is therefore critical that the effect of

turned to texture as a means of extracting improved properties from the alloy.

through slip behaviour.

considering R ratios from R=-1 to R=0.8.

the three orthogonal pole figures shown in Figure 2.

texture under this mode of deformation is accurately characterised.

Fig. 2. Microstructure and texture of Titanium 6-4 UD plate.

#### **3.1.1 Plain specimen behaviour**

Figure 3 illustrates the type of effect that may be seen in textured alloys. Under strain control testing at R=0 (20°C, 1-1-1-1 second trapezoid cycle) it can be seen that specimens taken parallel to the rolling direction of the material show longer fatigue lives than counterparts taken from the transverse direction of the plate (TD), i.e. at 90° to the rolling direction. Initially this result may be considered counterintuitive; the increased density of basal planes perpendicular to the TD direction results in a higher modulus in the TD specimens, and also higher yield stress and UTS values. However, in order to understand this effect, it is necessary to consider the behaviour of the material under strain control, and the effect of the material texture.

Figure 4 illustrates the evolution of the maximum and minimum stress values achieved during a strain control test with a peak strain of 1.4%. It can be seen that there are significant differences between the RD and TD specimens, with considerably more stress relaxation in the early part of the test in the RD specimen. This is related to the availability of prismatic planes (which have a low critical resolved shear stress in titanium at room temperature) for slip in the early part of the test (see Figure 2). This leads to the TD specimen operating at a higher maximum stress. In combination with this, the TD specimen also has a higher modulus, and consequently a lower minimum stress during the test. The TD specimen therefore operates over a larger stress range for a given strain range and shows a reduced fatigue life. When considered only on a stabilised stress range basis, previous work has shown that no difference in fatigue life occurs between the RD and TD specimens in this material.

Titanium in the Gas Turbine Engine 321

(Walker, 1970) is an appropriate method for these types of predictions. The Walker strain relationship is an empirical method for correlating R values and involves correlating strain control data of different R ratios, allowing for the derivation of a 'master curve'. As stated earlier, the material at a notch root is assumed to experience strain control type conditions, due to restraint from material surrounding the critically stressed volume of material. Through application of Neuber's rule (Neuber, 1968), that the product of stress and strain is a constant, conditions at the notch root can be approximated allowing for the calculation of the individual Walker strain value for that specimen. Subsequently a predicted life can be inferred from the 'master curve' based on the strain control data. This approach has been found to be accurate for similar titanium alloys to Ti6-4 (Whittaker et. al., 2007), but has not

During the course of the work, two notched specimen geometries were tested, both with cylindrical notches; the first was a V shaped cylindrical notch (VCN) which has a stress concentration factor, Kt, of 2.8, the second a round cylindrical notch (RCN) with a Kt of 1.4. Initially apparent from Figure 5 is the fact that no orientation effect appears to exist in the RCN specimen, with both RD and TD specimens showing similar fatigue lives to the plain specimen data. However, this is not the case with the VCN specimen, as shown in Figure 6,

Fig. 5. Comparison of notched (RCN) and plain specimen response showing no orientation

To interpret these results, it should be noted that these notched specimen tests are performed under load control; it is the geometry of the notch which imposes strain control type conditions on the material at the root of the notch. Figure 5, showing the results for the RCN specimen, is illustrative in a number of ways. Along with the fact that no orientation

with the RD specimens showing longer fatigue lives than the TD specimens.

previously been tested on a textured alloy.

effect.

Fig. 3. The effect of orientation on the strain-life of Ti6-4.

Fig. 4. The effect of orientation of the stress response of Ti6-4 under strain control loading (max=1.4%, R=0).

#### **3.1.2 Notched specimen behaviour**

In considering notched specimen behaviour, it is important to acknowledge the requirement for a predictive methodology, to enable designers to extrapolate to conditions for which reliable test data does not exist. Previous work has shown that the Walker strain approach

Fig. 4. The effect of orientation of the stress response of Ti6-4 under strain control loading

In considering notched specimen behaviour, it is important to acknowledge the requirement for a predictive methodology, to enable designers to extrapolate to conditions for which reliable test data does not exist. Previous work has shown that the Walker strain approach

Fig. 3. The effect of orientation on the strain-life of Ti6-4.

(max=1.4%, R=0).

**3.1.2 Notched specimen behaviour** 

(Walker, 1970) is an appropriate method for these types of predictions. The Walker strain relationship is an empirical method for correlating R values and involves correlating strain control data of different R ratios, allowing for the derivation of a 'master curve'. As stated earlier, the material at a notch root is assumed to experience strain control type conditions, due to restraint from material surrounding the critically stressed volume of material. Through application of Neuber's rule (Neuber, 1968), that the product of stress and strain is a constant, conditions at the notch root can be approximated allowing for the calculation of the individual Walker strain value for that specimen. Subsequently a predicted life can be inferred from the 'master curve' based on the strain control data. This approach has been found to be accurate for similar titanium alloys to Ti6-4 (Whittaker et. al., 2007), but has not previously been tested on a textured alloy.

During the course of the work, two notched specimen geometries were tested, both with cylindrical notches; the first was a V shaped cylindrical notch (VCN) which has a stress concentration factor, Kt, of 2.8, the second a round cylindrical notch (RCN) with a Kt of 1.4. Initially apparent from Figure 5 is the fact that no orientation effect appears to exist in the RCN specimen, with both RD and TD specimens showing similar fatigue lives to the plain specimen data. However, this is not the case with the VCN specimen, as shown in Figure 6, with the RD specimens showing longer fatigue lives than the TD specimens.

Fig. 5. Comparison of notched (RCN) and plain specimen response showing no orientation effect.

To interpret these results, it should be noted that these notched specimen tests are performed under load control; it is the geometry of the notch which imposes strain control type conditions on the material at the root of the notch. Figure 5, showing the results for the RCN specimen, is illustrative in a number of ways. Along with the fact that no orientation

Titanium in the Gas Turbine Engine 323

In considering the ability of the Walker strain method to accurately predict fatigue lives, only RD specimens are currently considered, although similar calculations can be made for TD specimens (Evans & Whittaker, 2006). Although the Walker strain method is a relatively simplistic method, and does not compensate for notch type, it is a useful approach that has previously been shown to give excellent results in titanium alloys (Whittaker et. al., 2007). Figure 7 shows the type of predictions which can be made using this approach, over a wide range of R ratios. In order to consider a total life prediction methodology it should be recognised that this type of approach predicts only fatigue crack initiation in notched specimens. In strain control specimens, when a crack initiates, it will propagate quickly to failure. This is not the case in a notched specimen where the crack will grow more slowly through material away from the notch root. Previous crack monitoring work has shown that assuming a propagation phase of 50% of the total life allows for reasonable predictions

Fig. 7. Predictions of notched fatigue lives in RCN and VCN notches by the Walker strain

Based on these assumptions it is clear that excellent predictions are made for R ratios of -1, 0 and 0.5. However, significant over predictions are made at an R=0.8, particularly for the RCN specimens. The reason for this lies in the introduction of additional failure mechanisms. Strain accumulation at low temperatures has been widely reported in near and titanium alloys and is loosely termed 'cold dwell'. Particularly at high mean stresses, these failures are characterised by the formation of quasi-cleavage facets which form due to stress redistribution from so called 'soft' (suitably orientated for slip) grains onto 'hard' grains (unsuitably orientated for slip), as shown by the Evans-Bache model in Figure 8(a) (Bache & Evans, 1996). Clear evidence of these facets was found in both RCN

(Whittaker et. al, 2010a).

method.

effect exists, it can also be seen that RD and TD specimens show similar fatigue lives. Furthermore, the notched specimen behaviour correlates well with the plain specimen response. The VCN specimens, however, do not follow either of these trends, Figure 6. Specimens in the RD orientation show longer fatigue lives than either plain or RCN specimens. This is consistent with previous experience since a lower volume of material is critically stressed in the VCN specimen. Since fatigue is essentially probabilistic in nature and relies on 'weak links' present in the material to initiate the fatigue process, a lower material volume infers a lower probability of a 'weak link' being present, and hence a longer fatigue life is statistically more likely.

The fact that the RCN specimen shows no orientation effect and correlated well with the plain specimen data when plotted on a stabilised stress basis indicates that a lack of constraint is occurring at the notch root. In this case a large volume of material is critically, or near critically stressed, similar to the plain specimens. Since the notch testing is performed under load control, the lack of constraint at the notch results in a shallow stress gradient, and hence the material at the notch root experiences conditions closer to load control than strain control. As a result these specimens behave like the plain specimens when considered on a stabilised stress basis, with no orientation effect. In the VCN specimens the stress gradient is far steeper, constraining material at the notch root, which then behave like the plain specimens, when considered on a strain range basis, and RD specimens show longer lives than TD specimens for the same reasons described in plain specimens (i.e. changes in relaxation behaviour and differences in modulus), as explained in the previous section.

Fig. 6. Comparison of notched specimen fatigue lives showing an orientation effect in the VCN notch, whereas no such effect exists in the RCN notch.

effect exists, it can also be seen that RD and TD specimens show similar fatigue lives. Furthermore, the notched specimen behaviour correlates well with the plain specimen response. The VCN specimens, however, do not follow either of these trends, Figure 6. Specimens in the RD orientation show longer fatigue lives than either plain or RCN specimens. This is consistent with previous experience since a lower volume of material is critically stressed in the VCN specimen. Since fatigue is essentially probabilistic in nature and relies on 'weak links' present in the material to initiate the fatigue process, a lower material volume infers a lower probability of a 'weak link' being present, and hence a longer

The fact that the RCN specimen shows no orientation effect and correlated well with the plain specimen data when plotted on a stabilised stress basis indicates that a lack of constraint is occurring at the notch root. In this case a large volume of material is critically, or near critically stressed, similar to the plain specimens. Since the notch testing is performed under load control, the lack of constraint at the notch results in a shallow stress gradient, and hence the material at the notch root experiences conditions closer to load control than strain control. As a result these specimens behave like the plain specimens when considered on a stabilised stress basis, with no orientation effect. In the VCN specimens the stress gradient is far steeper, constraining material at the notch root, which then behave like the plain specimens, when considered on a strain range basis, and RD specimens show longer lives than TD specimens for the same reasons described in plain specimens (i.e. changes in relaxation behaviour and differences in modulus), as explained in

Fig. 6. Comparison of notched specimen fatigue lives showing an orientation effect in the

VCN notch, whereas no such effect exists in the RCN notch.

fatigue life is statistically more likely.

the previous section.

In considering the ability of the Walker strain method to accurately predict fatigue lives, only RD specimens are currently considered, although similar calculations can be made for TD specimens (Evans & Whittaker, 2006). Although the Walker strain method is a relatively simplistic method, and does not compensate for notch type, it is a useful approach that has previously been shown to give excellent results in titanium alloys (Whittaker et. al., 2007). Figure 7 shows the type of predictions which can be made using this approach, over a wide range of R ratios. In order to consider a total life prediction methodology it should be recognised that this type of approach predicts only fatigue crack initiation in notched specimens. In strain control specimens, when a crack initiates, it will propagate quickly to failure. This is not the case in a notched specimen where the crack will grow more slowly through material away from the notch root. Previous crack monitoring work has shown that assuming a propagation phase of 50% of the total life allows for reasonable predictions (Whittaker et. al, 2010a).

Fig. 7. Predictions of notched fatigue lives in RCN and VCN notches by the Walker strain method.

Based on these assumptions it is clear that excellent predictions are made for R ratios of -1, 0 and 0.5. However, significant over predictions are made at an R=0.8, particularly for the RCN specimens. The reason for this lies in the introduction of additional failure mechanisms. Strain accumulation at low temperatures has been widely reported in near and titanium alloys and is loosely termed 'cold dwell'. Particularly at high mean stresses, these failures are characterised by the formation of quasi-cleavage facets which form due to stress redistribution from so called 'soft' (suitably orientated for slip) grains onto 'hard' grains (unsuitably orientated for slip), as shown by the Evans-Bache model in Figure 8(a) (Bache & Evans, 1996). Clear evidence of these facets was found in both RCN

Titanium in the Gas Turbine Engine 325

Fig. 9. Micrograph of Ti6246, showing a fine Widmanstatten type microstructure.

Fig. 10. Fatigue, creep and environmental effects in crack growth in Ti6246 (Evans et. al.,

load (1-1-120-1 waveform) as indicated. This increase in crack propagation rate is due to the effect of creep, with evidence seen of creep voids ahead of the crack tip. However it is also clear that at this temperature, creep and fatigue are not the only damage mechanisms in

2005b).

and VCN R=0.8 specimens, although an increased density was found in the RCN specimens. The result of this is the reduction in fatigue lives (when compared with the Walker predictions) seen in Figure 7. The effect is more pronounced in the RCN specimens because of the larger amount of material being critically or near-critically stressed.

Fig. 8. The Evans-Bache model for facet generation in titanium alloys, with an example facet from an RCN, R=0.8 notched specimen.

Whilst it is clear that it is possible to accurately life notched specimens in a textured alloy, it is also evident that there are limitations. In the current work predictions have been made based on strain control data from the same orientation. Without this it is impossible to make accurate predictions. It is also apparent that for Ti6-4 there is a limited range of R ratios over which predictions can be made, with additional failure mechanisms playing a role.

#### **3.2 High temperature lifing (Ti6246)**

As temperatures rise in the gas turbine engine designers turn to titanium alloys with a higher temperature capability than Ti6-4, for which operation is limited to less than approximately 350⁰C. Ti6246 (Ti-6Al-2Sn-4Zr-6Mo) is such an alloy with good low cycle fatigue properties and improved creep resistance over Ti6-4, Figure 9. It is immediately apparent that the microstructure of Ti6246 differs significantly to Ti6-4, showing a fine Widmanstatten microstructure that would be typical of a material processed above the beta transus. The fine nature of the microstructure infers the high strength of the material and also offers good resistance to crack propagation.

Widely used as a compressor disc alloy, Ti6246 has traditionally been employed at temperatures where creep effects would not be considered significant. However, it is not necessary for the alloy to be limited in this way provided appropriate lifing techniques are employed. The following work describes the construction of a total life prediction capability for fatigue at high temperatures in the alloy. Again, the focus of the work is on notched specimens, due to the importance of the stress raising features within the gas turbine engine. Figure 10 demonstrates the importance of considering additional failure mechanisms to fatigue by considering crack propagation rates at 550⁰C in Ti6246. The vacuum 1Hz sinewave data (square symbols) represent solely the influence of fatigue on the crack propagation rate whereas the circular symbols indicate that as a dwell period is added to the waveform, by employing a trapezoidal 1-1-1-1 waveform, a significant increase is seen in the crack propagation rate. This is further increased by adding a 2 minute dwell period at peak

and VCN R=0.8 specimens, although an increased density was found in the RCN specimens. The result of this is the reduction in fatigue lives (when compared with the Walker predictions) seen in Figure 7. The effect is more pronounced in the RCN specimens because

Fig. 8. The Evans-Bache model for facet generation in titanium alloys, with an example facet

Whilst it is clear that it is possible to accurately life notched specimens in a textured alloy, it is also evident that there are limitations. In the current work predictions have been made based on strain control data from the same orientation. Without this it is impossible to make accurate predictions. It is also apparent that for Ti6-4 there is a limited range of R ratios over

As temperatures rise in the gas turbine engine designers turn to titanium alloys with a higher temperature capability than Ti6-4, for which operation is limited to less than approximately 350⁰C. Ti6246 (Ti-6Al-2Sn-4Zr-6Mo) is such an alloy with good low cycle fatigue properties and improved creep resistance over Ti6-4, Figure 9. It is immediately apparent that the microstructure of Ti6246 differs significantly to Ti6-4, showing a fine Widmanstatten microstructure that would be typical of a material processed above the beta transus. The fine nature of the microstructure infers the high strength of the material and

Widely used as a compressor disc alloy, Ti6246 has traditionally been employed at temperatures where creep effects would not be considered significant. However, it is not necessary for the alloy to be limited in this way provided appropriate lifing techniques are employed. The following work describes the construction of a total life prediction capability for fatigue at high temperatures in the alloy. Again, the focus of the work is on notched specimens, due to the importance of the stress raising features within the gas turbine engine. Figure 10 demonstrates the importance of considering additional failure mechanisms to fatigue by considering crack propagation rates at 550⁰C in Ti6246. The vacuum 1Hz sinewave data (square symbols) represent solely the influence of fatigue on the crack propagation rate whereas the circular symbols indicate that as a dwell period is added to the waveform, by employing a trapezoidal 1-1-1-1 waveform, a significant increase is seen in the crack propagation rate. This is further increased by adding a 2 minute dwell period at peak

which predictions can be made, with additional failure mechanisms playing a role.

of the larger amount of material being critically or near-critically stressed.

from an RCN, R=0.8 notched specimen.

**3.2 High temperature lifing (Ti6246)** 

also offers good resistance to crack propagation.

Fig. 9. Micrograph of Ti6246, showing a fine Widmanstatten type microstructure.

Fig. 10. Fatigue, creep and environmental effects in crack growth in Ti6246 (Evans et. al., 2005b).

load (1-1-120-1 waveform) as indicated. This increase in crack propagation rate is due to the effect of creep, with evidence seen of creep voids ahead of the crack tip. However it is also clear that at this temperature, creep and fatigue are not the only damage mechanisms in

Titanium in the Gas Turbine Engine 327

The predictions made for R=-1 notch tests have improved accuracy over the R=0 tests simply for the reason that it is easier to predict the stress/strain state at the notch root for these tests. The highest load which was employed in fatigue testing of the R=-1 tests resulted in a peak elastic stress of 800MPa, which would be below yield for Ti6246 at room temperature, at a typical strain rate of 0.5%/sec. As such the stress/strain conditions at the notch root are simply 800MPa and 0.0067 (from strain = stress/modulus). However, in the R=0 tests, significant plasticity is induced at the notch root. Whilst in Ti6-4 this plasticity could be accurately approximated by Neuber's rule, clearly more accurate description is

In order to achieve greater accuracy a model was developed in the modelling suite ABAQUS based upon open hysteresis loops generated under fully reversed strain control loading of Ti6246, over a range of temperatures. The loops were generated under laboratory air conditions so that fatigue/environment and subsequently fatigue/creep/environment interactions could be studied. The model was based around the Mroz multilayer kinematic hardening model (Mroz, 1969) which compared well with experimental observations that stress redistribution within the material allowed for the stabilization of the peak/minimum stress during the initial cycles of a strain control test. A typical stress-strain loop generated by the model is shown in Figure 12. It can be seen that the loop generated in ABAQUS accurately describes the test data generated for a strain control test with a peak strain of

Modelling of the double-edged notch specimen was achieved through the construction of a three dimensional 1/8 symmetrical FE model using 20-noded isoparametric rectangular elements (C3D20) with 18833 nodes and 4032 elements, with element size reduced near to the notch to improve accuracy. Calculations of the fatigue life were then based on the

stabilised conditions of stress and strain at the node adjacent to the notch root.

Fig. 12. ABAQUS modelling of a stress-strain loop at 20⁰C in Ti6246 (Whittaker et. al.,

required in the current case.

1.5%.

2010a).

**3.2.2 Development of FEA model in ABAQUS** 

operation. For tests conducted in air, rather than under high vacuum (10-6 mbar) conditions, a significant further increase in propagation rate is seen when the same 1-1-120-1 second trapezoid waveform is applied. This effect is environmental damage and as indicated by the graph, also requires consideration, since the increases in crack growth can be similar to, or even surpass those due to creep.

Whilst these results give an indication of the roles of fatigue, creep and environmental damage, it is clear that in order to build a total life prediction capability, their effects on fatigue crack initiation must be considered.

#### **3.2.1 Fatigue modelling**

As described previously the Walker strain method (Walker, 1970) has been shown to be a useful approach to the prediction of notched specimen behaviour, particularly in terms of predictions over a wide range of R ratios. However, the previous analysis was performed only at room temperature and it is necessary to investigate whether the Walker strain approach still offers accurate results at higher temperatures. In this work the notch considered is a double edged notch (DEN) with a Kt = 1.9.

Figure 11 illustrates predictions made using the Walker strain approach at 20⁰C and 450⁰C, with notch root conditions again approximated by use of Neuber's rule (Neuber, 1968). As described previously, these predictions do not account for the crack propagation phase of a notch test and assuming a propagation phase of approximately 50% of the total life has previously been shown to be a reasonable assumption (Whittaker, 2010a). Whilst predictions under R=-1 loading conditions are excellent, it can be seen that predictions for R=0 tests at 20⁰C and 450⁰C tend to be non-conservative when the propagation phase is added. This is obviously undesirable for designers of critical parts.

Fig. 11. Predictions of notched specimen behaviour at 20⁰C and 450⁰C using the Walker strain method.

operation. For tests conducted in air, rather than under high vacuum (10-6 mbar) conditions, a significant further increase in propagation rate is seen when the same 1-1-120-1 second trapezoid waveform is applied. This effect is environmental damage and as indicated by the graph, also requires consideration, since the increases in crack growth can be similar to, or

Whilst these results give an indication of the roles of fatigue, creep and environmental damage, it is clear that in order to build a total life prediction capability, their effects on

As described previously the Walker strain method (Walker, 1970) has been shown to be a useful approach to the prediction of notched specimen behaviour, particularly in terms of predictions over a wide range of R ratios. However, the previous analysis was performed only at room temperature and it is necessary to investigate whether the Walker strain approach still offers accurate results at higher temperatures. In this work the notch

Figure 11 illustrates predictions made using the Walker strain approach at 20⁰C and 450⁰C, with notch root conditions again approximated by use of Neuber's rule (Neuber, 1968). As described previously, these predictions do not account for the crack propagation phase of a notch test and assuming a propagation phase of approximately 50% of the total life has previously been shown to be a reasonable assumption (Whittaker, 2010a). Whilst predictions under R=-1 loading conditions are excellent, it can be seen that predictions for R=0 tests at 20⁰C and 450⁰C tend to be non-conservative when the propagation phase is added. This is

Fig. 11. Predictions of notched specimen behaviour at 20⁰C and 450⁰C using the Walker

even surpass those due to creep.

**3.2.1 Fatigue modelling** 

strain method.

fatigue crack initiation must be considered.

considered is a double edged notch (DEN) with a Kt = 1.9.

obviously undesirable for designers of critical parts.

The predictions made for R=-1 notch tests have improved accuracy over the R=0 tests simply for the reason that it is easier to predict the stress/strain state at the notch root for these tests. The highest load which was employed in fatigue testing of the R=-1 tests resulted in a peak elastic stress of 800MPa, which would be below yield for Ti6246 at room temperature, at a typical strain rate of 0.5%/sec. As such the stress/strain conditions at the notch root are simply 800MPa and 0.0067 (from strain = stress/modulus). However, in the R=0 tests, significant plasticity is induced at the notch root. Whilst in Ti6-4 this plasticity could be accurately approximated by Neuber's rule, clearly more accurate description is required in the current case.

#### **3.2.2 Development of FEA model in ABAQUS**

In order to achieve greater accuracy a model was developed in the modelling suite ABAQUS based upon open hysteresis loops generated under fully reversed strain control loading of Ti6246, over a range of temperatures. The loops were generated under laboratory air conditions so that fatigue/environment and subsequently fatigue/creep/environment interactions could be studied. The model was based around the Mroz multilayer kinematic hardening model (Mroz, 1969) which compared well with experimental observations that stress redistribution within the material allowed for the stabilization of the peak/minimum stress during the initial cycles of a strain control test. A typical stress-strain loop generated by the model is shown in Figure 12. It can be seen that the loop generated in ABAQUS accurately describes the test data generated for a strain control test with a peak strain of 1.5%.

Modelling of the double-edged notch specimen was achieved through the construction of a three dimensional 1/8 symmetrical FE model using 20-noded isoparametric rectangular elements (C3D20) with 18833 nodes and 4032 elements, with element size reduced near to the notch to improve accuracy. Calculations of the fatigue life were then based on the stabilised conditions of stress and strain at the node adjacent to the notch root.

Fig. 12. ABAQUS modelling of a stress-strain loop at 20⁰C in Ti6246 (Whittaker et. al., 2010a).

Titanium in the Gas Turbine Engine 329

Fig. 14. Influence of environment on the fatigue lives of notched specimens in Ti6246

Clearly the interactions of fatigue, creep and environment within the material are complex and offer a challenge to designers who wish to make accurate life predictions. However, some limited success has been achieved by the development of a fatigue-creep-environment model for crack growth. To construct the model it was necessary to combine fatigue crack predictions based on laboratory air conditions with damage due to creep effects, in the form

> *f da da da dt dN dN dN*

where the first term on the right hand side represents fatigue damage and the second term

To calculate the creep damage, a suitable creep model was required. Previous experience had shown that the theta projection method offered acceptable results of creep behaviour in Ti6246, so an ABAQUS subroutine for the relationship was compiled which included creep rupture based on a Kachanov type failure process. In order to predict fatigue crack growth at high temperatures fatigue and creep damage were then calculated separately and combined at each time increment to give the total growth rate. Figure 15 indicates the results of this approach for growth rates at 500°C, R=0.1. It is clear that a prediction based purely on fatigue significantly underestimates the growth rate, but when the combined fatigue-creepenvironment prediction is made, predictions are accurate. The effect of further creep damage is represented by the growth rates under a waveform with a two minute dwell at

Whilst it is acknowledged that there is still much work to be completed in developing a total life prediction methodology for fatigue performance at high temperatures, the results of the

peak stress, although currently predictions have not been made for this data.

**3.2.4 Combining fatigue, creep and environmental damage** 

(Whittaker et. al., 2010a).

represents creep damage.

#### **3.2.3 Creep and environmental damage**

Figure 13 shows the predictions made by the model under 20⁰C R=-1 loading conditions, and also 500⁰C R=0 loading conditions. It can be seen that the low temperature predictions of initiation life are again extremely accurate. At 500⁰C the predictions are slightly conservative, but clearly more acceptable than those previously demonstrated without the use of FEA. Previous work (Whittaker et. al., 2010a) has in fact shown that in this material, using DEN specimens, fatigue lives at 500⁰C are actually longer than at 450⁰C. This is due to the effect of creep within the vicinity of the notch root. At 450⁰C creep has a limited effect, whereas at 500⁰C it becomes more prevalent, and acts to decrease the stresses around the notch root, creating a shallower stress gradient and hence an improved fatigue life. Further increases in temperature to 550⁰C however, lead to a reduction in fatigue life as creep and environmental effects become more damaging.

Fig. 13. Predictions of notched fatigue life made by ABAQUS model at 20⁰C and 500⁰C.

Further evidence of the significance of environment is demonstrated in Figure 14. Previous authors have described the development of a marked transition in the fatigue life curve of Ti6246 when tested under strain control (Mailly, 1999). Similar effects have been observed in the current work, where for lives greater than approximately 104 cycles the fatigue lives of the material may be highly variable as the curve becomes very flat. At this point the material is protected by an oxide layer which forms during the test, preventing further oxidation. However, as material strain increases as the applied stress is raised, the oxide layer cracks and allows further ingress of oxygen, causing damage to the material and resulting in a more typical fatigue curve. The effect is not observed at 20⁰C, but interestingly has been seen in strain control tests at temperature as low as 80⁰C.

Figure 13 shows the predictions made by the model under 20⁰C R=-1 loading conditions, and also 500⁰C R=0 loading conditions. It can be seen that the low temperature predictions of initiation life are again extremely accurate. At 500⁰C the predictions are slightly conservative, but clearly more acceptable than those previously demonstrated without the use of FEA. Previous work (Whittaker et. al., 2010a) has in fact shown that in this material, using DEN specimens, fatigue lives at 500⁰C are actually longer than at 450⁰C. This is due to the effect of creep within the vicinity of the notch root. At 450⁰C creep has a limited effect, whereas at 500⁰C it becomes more prevalent, and acts to decrease the stresses around the notch root, creating a shallower stress gradient and hence an improved fatigue life. Further increases in temperature to 550⁰C however, lead to a reduction in fatigue life as creep and

Fig. 13. Predictions of notched fatigue life made by ABAQUS model at 20⁰C and 500⁰C.

seen in strain control tests at temperature as low as 80⁰C.

Further evidence of the significance of environment is demonstrated in Figure 14. Previous authors have described the development of a marked transition in the fatigue life curve of Ti6246 when tested under strain control (Mailly, 1999). Similar effects have been observed in the current work, where for lives greater than approximately 104 cycles the fatigue lives of the material may be highly variable as the curve becomes very flat. At this point the material is protected by an oxide layer which forms during the test, preventing further oxidation. However, as material strain increases as the applied stress is raised, the oxide layer cracks and allows further ingress of oxygen, causing damage to the material and resulting in a more typical fatigue curve. The effect is not observed at 20⁰C, but interestingly has been

**3.2.3 Creep and environmental damage** 

environmental effects become more damaging.

Fig. 14. Influence of environment on the fatigue lives of notched specimens in Ti6246 (Whittaker et. al., 2010a).

#### **3.2.4 Combining fatigue, creep and environmental damage**

Clearly the interactions of fatigue, creep and environment within the material are complex and offer a challenge to designers who wish to make accurate life predictions. However, some limited success has been achieved by the development of a fatigue-creep-environment model for crack growth. To construct the model it was necessary to combine fatigue crack predictions based on laboratory air conditions with damage due to creep effects, in the form

$$
\left(\frac{da}{dN}\right) = \left(\frac{da}{dN}\right)\_f + \iint \frac{da}{dN} \bigg| dt
$$

where the first term on the right hand side represents fatigue damage and the second term represents creep damage.

To calculate the creep damage, a suitable creep model was required. Previous experience had shown that the theta projection method offered acceptable results of creep behaviour in Ti6246, so an ABAQUS subroutine for the relationship was compiled which included creep rupture based on a Kachanov type failure process. In order to predict fatigue crack growth at high temperatures fatigue and creep damage were then calculated separately and combined at each time increment to give the total growth rate. Figure 15 indicates the results of this approach for growth rates at 500°C, R=0.1. It is clear that a prediction based purely on fatigue significantly underestimates the growth rate, but when the combined fatigue-creepenvironment prediction is made, predictions are accurate. The effect of further creep damage is represented by the growth rates under a waveform with a two minute dwell at peak stress, although currently predictions have not been made for this data.

Whilst it is acknowledged that there is still much work to be completed in developing a total life prediction methodology for fatigue performance at high temperatures, the results of the

Titanium in the Gas Turbine Engine 331

mechanical testing aimed at detailing these variations in a range of prestrained Ti834

Fig. 16. Micrograph of Ti834, indicating a bimodal microstructure with primary alpha grains

Figure 16 shows the microstructure of Ti834 tested, with a bimodal microstructure clearly evident encompassing primary grains of 20-200m in diameter. Total tensile prestrains of 2% and 8% (resulting in 1.25% and 7.25% plastic prestrains respectively) were applied to individual batches of specimens along with a compressive prestrain of 2%, denoted as -2% (plastic prestrain of -1.25%). These four different conditions, -2%, 0% (as received), 2% and 8% could then be compared under different loading conditions such as fully reversed strain control fatigue (20°C), stress relaxation at 1% strain (20°C) and creep (20°C and 600°C).

Fig. 17. Effect of prestrain on the creep rate of Ti834 at 20°C (Whittaker et. al. 2010b)

specimens.

ranging in size from 20-200m.

work are encouraging. It has been demonstrated that interactions between fatigue, creep and environment are complex and produce many non-linear effects which are difficult to model. However, some success has been achieved in the production of a fatigue crack growth model at high temperatures and the belief is that similar models could be produced to describe crack initiation lives based on suitable deformation data for the alloy. Clearly, to accurately build the model, all three damage mechanisms should be considered independently before coupling in a model which considers their effects. However, in order to achieve this, a greater proportion of vacuum data will be required, particularly under strain control conditions, which will be experimentally challenging.

Fig. 15. Predictions of crack growth behaviour of Ti6246 at 500°C, R=0.1 (Whittaker et. al., 2010a).

#### **3.3 Application of prestrain (Ti834)**

Ti834 is a near titanium alloy which was developed with a carefully controlled microstructure to enable exceptional mechanical properties at temperatures up to approximately 630°C. Combined with the high strength to weight ratio of the alloy, this excellent elevated temperature behaviour makes the alloy a popular choice for applications such as compressor discs and blades.

Clearly it is critical that the alloy is utilised under well understood conditions where any effect of the processing history can be accounted for. This may be a complex issue with different forging/machining/peening parameters influencing the surface condition of the material. During processing of components it is highly likely that surface roughness variations may occur, along with the possibility of further 'damage' to the material. Combined with the effect of residual stresses brought about by the peening process (commonly used to extend fatigue life), it is clear that significant variation may occur in the material mechanical properties. It is therefore necessary to understand these effects through a detailed investigation. In the current work, this was undertaken through a programme of

work are encouraging. It has been demonstrated that interactions between fatigue, creep and environment are complex and produce many non-linear effects which are difficult to model. However, some success has been achieved in the production of a fatigue crack growth model at high temperatures and the belief is that similar models could be produced to describe crack initiation lives based on suitable deformation data for the alloy. Clearly, to accurately build the model, all three damage mechanisms should be considered independently before coupling in a model which considers their effects. However, in order to achieve this, a greater proportion of vacuum data will be required, particularly under

Fig. 15. Predictions of crack growth behaviour of Ti6246 at 500°C, R=0.1 (Whittaker et. al.,

Ti834 is a near titanium alloy which was developed with a carefully controlled microstructure to enable exceptional mechanical properties at temperatures up to approximately 630°C. Combined with the high strength to weight ratio of the alloy, this excellent elevated temperature behaviour makes the alloy a popular choice for applications

Clearly it is critical that the alloy is utilised under well understood conditions where any effect of the processing history can be accounted for. This may be a complex issue with different forging/machining/peening parameters influencing the surface condition of the material. During processing of components it is highly likely that surface roughness variations may occur, along with the possibility of further 'damage' to the material. Combined with the effect of residual stresses brought about by the peening process (commonly used to extend fatigue life), it is clear that significant variation may occur in the material mechanical properties. It is therefore necessary to understand these effects through a detailed investigation. In the current work, this was undertaken through a programme of

2010a).

**3.3 Application of prestrain (Ti834)** 

such as compressor discs and blades.

strain control conditions, which will be experimentally challenging.

mechanical testing aimed at detailing these variations in a range of prestrained Ti834 specimens.

Fig. 16. Micrograph of Ti834, indicating a bimodal microstructure with primary alpha grains ranging in size from 20-200m.

Figure 16 shows the microstructure of Ti834 tested, with a bimodal microstructure clearly evident encompassing primary grains of 20-200m in diameter. Total tensile prestrains of 2% and 8% (resulting in 1.25% and 7.25% plastic prestrains respectively) were applied to individual batches of specimens along with a compressive prestrain of 2%, denoted as -2% (plastic prestrain of -1.25%). These four different conditions, -2%, 0% (as received), 2% and 8% could then be compared under different loading conditions such as fully reversed strain control fatigue (20°C), stress relaxation at 1% strain (20°C) and creep (20°C and 600°C).

Fig. 17. Effect of prestrain on the creep rate of Ti834 at 20°C (Whittaker et. al. 2010b)

Titanium in the Gas Turbine Engine 333

temperatures increased prestrain restricts further creep damage because of the high dislocation densities and apparent difficulty in processes such as climb and cross slip. Conversely, these processes occur more readily at 600°C and increases in prestrain lead to

Fig. 19. Effect of prestrain on the fatigue properties of Ti834 at 20°C (Whittaker & Evans,

However, stress states in the gas turbine are rarely static and as such further consideration must be given to the effect on fatigue performance of the material, Figure 19. In the current work it was found that a small period of stress relaxation (<2 seconds) occurred at the end of the prestrain process before the specimen was unloaded. Previous work (Evans, 1998) has demonstrated that near and titanium alloys tend to form facets under stress relaxation, and fractographic analysis of failed specimens showed that this was indeed the case here. These facets offer initiation sites for fatigue cracks, which along with the increased dislocation density contributes to the 8% prestrain specimens showing significantly shorter

It is clear that whilst the rate of development of new titanium alloys has slowed in recent years, there are further areas which may be explored in order to achieve further improvements in mechanical properties. The research described here has shown that there is definite potential through the harnessing of texture, improved high temperature lifing

Of these, perhaps improvements in high temperature lifing offer designers the greatest reward. Since the development of the gas turbine engine, increased efficiency has acted as the driver which has led to operation at higher and higher temperatures. Enabling components to operate at higher temperatures whilst retaining low density/low cost materials in their manufacture is obviously desirable and operation at temperatures where

techniques or improved understanding of processing effects.

an acceleration in creep damage.

2009).

fatigue lives.

**4. Discussion** 

Fully reversed strain control loading at a peak strain of 1% was shown not to result in the formation of quasi-cleavage facets and as such was used as a control mechanism to produce eventual failure in test samples. In this way specimens could be separated and fracture surfaces investigated with the confidence that any facets generated would have been generated by the previous loading conditions and not the strain control fatigue.

Creep rates at room temperature (tested at 950MPa) were shown to be significantly affected by the application of prestrain, Figure 17. It can be seen that for 2% and -2% tests the primary creep is greatly reduced, although the creep strain rate increases and the strain at failure and creep life are markedly reduced. These effects offer only a limited improvement window for the material, in which creep strain is reduced over first few hours. However, the specimen which had undergone 8% prestrain showed a dramatic reduction in creep rate, eventually being removed from test after 250 hours, in which little creep was seen.

Fig. 18. Effect of prestrain on the creep rate of Ti834 at 600°C (Whittaker et. al., 2010b).

Clearly for designers interested in reduced creep rates at room temperature, this effect is attractive. Figure 18 illustrates though that these advantages will be temperature dependent. At 600°C the 8% prestrain specimen now shows a faster creep rate, shorter lives and reduced strain at failure. The reason for both of these effects will be related to the dislocation structure following prestrain. During the prestrain process, dislocations are generated as the yield stress is exceeded, which occurs at approximately 0.75% strain. It is clear that as the specimen continues to extend towards 8% strain, the dislocations will continue to multiply and as a result a high dislocation density occurs in the material. At room temperature, under creep conditions, dislocation mobility in the structure is significantly reduced, and the creep rate remains very low. However, at 600°C the increased thermal energy means that processes such as climb and cross slip become more prevalent, increasing dislocation mobility. This results in an increased creep rate when compared with the as received (0% prestrain) material.

Based on these results, it is clear that the effects of prestrain on the creep performance of the alloy vary significantly with temperature, and as such, dislocation mobility. At low

Fully reversed strain control loading at a peak strain of 1% was shown not to result in the formation of quasi-cleavage facets and as such was used as a control mechanism to produce eventual failure in test samples. In this way specimens could be separated and fracture surfaces investigated with the confidence that any facets generated would have been

Creep rates at room temperature (tested at 950MPa) were shown to be significantly affected by the application of prestrain, Figure 17. It can be seen that for 2% and -2% tests the primary creep is greatly reduced, although the creep strain rate increases and the strain at failure and creep life are markedly reduced. These effects offer only a limited improvement window for the material, in which creep strain is reduced over first few hours. However, the specimen which had undergone 8% prestrain showed a dramatic reduction in creep rate,

generated by the previous loading conditions and not the strain control fatigue.

eventually being removed from test after 250 hours, in which little creep was seen.

Fig. 18. Effect of prestrain on the creep rate of Ti834 at 600°C (Whittaker et. al., 2010b).

prestrain) material.

Clearly for designers interested in reduced creep rates at room temperature, this effect is attractive. Figure 18 illustrates though that these advantages will be temperature dependent. At 600°C the 8% prestrain specimen now shows a faster creep rate, shorter lives and reduced strain at failure. The reason for both of these effects will be related to the dislocation structure following prestrain. During the prestrain process, dislocations are generated as the yield stress is exceeded, which occurs at approximately 0.75% strain. It is clear that as the specimen continues to extend towards 8% strain, the dislocations will continue to multiply and as a result a high dislocation density occurs in the material. At room temperature, under creep conditions, dislocation mobility in the structure is significantly reduced, and the creep rate remains very low. However, at 600°C the increased thermal energy means that processes such as climb and cross slip become more prevalent, increasing dislocation mobility. This results in an increased creep rate when compared with the as received (0%

Based on these results, it is clear that the effects of prestrain on the creep performance of the alloy vary significantly with temperature, and as such, dislocation mobility. At low temperatures increased prestrain restricts further creep damage because of the high dislocation densities and apparent difficulty in processes such as climb and cross slip. Conversely, these processes occur more readily at 600°C and increases in prestrain lead to an acceleration in creep damage.

Fig. 19. Effect of prestrain on the fatigue properties of Ti834 at 20°C (Whittaker & Evans, 2009).

However, stress states in the gas turbine are rarely static and as such further consideration must be given to the effect on fatigue performance of the material, Figure 19. In the current work it was found that a small period of stress relaxation (<2 seconds) occurred at the end of the prestrain process before the specimen was unloaded. Previous work (Evans, 1998) has demonstrated that near and titanium alloys tend to form facets under stress relaxation, and fractographic analysis of failed specimens showed that this was indeed the case here. These facets offer initiation sites for fatigue cracks, which along with the increased dislocation density contributes to the 8% prestrain specimens showing significantly shorter fatigue lives.

#### **4. Discussion**

It is clear that whilst the rate of development of new titanium alloys has slowed in recent years, there are further areas which may be explored in order to achieve further improvements in mechanical properties. The research described here has shown that there is definite potential through the harnessing of texture, improved high temperature lifing techniques or improved understanding of processing effects.

Of these, perhaps improvements in high temperature lifing offer designers the greatest reward. Since the development of the gas turbine engine, increased efficiency has acted as the driver which has led to operation at higher and higher temperatures. Enabling components to operate at higher temperatures whilst retaining low density/low cost materials in their manufacture is obviously desirable and operation at temperatures where

Titanium in the Gas Turbine Engine 335

fully understood. Ti6-4 sits between these two more extreme cases, showing cold dwell sensitivity which is often affected by microstructural form. However, since it is an alloy that tends to be used for more low temperature applications than Ti6246 or Ti834, it is clear that

Whilst titanium alloys are under pressure to improve from either lighter, more complex materials (such as composite fan blades) or materials with higher temperature capability (polycrystalline nickel alloys in the HP compressor) it is clear that there are areas of development which have not yet been fully explored, which may offer significant

a. The harnessing of crystallographic texture is capable of providing improved properties, providing the most effective orientations are aligned with direction of loading. It has been shown that despite this anisotropy, predictions of fatigue life can still be made accurately provided the input data for approaches such as the Walker strain method is

b. High temperature lifing allows for extending the operational envelopes of alloys and is an area which requires further research. It has been shown that predictive models can be accurate provided all damage mechanisms (i.e. fatigue, creep and environment) are considered. Whilst the work here is only a start, it is clear that there is scope for further model development/refinement which may result in the safe operation of alloys such

c. Improved understanding of the effect of processing conditions, and the resultant surface finish/damage and residual stress effects in alloys such as Ti834 can lead to increased minimum properties and hence a reduction in safety factors. It has been demonstrated that these effects, represented by prestrained specimens, can significantly alter mechanical properties, and show significant variation with temperature. Whereas room temperature creep rates may be reduced, the creep rates at 600°C are increased and would need to be accurately accounted for in deformation modelling of components Furthermore, fatigue properties at room temperature are reduced, as a result of the formation of quasi-cleavage facets under stress relaxation. This is the type

of critical result which designers must account for when lifing such components.

The author would like to acknowledge funding and financial assistance from EPSRC, Rolls-Royce plc, TIMET UK, Cosworth Racing and QinetiQ during the course of this work.

Bache, MR; Evans, WJ. The effect of tension-torsion loading on low temperature dwell-

Bache, MR; Cope, M; Davies, HM; Evans, WJ; Harrison, G. Dwell sensitive fatigue in a near

Supp. 1, (1997), pp. S83-S88. ISSN: 0142-1123

sensitive fatigue in titanium alloys. *Mechanical Engineering publications* (1996) pp.

alpha titanium alloy at ambient temperature. *International Journal of Fatigue* 19,

a good understanding of the effects is critical for safe utilisation of the alloy.

**5. Conclusions** 

**6. Acknowledgements** 

229-242.

**7. References** 

opportunities for titanium alloys. In particular:-

from the same orientation and of a high enough quality.

as Ti6246 at temperatures in excess of those currently used in service.

creep and environmental damage operate need not be ruled out. Indeed, provided a robust methodology is developed, it need not only apply to titanium alloys, and could be utilised throughout the engine.

It is useful, however, that the method is developed using Ti6246. The alloy is well understood and offers relatively few surprises to design engineers, particularly in its lack of susceptibility to cold dwell. The creep behaviour of the alloy is well detailed and suitably described by techniques such as the theta projection method. It also shows relatively good environmental resistance to approximately 500°C, giving a desirable combination of properties.

The research described here has shown that accurate predictions can be made for the fatigue behaviour of the alloy at high temperature, in the form of a model that described fatigue crack growth. It is recognised that this is only an initial step in the process of developing a total life prediction capability, but it is at least encouraging. Further work would seek to produce strain control deformation behaviour under vacuum conditions in order to isolate the effects of environment, which has been shown to be at least as significant as creep. It is also recognised that the model requires further refinement in order to describe the effects of the type that cause notch lives to be extended as the temperature is raised from 450°C to 500°C. To allow for this the creep deformation should be integrated more closely to fatigue damage in each cycle to describe the stress state of the notched specimen. There is no doubt that these requirements are challenging, but as described, the benefits are clear.

Textured alloys have been trialled for some applications previously, but the current research has demonstrated that it may be possible to widen the field of opportunity. By showing that textured alloys provide a predictable, consistent response, even in the presence of stress raising features, confidence can be gained towards applications in more complex components. Furthermore the work demonstrates that in this alloy, techniques such as the Walker strain approach are able to deal with a wide range of R ratios, although limitations do exist at high R ratios when cold dwell failure mechanisms become apparent. However, it is recognised that this may not be the case across all titanium alloys or process conditions.

The work in fact indicates how cold dwell still acts as a limiting factor in a number of titanium based applications. Essentially it can be seen that a threshold stress exists above which prediction becomes difficult and an extremely shallow S-N curve develops. Recent research at Swansea has shown that in attempting to model this type of behaviour, it is often more appropriate to base predictions on the creep behaviour of the alloy at room temperature, rather than its fatigue response, although further work is still required to characterise behaviour in this area. Such predictions, based on time at a high stress, rather than cyclic fluctuations, have been shown to capture the shape of the curve more accurately and to give reasonable estimates of life.

Indeed, one of the goals of the gas turbine industry is the increased understanding of cold dwell in order to further raise safe operating stresses. However, whilst a good understanding of the mechanisms of facet formation exists (Sinha et. al, 2009, Bache et. al., 1996) further work on the sensitivity of particular alloys would be useful. The three alloys considered here accurately demonstrate the range of effects seen in near titanium alloys. Ti834 has always shown a high sensitivity to cold dwell, particularly in disc form (Bache et. al., 1997) and as such has required designers to carefully consider operating conditions in components where it is utilised. Ti6246 on the other hand, has shown almost no sensitivity to cold dwell (Bache et al, 2007). Whilst this lack of sensitivity is possibly related to the fine Widmanstatten microstructure of the alloy, the mechanisms are still not fully understood. Ti6-4 sits between these two more extreme cases, showing cold dwell sensitivity which is often affected by microstructural form. However, since it is an alloy that tends to be used for more low temperature applications than Ti6246 or Ti834, it is clear that a good understanding of the effects is critical for safe utilisation of the alloy.

## **5. Conclusions**

334 Advances in Gas Turbine Technology

creep and environmental damage operate need not be ruled out. Indeed, provided a robust methodology is developed, it need not only apply to titanium alloys, and could be utilised

It is useful, however, that the method is developed using Ti6246. The alloy is well understood and offers relatively few surprises to design engineers, particularly in its lack of susceptibility to cold dwell. The creep behaviour of the alloy is well detailed and suitably described by techniques such as the theta projection method. It also shows relatively good environmental resistance to approximately 500°C, giving a desirable combination of

The research described here has shown that accurate predictions can be made for the fatigue behaviour of the alloy at high temperature, in the form of a model that described fatigue crack growth. It is recognised that this is only an initial step in the process of developing a total life prediction capability, but it is at least encouraging. Further work would seek to produce strain control deformation behaviour under vacuum conditions in order to isolate the effects of environment, which has been shown to be at least as significant as creep. It is also recognised that the model requires further refinement in order to describe the effects of the type that cause notch lives to be extended as the temperature is raised from 450°C to 500°C. To allow for this the creep deformation should be integrated more closely to fatigue damage in each cycle to describe the stress state of the notched specimen. There is no doubt

Textured alloys have been trialled for some applications previously, but the current research has demonstrated that it may be possible to widen the field of opportunity. By showing that textured alloys provide a predictable, consistent response, even in the presence of stress raising features, confidence can be gained towards applications in more complex components. Furthermore the work demonstrates that in this alloy, techniques such as the Walker strain approach are able to deal with a wide range of R ratios, although limitations do exist at high R ratios when cold dwell failure mechanisms become apparent. However, it is recognised that this may not be the case across all titanium alloys or process conditions. The work in fact indicates how cold dwell still acts as a limiting factor in a number of titanium based applications. Essentially it can be seen that a threshold stress exists above which prediction becomes difficult and an extremely shallow S-N curve develops. Recent research at Swansea has shown that in attempting to model this type of behaviour, it is often more appropriate to base predictions on the creep behaviour of the alloy at room temperature, rather than its fatigue response, although further work is still required to characterise behaviour in this area. Such predictions, based on time at a high stress, rather than cyclic fluctuations, have been shown to capture the shape of the curve more accurately

Indeed, one of the goals of the gas turbine industry is the increased understanding of cold dwell in order to further raise safe operating stresses. However, whilst a good understanding of the mechanisms of facet formation exists (Sinha et. al, 2009, Bache et. al., 1996) further work on the sensitivity of particular alloys would be useful. The three alloys considered here accurately demonstrate the range of effects seen in near titanium alloys. Ti834 has always shown a high sensitivity to cold dwell, particularly in disc form (Bache et. al., 1997) and as such has required designers to carefully consider operating conditions in components where it is utilised. Ti6246 on the other hand, has shown almost no sensitivity to cold dwell (Bache et al, 2007). Whilst this lack of sensitivity is possibly related to the fine Widmanstatten microstructure of the alloy, the mechanisms are still not

that these requirements are challenging, but as described, the benefits are clear.

throughout the engine.

and to give reasonable estimates of life.

properties.

Whilst titanium alloys are under pressure to improve from either lighter, more complex materials (such as composite fan blades) or materials with higher temperature capability (polycrystalline nickel alloys in the HP compressor) it is clear that there are areas of development which have not yet been fully explored, which may offer significant opportunities for titanium alloys. In particular:-


## **6. Acknowledgements**

The author would like to acknowledge funding and financial assistance from EPSRC, Rolls-Royce plc, TIMET UK, Cosworth Racing and QinetiQ during the course of this work.

## **7. References**


**15** 

*South Africa* 

**Platinum-Based Alloys and Coatings:** 

Since platinum has a similar chemistry and atomic structure to nickel, platinum-based alloys are possible contenders for partial or complete substitution of nickel-based superalloys. Although the major disadvantages are high price and density, platinum-based alloys have many advantages, including excellent chemical and oxidation resistance and high strength at high temperatures. Since the melting point is higher than nickel, there is potential that Ptbased alloys can exhibit mechanical properties that surpass those of the nickel-based superalloys. Pt-modified coatings are already employed on turbine blades. These can be modified with the addition of different elements and various coating procedures can be used so that the coatings can complement different substrates. This chapter covers the basic properties of a range of Pt-based alloys and describes the different strengthening mechanisms that exist in these alloys, mainly through a structural approach. The oxidation resistance and corrosion resistance are also described. Further, new alloying additions and

Nickel-based superalloys (NBSAs) have excellent mechanical properties due to precipitation strengthening. The microstructure comprises many small, strained-coherent, particles of the ' phase based on Ni3Al, in a softer matrix of the phase, the solid solution (Ni) of nickel (Sims et al., 1987). The strengthening originates from dislocations being slowed down as they negotiate the small ordered ' particles. There are several mechanisms whereby a unit dislocation has to split into partial dislocations to pass through the ordered precipitates and then re-associate to pass into the random matrix. Each stage requires energy, thus slowing the dislocation movement and providing strengthening. The strengthening depends on the amount of interfacial boundary between the two phases, and is highest when the amount of boundary to be negotiated is highest. This occurs when there are many small precipitates densely distributed, comprising at least 70 vol. % in the alloy (Vattré et al., 2009). For NBSAs, these are usually cubic ' precipitates aligned on the {100} planes (Kear & Wilsdorf, 1962). Additionally, there is solid solution strengthening in the (Ni) matrix, as other elements are dissolved into the nickel, forming a random solid solution. The (Ni) strengthening depends mainly on the difference in elastic and bulk moduli between the Ni and solute atoms (Gypen & Deruyttere, 1981), as well as the size misfit, since each atom varies in size according to its atomic number. There are also bonding effects, all of which

their effects on the structure and properties are identified.

make it more difficult for the dislocations to pass.

**1. Introduction** 

**Materials for the Future?** 

Lesley A. Cornish and Lesley H. Chown *DST/NRF Centre of Excellence in Strong Materials, and School of Chemical and Metallurgical Engineering,* 

 *University of the Witwatersrand,* 

