Gamma Titanium Aluminide Alloys - Fritz Appel - E-Book

Gamma Titanium Aluminide Alloys E-Book

Fritz Appel

214,99 €


The first book entirely dedicated to the topic emphasizes the relation between basic research and actual processing technologies. As such, it covers complex microstructures down to the nanometer scale, structure/property relationships and potential applications in key industries. From the contents: * Constitution * Thermophysical Constants * Phase Transformations and Microstructures * Deformation Behaviour * Strengthening Mechanisms * Creep * Fracture Behaviour * Fatigue * Oxidation Resistance and Related Issues * Alloy Design * Ingot Production and Component Casting * Powder Metallurgy * Wrought Processing * Joining * Surface Hardening * Applications and Component Assessment

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Table of Contents


Related Titles

Title page

Copyright page


Figures – Tables Acknowledgement List

1 Introduction

2 Constitution

2.1 The Binary Ti–Al Phase Diagram

2.2 Ternary and Multicomponent Alloy Systems

3 Thermophysical Constants

3.1 Elastic and Thermal Properties

3.2 Point Defects

3.3 Diffusion

4 Phase Transformations and Microstructures

4.1 Microstructure Formation on Solidification

4.2 Solid-State Transformations

5 Deformation Behavior of Single-Phase Alloys

5.1 Single-Phase γ(TiAl) Alloys

5.2 Deformation Behavior of Single-Phase α2(Ti3Al) Alloys

5.3 β/B2 Phase Alloys

6 Deformation Behavior of Two-Phase α2(Ti3Al) + γ(TiAl) Alloys

6.1 Lamellar Microstructures

6.2 Deformation Mechanisms, Contrasting Single-Phase and Two-Phase Alloys

6.3 Generation of Dislocations and Mechanical Twins

6.4 Glide Resistance and Dislocation Mobility

6.5 Thermal and Athermal Stresses

7 Strengthening Mechanisms

7.1 Grain Refinement

7.2 Work Hardening

7.3 Solution Hardening

7.4 Precipitation Hardening

7.5 Optimized Nb-Bearing Alloys

8 Deformation Behavior of Alloys with a Modulated Microstructure

8.1 Modulated Microstructures

8.2 Misfitting Interfaces

8.3 Mechanical Properties

9 Creep

9.1 Design Margins and Failure Mechanisms

9.2 General Creep Behavior

9.3 The Steady-State or Minimum Creep Rate

9.4 Effect of Microstructure

9.5 Primary Creep

9.6 Creep-Induced Degradation of Lamellar Structures

9.7 Precipitation Effects Associated with the α2→γ Phase Transformation

9.8 Tertiary Creep

9.9 Optimized Alloys, Effect of Alloy Composition and Processing

9.10 Creep Properties of Alloys with a Modulated Microstructure

10 Fracture Behavior

10.1 Length Scales in the Fracture of TiAl Alloys

10.2 Cleavage Fracture

10.3 Crack-Tip Plasticity

10.4 Fracture Toughness, Strength, and Ductility

10.5 Fracture Behavior of Modulated Alloys

10.6 Requirements for Ductility and Toughness

10.7 Assessment of Property Variability

11 Fatigue

11.1 Definitions

11.2 The Stress–Life (S–N) Behavior

11.3 HCF

11.4 Effects of Temperature and Environment on the Cyclic Crack-Growth Resistance

11.5 LCF

11.6 Thermomechanical Fatigue and Creep Relaxation

12 Oxidation Behavior and Related Issues

12.1 Kinetics and Thermodynamics

12.2 General Aspects Concerning Oxidation

12.3 Summary

13 Alloy Design

13.1 Effect of Aluminum Content

13.2 Important Alloying Elements – General Remarks

13.3 Specific Alloy Systems

13.4 Summary

14 Ingot Production and Component Casting

14.1 Ingot Production

14.2 Casting

14.3 Summary

15 Powder Metallurgy

15.1 Prealloyed Powder Technology

15.2 Elemental-Powder Technology

15.3 Mechanical Alloying

16 Wrought Processing

16.1 Flow Behavior under Hot-Working Conditions

16.2 Conversion of Microstructure

16.3 Workability and Primary Processing

16.4 Texture Evolution

16.5 Secondary Processing

17 Joining

17.1 Diffusion Bonding

17.2 Brazing and Other Joining Technologies

18 Surface Hardening

18.1 Shot Peening and Roller Burnishing

18.2 Residual Stresses, Microhardness, and Surface Roughness

18.3 Surface Deformation Due to Shot Peening

18.4 Phase Transformation, Recrystallization, and Amorphization

18.5 Effect of Shot Peening on Fatigue Strength

18.6 Thermal Stability of the Surface Hardening

19 Applications, Component Assessment, and Outlook

19.1 Aerospace

19.2 Automotive

19.3 Outlook

Subject Index

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The Authors

Dr. habil. Fritz Appel

Helmholtz-Zentrum Geesthacht

Institute for Materials Research

Max-Planck-Str. 1

21502 Geesthacht

Dr. Jonathan David Heaton Paul

Helmholtz-Zentrum Geesthacht

Institute for Materials Research

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21502 Geesthacht

Dr. Michael Oehring

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Foreground shows the General Electric GEnx-1B engine (photo courtesy of General Electric). The background is an artificially colored high resolution TEM image of a deformation twin/matrix interface in TiAl.

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There is an ever-increasing demand for the development of energy-conversion systems towards improved thermodynamic efficiency and ecological compatibility. Advanced design concepts are based on higher service temperatures, lower weight, and higher operational speeds. For example, the operating efficiency of a gas-turbine engine will increase by over 1% for every 10 °C increase in the turbine-inlet temperature. Substantial fuel savings in aircraft and power generation can be achieved through the introduction of new materials that can provide higher temperatures or reduced component weight. The conventional metallic systems that are currently in use have been developed over the last 50 years to near the limits of their capability. If further advances are to be made, new classes of materials will be required.

Titanium aluminide alloys based on the intermetallic gamma phase are widely recognized as having the potential to meet the design requirements mentioned above. Undoubtedly, the development of such a material system has important implications for spin-offs to other high-temperature technologies, as well as for the general economy. For example, General Electric has recently made public that its most recent engine, the GEnx, includes the use of titanium aluminide as a blade material. This is a significant milestone for a relatively new, advanced engineering material.

Although there is a vast body of TiAl literature going back over 20 years, there have only been a few review articles published in the recent past, the latest nearly a decade ago. Since that time, considerable advances have been made, both in the basic understanding of the physical metallurgy and in processing technology. It is our intention that the publication of this book will, for the first time, give a wide-ranging interpretation and discussion of the voluminous amounts of data documented in the literature. For TiAl to be successfully employed as a structural material requires a comprehensive understanding of the complex microstructures, down to the nanometer scale, and knowledge concerning how the structure–property relationships are determined by, for example, the atomic details of interface-related phenomena.

The overview of all relevant research topics that are presented in this book is intended to form a link between scientific findings and alloy development, material properties, industrial processing technologies, and engineering applications. The metallurgy of TiAl alloys undoubtedly has several features in common with other intermetallic system. Thus, in that we have chosen to emphasize the scientific principles, the book will provide a treatment of the subject for researchers and advanced students who need a more detailed coverage than is found in physical metallurgy textbooks. We expect that our compilation of the current state of titanium aluminide science and technology will not only serve as a guide through the huge body of literature to the TiAl community, but will also be of interest to materials scientists, engineers, and technical managers who are involved in areas where low-density, high-temperature resistant materials are required. The detailed description of interfaces and interface related phenomena will certainly be of interest to an extended scientific community.

It would not have been possible to write such a book without the help and support from numerous people and organizations. First, we would like to acknowledge the generous support and the excellent research conditions provided by the Helmholtz-Zentrum Geesthacht (formerly GKSS) under its Scientific Director Prof. Wolfgang Kaysser, Prof. Andreas Schreyer as the Director of the Institute for Materials Research, and Prof. Florian Pyczak as group leader.

We also thank the BMBF (German Ministry for Education and Research), DFG (German Science Foundation), Helmholtz Gemeinschaft (Helmholtz Association), Rolls-Royce Deutschland, and CBMM (Companhia Brasileira de Metalurgia e Mineração) for financial support through their funding of numerous research projects.

We would particularly like to thank Prof. Richard Wagner (now Director at the Institute Laue-Langevin, Grenoble, France) who initiated the work on TiAl in the late 1980s while he was director of our institute. Additionally, we would like to thank our colleagues and former students, Ulrich Brossmann, Stefan Eggert, Dirk Herrmann, Roland Hoppe, Ulrich Fröbel, Viola Küstner, Uwe Lorenz, the late Johann Müllauer, Thorsten Pfullmann, and Ulf Sparka for their interest, support, and for contributing to an excellent group atmosphere. The generous help from the HZG library personnel is also acknowledged.

A very special mention must be made to acknowledge Dr. Young-Won Kim (Universal Energy Systems, Dayton, USA) for his achievement in keeping the titanium aluminide community together for very many years and his friendship. Fritz Appel would like to thank his wife, Bärbel, for her support. Finally, the authors would like to expresss their gratitude to Wiley-VCH for the opportunity to write the book and in particular gratefully acknowledge the patient support by Waltraud Wüst and Ulrike Werner and careful copyediting of Bernadette Cabo.

Fritz Appel

Jonathan David Heaton Paul

Michael Oehring

Geesthacht, January 2011

Figures – Tables Acknowledgement List

In order to cover the wide range of literature, the authors have used copies or slightly modified copies of figures/micrographs from previously published work. Where this has been done the figure has been referenced so that the source paper, authors, and journal of the original work are credited. The table below is intended to indicate and thank the publishers, companies, or individuals who hold the copyright to these figures and acknowledges their generous permission for reuse.

Publishing sourceFigures/tablesACCESS e.V.Reprinted with permission, Copyright ACCESS.Figure: 14.21.American Physical Society (APS)Reprinted from “Phys. Rev. B” with permission, Copyright (1998) APS.Figures: 3.1 and 3.2.ASM International®.Reprinted from “Castings, Metals Handbook” with permission, Copyright ASM. www.asminternational.orgFigure: 14.25.Cambridge University Press.Reprinted with permission, Copyright Cambridge University Press.Figures: 5.7, 12.1, 15.15, and 16.47.Deutsche Gesellschaft für Materialkunde (DGM). Reprinted with permission, Copyright DGM.Figure: 15.12.General Electric Company (Aviation).Reprinted with permission, Copyright GE.Figure: 1.2 and cover.Elsevier Publishing.Reprinted with permission, Copyright Elsevier.Figures: 1.1a, 2.5–2.10, 3.2, 4.1, 4.13–4.15, 6.6, 6.26, 6.30, 6.31, 6.60, 6.65, 6.69, 6.70, 6.72–6.74, 6.77, 6.78, 7.6, 7.12, 7.13, 7.20, 7.21, 7.23, 7.25, 8.1–8.3, 8.9, 8.10, 9.3, 9.27, 9.31–9.33, 10.14, 10.17, 10.25, 10.26, 11.5, 11.6, 11.8, 11.9, 11.11–11.13, 11.16, 11.17, 12.6, 12.8, 12.9, 12.11, 12.12, 12.16, 13.4, 13.5, 13.16, 14.12–14.15, 14.26, 15.7, 15.16, 16.57, 18.2–18.4, 18.5, 18.14, 18.15, 19.3, and 19.4. Tables: 6.3, 7.4, 11.1, and 11.2.IOP Publishing.Reprinted with permission, Copyright IOP Publishing.Figure: 7.2.The Japanese Institute of Metals (JIM).Reprinted with permission, Copyright JIM.Figure: 9.2.Metal Powder Industries Federation (MPIF), 105 College Road East, Princeton, New Jersey, USA. Reprinted with permission, Copyright MPIF.Figures: 7.1, 15.2, 15.6, 15.8, and 15.9.Springer Publishing.Reprinted with permission, Copyright Springer.Figures: 2.2–2.4, 2.11, 6.67, 7.15, 7.16, 10.16, 15.19–15.21, 16.3–16.5, 16.8–16.10, 16.12, 16.14–16.17, 16.29, 16.44–16.46, 16.48–16.50, 16.52, 16.58, and 17.4–17.16. Table: 7.6.Taylor & Francis Publishing.Reprinted with permission, Copyright Taylor & Francis.Figures: 5.11, 5.12, 5.16, 5.20–5.22, 6.11, 6.49, 6.50, 6.53, 6.59, 7.27, 10.8, 16.21, and 16.22.The Minerals, Metals and Materials Society, Warrendale, PA, (TMS).Reprinted with permission, Copyright TMS.Figures: 4.2–4.4, 4.6, 6.29, 6.49, 7.17, 9.1, 9.22, 9.24, 10.13, 10.15, 10.18, 10.21, 10.30, 10.31, 11.2–11.4, 11.7, 12.15, 13.2, 13.3, 14.2, 14.3, 14.5, 14.6, 14.8, 14.9, 14.16–14.20, 15.3, 15.11, 15.17, 15.18, 16.11, 16.18, 16.55, 19.1, and 19.2.Wiley Publishing.Reprinted with permission, Copyright Wiley.Figures: 6.40, 6.42, 10.2, 12.2–12.5, 12.7, 12.10, 12.14, 14.1, 15.10, 15.13, 15.14, 16.33, 16.35–16.37, 16.39 and 16.53. Table: 14.1.G. Hug, PhD Thesis, Université de Paris-Sud, France, 1988.Figure: 5.9.



The reason why gamma TiAl has continued to attract so much attention from the research community including universities, publicly funded bodies, indus­trial manufactures, and end-product users is that it has a unique combination of mechanical properties when evaluated on a density-corrected basis. In particular, the elevated temperature properties of some alloys can be superior to those of superalloys.

Dimiduk [1] has assessed gamma TiAl with other aerospace structural materials and shown that new capabilities become available on account of its properties. The most important pay-offs involve

high melting point;low density;high specific strengths and moduli;low diffusivity;good structural stability;good resistance against oxidation and corrosion;high ignition resistance (when compared with conventional titanium alloys).

Figure 1.1 shows how the specific modulus and specific strength of gamma TiAl alloys compare to other materials. As a result of these properties TiAl alloys could ultimately find use in a wide range of components in the automotive, aero-engine and power-plant turbine industries.

Figure 1.1 Graphs showing the (a) specific moduli and (b) specific strengths of TiAl and other structural materials, as a function of temperature [1]. The data indicates that TiAl compares favorably with the other materials. The data has been redrawn based on the original diagrams.

For a material to be ready for introduction, the whole production chain and supplier base, from material manufacture through processing and heat treatment must have achieved “readiness”. This includes detailed knowledge of how component properties are related to alloy chemistry, microstructure, and processing technology. In addition, TiAl-specific component design and lifing methodologies need to be developed and give reliable predictions [2]. At the implementation stage no unforeseen technical problems concerning the processing route or component behavior, which may be very costly to remedy, should arise. In 1999, a time when fuel costs were relatively low compared to the current day, Austin [3] discussed how introduction of gamma would depend on economic viability. This was identified as the chief obstacle for the use of gamma, with marketplace factors dominating implementation decisions.

Due to its intermetallic nature, the complex constitution and microstructure, and the inherent brittleness, the physical metallurgy of TiAl alloys is very demanding. Nevertheless, we will attempt to discuss the broad literature that has been published over the last two decades concerning synthesis, processing and characterization. In our opinion, significant advances have been made, in particular General Electric has made public its intention [4, 5] to use gamma TiAl in its latest engine, the GEnx-1B (Figure 1.2), which best illustrates the present state that has been achieved in TiAl technology. Gamma TiAl has also been successfully introduced into at least one automotive series production, used in formula 1 racing engines, and a variety of components have been manufactured and successfully tested. In the following chapters we will present a comprehensive assessment of both the science and the related technology that has enabled TiAl to be used in the real world.

Figure 1.2 The General Electric GEnx-1B engine for the Boeing 787 Dreamliner. The blades in the last 2 stages of the low-pressure turbine in this engine are made from cast TiAl, making this engine the first to use TiAl in the real world.

Photo courtesy of General Electric.


1 Dimiduk, D.M. (1999) Mater. Sci. Eng., A263, 281.

2 Prihar, R.I. (2001) Structural Intermetallics 2001 (eds K.J. Hemker, D.M. Dimiduk, H. Clemens, R. Darolia, H. Inui, J.M. Larsen, V.K. Sikka, M. Thomas, and J.D. Whittenberger), TMS, Warrendale, PA, p. 819.

3 Austin, C.M. (1999) Curr. Opin. Solid State Mater. Sci., 4, 239.

4 Weimer, M., and Kelly, T.J. Presented at the 3rd international workshop on γ-TiAl technologies, 29th to 31st May 2006, Bamberg, Germany.

5 Norris, G. (2006) Flight International Magazine.



2.1 The Binary Ti–Al Phase Diagram

The binary Ti–Al phase diagram contains several intermetallic phases, which represent superlattices of the terminal solid solutions and have been recognized to be an attractive basis for lightweight high-temperature materials for many years [1]. However, over the past two decades of research, only alloys based on the α2(Ti3Al) phase with the hexagonal D019 structure or the γ(TiAl) phase with the tetragonal L10 structure (Figure 2.1) have emerged as structural materials. Among these, interest has been strongly focused on γ titanium aluminide alloys, which, for engineering applications always contain minor fractions of the α2(Ti3Al) phase. Further, the high-temperature β phase with the bcc A2 structure and its ordered B2 variant (Figure 2.1) play a significant role in some engineering alloys. Despite intensive research, the binary Ti–Al phase diagram still remains a matter of debate and thus has been the subject of experimental work and critical assessment in recent years [2–5]. The impact of such work cannot be overestimated as a full understanding of the microstructural evolution in an alloy is limited by knowledge of the relevant phase equilibria. The discrepancies between different versions of the phase diagram might predominantly have been caused by the high sensitivity of phase equilibria to nonmetallic impurities, in particular oxygen [6–8]; but experimental difficulties, for example, problems in the identification of superlattices and sluggish phase transformations [9] may also play a role. Historically, investigations date back to the 1920s, as reported by Mishurda and Perepezko [6] and Schuster and Palm [2]. The first phase diagrams that covered the whole concentration range were published in the 1950s [10, 11]. More information on the historical development can be found in the article by Mishurda and Perepezko [6] cited earlier. The first critical and thorough assessment of the binary phase diagram was published by Murray [12] and has been considered as a standard reference [2]. Although Murray’s assessment does not reflect current knowledge on the phase equilibria in the Ti–Al system, it is a very useful compilation of phase diagram and physical data. Published experimental data was recently reassessed in a comprehensive study by Schuster and Palm [2]. This publication, together with the thermodynamic re-evaluations [3, 7–9, 13–15] constitute the current state of knowledge on this phase diagram. Figure 2.2 shows the phase diagram that was constructed by Schuster and Palm [2] after their critical assessment of all available experimental data. Figures 2.3 and 2.4 show sections of this diagram covering Al concentrations that are relevant for titanium aluminide alloys based on the γ(TiAl) phase. Experimental data has been plotted on these sections to give an impression on the reliability of phase boundaries. Figure 2.5 shows the most recent result for a thermodynamic evaluation of the binary Ti–Al phase diagram [3] obtained using the CALPHAD approach [16, 17]. Despite many similarities this diagram differs from that of Schuster and Palm [2], in particular for Al concentrations above 60 at.%. These discrepancies mainly arise due to the occurrence of the Ti3Al5 and Ti2+xAl5−x phases in the diagram of Witusiewicz et al. [3] and will be briefly discussed below. The crystallographic data of the stable and metastable phases considered by Schuster and Palm [2] and Witusiewicz et al. [3] are given in Table 2.1. Thermodynamic data and data on phase equilibria can be found in these two recent publications as well as in the references cited therein.

Table 2.1 Crystallographic data of phases occurring in the binary Ti–Al system.

The data has been taken from the publications of Schuster and Palm [2] and Witusiewicz et al. [3] and partially supplemented. 1d-APS: one-dimensional antiphase domain structures.

Figure 2.1 Crystal structures of binary Ti aluminide phases. (a) Hexagonal α2(Ti3Al) phase (Strukturbericht designation D019, prototype Ni3Sn, Pearson symbol hP8, space group P63/mmc), (b) tetragonal γ(TiAl) phase (Strukturbericht designation L10, prototype AuCu, Pearson symbol tP4, space group P4/mmm), (c) cubic high temperature B2 phase (Strukturbericht designation B2, prototype CsCl, Pearson symbol cP2, space group ). As explained in the text, and by comparing Figures 2.2 to 2.4 with Figure 2.5, the occurrence of the B2 phase in binary alloys is not yet fully clarified.

Figure 2.2 Binary Ti–Al phase diagram according to the assessment of Schuster and Palm [2].

Figure 2.3 Section of the binary Ti–Al phase diagram according to the assessment of Schuster and Palm [2].

Figure 2.4 Section of the binary Ti–Al phase diagram according to the assessment of Schuster and Palm [2].

Figure 2.5 Calculated binary Ti–Al phase diagram according to the thermodynamic evaluation of Witusiewicz et al. [3].

As already mentioned, many details of the Ti–Al phase diagram have been debated for a long time. One prominent example is the peritectic reaction L + β → α, which is particularly important for an understanding of the solidification of γ alloys but was not accepted in the assessment of Murray [12]. However, the work of McCullough et al. [32], Mishurda et al. [33], Mishurda and Perepezko [6], Kattner et al. [7] and Jung et al. [34] clearly confirmed this reaction. A further question that has attracted attention is whether the α2 phase is formed congruently from α or in a peritectoid reaction from α + β as shown in Figures 2.2, 2.4 and 2.5. The work of Kainuma et al. [35] together with other studies, cited by Schuster and Palm [2], has shown that the peritectoid reaction and, as a consequence, a second peritectoid reaction β + α2 → α occurs. However, other studies by Veeraghavan et al. [36] and Suzuki et al. [37] cast some doubts. Another issue that has been discussed in the literature concerns the question as to whether the β phase occurs as an ordered B2 variant in the binary phase diagram. The B2 ordering has been proposed by Kainuma et al. [35] and would result in a better agreement of the β phase boundary with the experimental results. The two possible versions of the phase diagram are seen in Figures 2.2 and 2.5. By combining DSC measurements with a theoretical extrapolation of the ordering temperature from ternary phase diagrams, Ohnuma et al. [8] confirmed that the β phase shows a second-order transition to the B2 phase. In contrast, Suzuki et al. [37] found no experimental evidence for B2 ordering. Schuster and Palm [2] concluded that B2 ordering in binary alloys is rather unlikely but that it could not be definitely ruled out. This subject could be the topic of future research work.

In Al-rich alloys containing 65 to 72 at.% Al that were quenched from temperatures around 1215 °C, the presence of so-called one-dimensional antiphase domain structures (1d APS) was observed and were designated as Ti5Al11, γ2, Ti2Al5 or long-period superstructures [2]. However, according to the authors it is an open question as to whether the structures are the result of a second-order transition, or if narrow two-phase regions occur in the phase diagram, or if the 1d-APS are transient metastable structures. This open question, together with the not unambiguous existence of the Ti3Al5 phase [3] results in the discrepancies between the phase diagrams shown in Figures 2.2 and 2.5. Nevertheless, these phase diagrams reflect fairly well current knowledge concerning the constitution of binary Ti–Al alloys and the remaining uncertainties.

Finally, some specific features of the Ti–Al phase diagram should be noted. Engineering alloys based on the γ(TiAl) phase usually have Al concentrations of 44 to 48 at.% and thus solidify, according to the phase diagram, either through the β phase or peritectically. Depending on processing conditions and alloy composition, it is even possible that two peritectic reactions could occur. Indeed small differences in the Al concentration can result in very different solidification microstructures and textures. Further, the existence of one (α) or two (β and α) single-phase high-temperature regions is a characteristic feature of γ(TiAl) alloys. Similar to steels after heat treatment in the austenite region, a variety of different phase transformations can occur during cooling from high-temperatures or after subsequent heat treatments. In principle, this enables one to obtain a wide range of microstructures [38, 39]. The complexity of possible phase transformations is further increased in multicomponent alloys and has only just begun to be systematically studied. Thus, the mechanical properties may be tailored to some degree using conventional metallurgical processing and the limited damage tolerance of γ alloys may be controlled to some extent. It should be further mentioned that the eutectoid transformation α → α2 + γ that takes place on cooling, occurs in all engineering γ(TiAl) alloys. The mechanism of this reaction seems to be identical to the reaction α → α + γ that proceeds via nucleation and growth of single γ lamellae. To maintain thermodynamic equilibrium in alloy compositions that deviate from the eutectoid composition, the volume fraction of the γ phase has to increase abruptly when the temperature falls below that of the eutectoid transformation. However, the cooling rates usually employed are often too fast to result in thermodynamic equilibrium, and therefore the microstructures obtained may be not stable at the intended service temperature of around 700 °C. Moreover, since the final microstructure in γ(TiAl) alloys is formed during cooling from high-temperature heat treatments, the variation of the volume fraction of γ and the other phases with temperature is a general problem that can result in nonequilibrium phase constitutions. For this reason, a microstructural stabilization treatment or appropriate cooling conditions should be considered.

2.2 Ternary and Multicomponent Alloy Systems

Over the past two decades a broad range of engineering alloys has emerged from a number of alloy development programs, each being developed with respect to different processing routes and applications. The alloys can be described by the general composition:


with X designating the elements Cr, Nb, V, Mn, Ta, Mo, Zr, W, Si, C, Y and B [40–43]. For the alloy compositions described by Equation 2.1, the α2 phase usually exists in addition to the γ phase and sometimes other phases. Considering the effect of the various alloying elements on the constitution, two alloying strategies can be distinguished. Alloying elements that go into solid solution can be added to γ(TiAl) alloys. Such additions can influence the properties of the γ phase; like the energies of planar defects or the diffusion coefficient. In contrast, other alloying elements are aimed towards the formation of third (or even further) phases to obtain for example, precipitation hardening, grain refinement during casting, a stabilization of the microstructure against grain growth, or transient phases that decompose into fine structures. In Table 2.2 the crystallographic data of some stable and metastable phases occurring in ternary or multicomponent systems are given. When the effect of alloying additions on the intrinsic properties is considered, the solubility in the γ phase, partitioning of alloying elements between the α2 and the γ phases, and the occupation of the alloying element on the two sublattices of the γ phase are of interest. Information on the site occupancy of different alloying elements in the γ phase can be found in studies by Mohandas and Beaven [58], Rossouw et al. [59] and Hao et al. [60]. In an interesting study, Hao et al. [61] have shown that the lattice occupancy of different alloying elements is related to the boundaries of the (α2 + γ) two-phase field in the respective (Ti–Al–X) systems. A more detailed discussion of site occupancy and its influence on mechanical properties can be found in Chapter 7. With respect to solubility in the γ phase, most of the metals mentioned above are only soluble up to a limited amount of around 2 to 3 at.%, as the respective ternary phase diagrams show [62–68]. For higher concentrations, the bcc β solid solution or its ordered variant with a B2 structure usually form as a third phase. A good overview on the constitution of ternary alloy systems in relevant composition ranges has been given by Kainuma et al. [68]. In a comprehensive study the authors investigated the phase equilibria between the α, β and γ phases for most of the Ti–Al–X systems according to the general composition described by Equation 2.1. Figure 2.6 shows the three-phase equilibria at 1200 °C obtained in this study. From this figure it is obvious that the alloying elements Zr, Nb and Ta are exceptional in respect that they are soluble in higher contents in the γ phase, in the case of Nb up to around 9 at.% at 1200 °C [68]. The study of Kainuma et al. [68] is also interesting because attention was paid to the partitioning coefficient between the different phases. With respect to partitioning between the α2 and the γ phases, V, Cr, Mo, Ta and W were enriched in the α2 phase, Nb and Mn were distributed in equal amounts between the α2 and γ phases, whereas Zr was concentrated in the γ phase. The elements Fe, Cr, Mo and W strongly partitioned to the β phase, with respect to both the α2 and the γ phases. In Figures 2.7 and 2.8 a number of isothermal sections from important ternary phase diagrams are displayed that were published by Kainuma et al. [68]. These diagrams give an impression regarding the constitution of the different ternary systems. In recent years, some alloy development programs have been directed to alloy systems in which higher amounts of β phase occur [69–76]. However, lack of knowledge regarding the constitution of such multicomponent systems is one of the major obstacles to alloy development and significant work has to be expended to obtain reliable information on phase relationships. This originates from the fact that the constitution of multicomponent systems can be intricate and rich in detail. Similar to the case for the binary Ti–Al system some ternary phase diagrams, such as the Ti–Al–Nb diagram, have remained a matter of debate for a long time [61, 66, 67, 77–88]. With regard to other ternary systems a vast body of valuable information can be found in the literature, for example, Ti–Al–Cr [62, 63, 67, 68, 89, 90], Ti–Al–Mo [64, 65, 68, 71, 91], Ti–Al–V [61, 68, 71, 92, 93], Ti–Al–Mn [61, 68, 94], Ti–Al–Ta [61, 68, 95], Ti–Al–W [68], Ti–Al–Fe [61, 68, 96, 97], Ti–Al–Si [98, 99], Ti–Al–O [67, 100, 101], Ti–Al–N [102], Ti–Al–C [103] and Ti–Al–B [104–106].

Table 2.2 Crystallographic data of some phases arising in ternary or multicomponent TiAl alloys of engineering relevance.

Figure 2.6 Position of the α + β + γ three-phase field at 1200 °C for different ternary Ti–Al–X systems [68].

Figure 2.7 Isothermal sections of different ternary Ti–Al–X systems [68]. (a) Ti–Al–V, (b) Ti–Al–Cr, (c) Ti–Al–Mn, (d) Ti–Al–Fe.

Figure 2.8 Isothermal sections of different ternary Ti–Al–X systems [68]. (a) Ti–Al–Nb, (b) Ti–Al–Mo, (c) Ti–Al–Ta, (d) Ti–Al–W.

As mentioned above, Nb is soluble in comparatively large amounts in the γ as well as the α2 phase and has been found to be particularly advantageous for γ(TiAl) engineering alloys [107–111]. The properties of such alloys with high Nb contents are addressed in Chapters 7 and 13. Here, it should be mentioned that the Ti–Al–Nb phase diagram is complicated since several ternary compounds with compositions that are not far away from the α + γ + β(B2) and α2 + γ + B2 three-phase fields occur. The occurrence of these compounds is one reason for the discrepancies between the different constitutional studies mentioned above. It is also noteworthy to mention that multiphase assemblies including orthorhombic phases were observed in high Nb-containing alloys [73] and these were not predicted by the phase diagrams available at that time. See Chapter 8 for more details on such alloys. In Figures 2.9 and 2.10 the constitution of the ternary Ti–Al–Nb system is depicted for temperatures and compositions that are relevant for engineering applications. The isothermal and vertical sections that are shown were taken from the most recent thermodynamic re-evaluation of the Ti–Al–Nb and the constituent binary systems [86] and thus represent current knowledge on the constitution of Ti–Al–Nb alloys.

Figure 2.9 Calculated isothermal section of the Ti-Al-Nb system at 700 °C according to the most recent thermodynamic re-evaluation [86]. The table below indicates which phases are present in the phase diagram:

Figure 2.10 Calculated vertical sections close to the Ti-Al side of the Ti-Al-Nb system according to the most recent thermodynamic re-evaluation [86]. (a) Isopleth for 45 at. % Al, (b) isopleth for 8 at. % Nb. The table below indicates which phases are present in the phase diagrams:

Besides Nb and Cr the most widely used alloying element in γ(TiAl) alloys is B. This is due to the grain-refining effect that B additions within the range of 0.1 to 2 at.% have on cast alloys [113, 114]. Additionally, B additions have also been found to be useful with respect to the hot-working behavior and the microstructure of wrought products (see Chapter 16). B is almost insoluble in all binary phases of the Ti–Al system [106, 115]. In an interesting study by Hyman et al. [104] it was shown that, depending on the Al concentration, either the β or α phase are the primary solidifying phase for B contents of 0.7 to 1.5 at.% (Figure 2.11). This finding is in agreement with very recent work by Witusiewicz et al. [106]. Taken together these results provide evidence that the simplest explanation for the mechanism of grain refinement in alloys with B concentrations less than 0.7 at.%, that is, that the borides serve as nuclei for the binary TiAl phases, is not correct. Thus, a full understanding of grain refinement in TiAl alloys, and higher multicomponent systems, requires detailed knowledge on the constitution. More on this subject can be found in Chapter 4.

Figure 2.11 Suggested liquidus projection for the ternary Ti–Al–B system in the vicinity of the equiatomic TiAl composition [104]. Dashed lines denote schematic solidification of some alloys.

To conclude this chapter, the influence of O on the constitution of TiAl alloys will be briefly considered. As the available ternary Ti–Al–O phase diagrams show [67, 100], the solubility of O is much higher in the α2 phase as compared to the γ phase. This has been confirmed by atom probe (AP) analysis, a method that is capable of quantitatively determining the O content with a very high spatial resolution, in fine α2 and γ lamellae for example. Using this technique, Nérac-Partaix et al. [116] showed that the solubility of O in the γ phase was about 300 at. ppm for Ti–Al two-phase alloys, with the α2 phase containing between 8000 and 22 000 at. ppm O depending on the Al concentration [116, 117]. In alloys with an Al content of 52 at.% that did not contain any α2 phase, the same solubility of O as in two-phase alloys was determined, even if the specimens contained 1 or 2 at.% O [116]. The authors proposed that the difference in the solubility of O in the α2 and the γ phase might originate from the presence of octahedral sites, surrounded by six titanium atoms, in the structure of the α2 phase [117]. The high solubility of O in the α2 phase obviously prevents the formation of oxides, which occur in higher Al containing alloys when α2 is not present, as early has been concluded by Vasudevan et al. [118]. The significantly different solubility of O in the two phases results in the phase boundaries between the α2 and the γ phases depending on the O content of the material. Similar results were obtained for alloys that contained additions of Cr, Mn and Nb [117, 119]. Taken together, these findings confirm the high sensitivity of binary Ti–Al and ternary Ti–Al–X systems on the O content. This also has been concluded indirectly from other studies [6–8].


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