Advances in gamma titanium aluminides and their manufacturing techniques

Progress in Aerospace Sciences 55 (2012) 1–16

Contents lists available at SciVerse ScienceDirect

Progress in Aerospace Sciences

0376-04

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journal homepage: www.elsevier.com/locate/paerosci

Advances in gamma titanium aluminides and theirmanufacturing techniques

Kunal Kothari 1, Ramachandran Radhakrishnan 2, Norman M. Wereley n

Composites Research Laboratory, Department of Aerospace Engineering, University of Maryland, College Park, MD 20742, USA

a r t i c l e i n f o

Available online 20 July 2012

Keywords:

Titanium aluminide

Intermetallics

Manufacturing techniques

Deformation mechanisms

Microstructure evolution

Alloy development

21/$ - see front matter & 2012 Elsevier Ltd. A

x.doi.org/10.1016/j.paerosci.2012.04.001

esponding author. Tel.: þ1 301 910 2272; fax

ail address: [email protected] (N.M. Wereley

raduate Research Assistant.

siting Assistant Scientist.

a b s t r a c t

Gamma titanium aluminides display attractive properties for high temperature applications. For over a

decade in the 1990s, the attractive properties of titanium aluminides were outweighed by difficulties

encountered in processing and machining at room temperature. But advances in manufacturing

technologies, deeper understanding of titanium aluminides microstructure, deformation mechanisms,

and advances in micro-alloying, has led to the production of gamma titanium aluminide sheets. An

in-depth review of key advances in gamma titanium aluminides is presented, including microstructure,

deformation mechanisms, and alloy development. Traditional manufacturing techniques such as

ingot metallurgy and investment casting are reviewed and advances via powder metallurgy based

manufacturing techniques are discussed. Finally, manufacturing challenges facing gamma titanium

aluminides, as well as avenues to overcome them, are discussed.

& 2012 Elsevier Ltd. All rights reserved.

Contents

1. Titanium aluminides: An overview . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.1. What are intermetallics? . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.2. Importance of titanium aluminides. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.3. Challenges . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.4. Titanium aluminides v. superalloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

1.5. Phases of titanium aluminides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3

1.6. Alloy composition and microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

2. Factors affecting mechanical properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.1. Ductility . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.2. Creep resistance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.3. Fatigue life . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.4. Fracture toughness. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

2.5. Tensile strength . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2.6. General remarks. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

3. Microstructural evolution and phase transformations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

3.1. Near gamma microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

3.2. Duplex microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

3.3. Nearly-lamellar microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

3.4. Fully lamellar microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

4. Crystal structure and deformation mechanisms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

4.1. Crystal structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

4.2. Deformation mechanisms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8

4.3. Polysynthetically twinned (PST) crystals. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

4.4. Effect of lamellar structure on deformation modes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

ll rights reserved.

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K. Kothari et al. / Progress in Aerospace Sciences 55 (2012) 1–162

4.5. Effect of a2-Ti3Al on deformation modes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

4.6. Anomalous yield strength behavior. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

4.7. General remarks. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

5. Alloy development . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

5.1. Effect of Nb. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

5.2. State of the art TiAl alloys . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

6. Manufacturing techniques . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

6.1. Ingot metallurgy and casting. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

6.2. Powder metallurgy. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12

6.3. State-of-the-art in TiAl sheet manufacturing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

6.4. Manufacturing challenges and opportunities . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

7. Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

1. Titanium aluminides: An overview

Titanium aluminides have attracted significant attention in thelast 20 years for their attractive properties that have the potentialto enable high temperature automobile and aerospace applica-tions [1–7]. A foremost application under consideration fortitanium aluminides is high performance gas turbine engines [8].Due to its low density and high strength, titanium aluminideshave become front-runners in replacing nickel-based superalloysin gas turbine engines. Replacement of Ni-based superalloy partswith titanium aluminides is expected to reduce the structuralweight of high performance gas turbine engines by 20–30% [9].Hence, a significant increase in engine performance and fuelefficiency may be realized by utilizing titanium aluminides ingas turbine engines. Titanium aluminides fall in the class ofmaterials known as intermetallics.

1.1. What are intermetallics?

Intermetallics are compounds formed from two metals. Theircrystal structure and properties are completely different fromtheir parent metals [10]. Usually after the formation of anintermetallic alloy, a long range ordering is developed in thematerial. This long range ordering places restriction on thedeformation modes. These restrictions usually are manifested asincreased strength (at least at elevated temperatures), reducedductility and fracture toughness [11]. Apart from titanium alu-minides, examples of other intermetallics include NiAl, and FeAl.

Fig. 1. TiAl turbocharger wheels made by Howmet Corporation [15].

1.2. Importance of titanium aluminides

Titanium aluminides display attractive properties such as lowdensity, high strength, high stiffness and good corrosion, creep,and oxidation resistance. As discussed earlier, due to the longrange order of intermetallics, titanium aluminides lack roomtemperature ductility and fracture toughness. Fatigue crackgrowth rates are also an area of concern [12,13]. In the past, theattractive properties of titanium aluminides were outweighed bytheir difficulties in processing and machining at room tempera-ture. However, advances in manufacturing technologies, deeperunderstanding of titanium aluminide’s microstructure, deforma-tion mechanisms, and advances in micro-alloying, has led to thefirst commercial use of titanium aluminides in high performanceturbochargers for Formula One and sports cars [14]. Exhaustvalves made of titanium aluminides have been used to replacethe existing Ti-6242, Ti-1100, and IMI 834 alloys. Turbochargerwheels made out of titanium aluminides have been prototypedand tested to replace Ni-based superalloys. Till date, there areabout thousand titanium aluminide turbocharger wheels andexhaust valves in service [15,16] Fig. 1.

Gamma titanium aluminides have also been slated for use inGeneral Electric’s GEnex gas turbine engine designed for theBoeing’s 787 Dreamliner [17]. Titanium aluminides are beingused in low pressure turbine (LPT) blades. In addition, titaniumaluminides were also investigated for use in High Speed CivilTransport (HSCT) aircraft [8]. High speed civil transport aircraftwere designed to fly at Mach 2.4, while they take-off and land atconventional airports and meet environmental protection agen-cy’s environmental goals of reduced exhaust and noise pollutions[18]. To meet these stringent requirements, a titanium aluminidebased divergent flap was proposed for noise attenuation andexhaust reduction Figs. 2–4.

1.3. Challenges

Titanium aluminides have found limited, but very challengingcommercial applications in the automobile and aerospace indus-try. A broader requirement for titanium aluminides is to matchNi-based superalloys in performance as well as cost. The densityof titanium aluminides is half that of superalloys, and, hence,replacement of Ni-based superalloys with titanium aluminides isexpected to produce leaner and efficient structural systems. Todate, titanium aluminides lag behind Ni-based superalloys inmechanical performance and significantly in production costs.This is mainly due to its low room temperature ductility and sothat its development has been further held back by a lack ofengineering design practices for low ductility materials. Even so,there have been significant advances made by the NASA GlennResearch Center, Plansee (Austria), and GKSS Research Center(Germany) in developing manufacturing techniques to producetitanium aluminide sheets. As a result, the cost of titanium

Fig. 2. Photos of LPT hardware scheduled for production, Ti–48%Al–2%Cr–2%Nb

(at%): (a) LPT blade casting (b) portion of disk and some blades ready for assembly

[10].

Fig. 3. Structural materials developed with temperature and speed for aircraft as a

function of the expected performances [3].

Fig. 4. Proposed divergent flap prototype fabricated from cast TiAl (Ti–48Al–2Cr–

2Nb) [7].

K. Kothari et al. / Progress in Aerospace Sciences 55 (2012) 1–16 3

aluminide sheets is expected to fall to as low as $150/lb,contingent on its wide acceptance in several commercial applica-tions [19]. Table 1 provides a list of potential applications fortitanium aluminides in the aerospace industry.

3 All the elemental concentrations mentioned henceforth will be in at%, unless

specified otherwise.

1.4. Titanium aluminides v. superalloys

As mentioned earlier, the density of titanium aluminides ishalf that of Ni-based superalloys. Titanium aluminides also dis-play good high temperature oxidation and creep resistance, aswell as excellent high temperature strength and modulus reten-tion. But, titanium aluminides lack room temperature ductilityand fracture toughness. Table 2 compares the mechanical proper-ties of titanium aluminides with Ni-based superalloys at roomtemperature and high temperatures.

As shown in Table 2 [20,21], titanium aluminides are con-siderably (65 times) more costly to fabricate than Ni-basedsuperalloys. Along with this, mechanical properties such as roomtemperature ductility, and yield strength are less favorable thanNi-based superalloys. Over the past 15 years, considerableresearch has been done to refine the microstructure of titaniumaluminides, so that improvements in the ductility, as well as

strength, in titanium aluminides has been realized. Below is ageneral overview of the phases of titanium aluminides.

1.5. Phases of titanium aluminides

Titanium aluminides have been found to exist in three differ-ent phases, namely Ti3Al (a2), TiAl (g), and TiAl3. Of these phases,only the Ti3Al (a2), and TiAl (g) phases have been found to be ofany engineering significance. The phase diagrams in Figs. 5 and 6show the aforementioned phases. As shown, the Ti3Al (a2) phasehas aluminum content between 22% and 39% (at%),3 while theTiAl (g) phase contains 48.5% to 66% (at%) aluminum. A dualphase of titanium aluminides with a mixture of a2 and g phaseexists between 37% and 49% aluminum content [22].

The a2 phase has been found to exhibit good high temperaturestrength, but has very low ductility. Along with this, it also has ahigh rate of Oxygen and Hydrogen absorption, which in turn leadsto further embrittlement at high temperatures [23,24]. On theother hand, the g phase exhibits excellent oxidation resistanceand has very low hydrogen absorption, but its room temperatureductility is close to none. Research has been conducted in refiningthe microstructure of these phases as well as their micro-alloying[25]. But none of the problems associated with these phases havebeen completely solved. Hence, these two phases by themselvesdo not have much of engineering significance. But a mixture ofthese two phases which exists between 40 wt% and 48 wt% ofaluminum has been found to be very viable for several structuralapplications [26,27].

Table 2Comparison of titanium aluminides to superalloys [20,21].

Property TiAl-base alloys Superalloys

Density (gm/cm3) 3.7–3.9 8.3

Room temperature modulus (GPa) 160–176 206

Yield strength (MPa) 400–630 1000

Tensile strength (MPa) 450–700 1200

Ductility at room temperature (%) 1–3 15

Creep limit (1C) 1000 1090

Oxidation (1C) 900–1000 1090

Cost ($/lb) 1300a 20

a Cost is for Gamma Met PX Titanium Aluminide Alloy fabricated by Plansee,

Austria.

Table 1Potential applications for titanium aluminides [18].

Vehicle Component Number of vehicles Projected usage

Global first strike/tomahawk missiles Hot structures, compressor, nozzle 2000þ 30,000 lbsþ

Comanche helicopter Nozzles, IR suppression system 500þ 10,000 lbsþ

F-22, JSF Nozzles, exhaust structures, engine components 700þ 100,000 lbsþ

Military space plane Thermal Protection System (TPS), wingbox and vertical tail hot structures 2þ 5,000 lbsþ

Re-usable launch vehicles (RLVs) TPS, wingbox and vertical tail hot structures 2–8 20,000–80,000 lbs

Fig. 5. Ti–Al phase diagram [9].

Fig. 6. Central portion of the Ti–Al phase diagram [27].


This dual a2þg phase has been found to be very sensitive to itsmicrostructure, grain size, and small volumes of micro-alloyingconstituents [28]. There have been several techniques, whichhave been established to refine the microstructure along withsmall micro-alloying additions by which the dual phase hasexhibited ductility as high as 6% [29]. Additionally, the dual phasealloys have also exhibited room temperature and high tempera-ture strengths equivalent to that of superalloys [30]. The creepand oxidation resistance have also been shown to be acceptablefor temperatures up to 850 1C [3,31].

1.6. Alloy composition and microstructure

The dual phase of titanium aluminides has exhibited a widerange of microstructures depending on the heat-treatment [32].These microstructures have been broadly classified into four cate-gories, namely (1) near-gamma, (2) duplex, (3) nearly-lamellar, and(4) fully lamellar [28,33]. Of these microstructures, the duplex and

fully lamellar microstructure have exhibited mechanical propertiesconducive for a range of commercial applications [34].

As shown in Fig. 6, the fully-lamellar microstructure isobtained by heat-treatment at temperature (T1 in Fig. 6) in thepure a phase field. Upon cooling to room temperature, the a-Tiphase precipitates into alternate plates of a2 and g plates forminga fully lamellar morphology. This microstructure is usuallycharacterized with coarse grains in the range of 200 mm to1000 mm.

The duplex microstructure is obtained by heat-treatment inthe aþg phase field at a temperature (T3 in Fig. 6) where the a/gphase volume ratio is equal to 1. The duplex microstructureconsists of fine fully lamellar colonies along with equiaxedgamma grains. The mixture of these two grain morphologiesforms a very fine microstructure with an average grain size in therange of 10 mm.

The nearly-lamellar microstructure forms at an intermediatetemperature (T2 in Fig. 6) between that of fully lamellar andduplex in the aþg phase field, where the a/g phase volume ratiois greater than 1. It is characterized by a majority of lamellarcolonies with some equiaxed gamma grains forming an averagegrain size in the range of 150 mm to 200 mm.

Finally, the near-gamma microstructure is formed by heat-treatment at much lower temperatures (T4 in Fig. 6) in the a2þgphase field. This microstructure is characterized by equiaxedgamma grains with a2 precipitates forming at the grain bound-aries. The average grain size for this microstructure usually rangesbetween 30 mm and 50 mm.

Due to its very fine grain size, the duplex microstructureexhibits the best ductility and strength at room temperature,but at higher temperatures creep and fatigue resistance is very


low. The fully lamellar microstructure on the other hand iscrippled with low ductility and strength due to its coarse grainsize, but has excellent creep and fatigue resistance properties.Hence, over the past decades, the development of gamma tita-nium aluminides have been geared towards development ofmicrostructures having the best features of duplex and fullylamellar microstructure properties [35–37].

As shown in the phase diagram, at 450% Al, titaniumaluminides exist in the pure gamma phase. The microstructurefor this alloy composition is characterized by equiaxed gammagrains exhibiting no ductility. Even after microstructural refine-ment and small alloy additions this alloy composition shows verylittle improvement in ductility. Therefore, this alloy compositionhas been considered to be of no engineering significance [38].

Fig. 8. Microstructures of a2 created by different thermo-mechanic

Fig. 7. Microstructure types in dual phase titanium aluminides (a) fully

The alloy composition with aluminum concentration o40%results in the formation of a2 (Ti3Al) phase. This phase exhibitsbetter room temperature ductility than alloy composition with Alconcentration 450% (pure g phase). Yet the ductility is unaccep-table from an engineering standpoint. The a2 (Ti3Al) phase, whenalloyed with high Nb content (410 wt%), is known to respond tothermo-mechanical treatment to produce a range of microstructuressimilar to the dual phase (a2þg) alloys. These microstructures asshown in Fig. 8, are fully lamellar, bi-modal and equiaxed.

Similar to dual phase (a2þg) alloys, the range of microstructuresin a2 alloy exhibit a host of desirable properties. The bi-modalmicrostructure exhibits good ductility, but the creep resistance over600 1C is an issue. Fully lamellar microstructure exhibits excellentcreep properties, but its room temperature ductility is very low.

al treatments, (a) Bi-modal (b) equiaxed, and (c) lamellar [10].

lamellar, (b) nearly lamellar, (c) duplex, and (d) near-gamma [37].

Fig. 9. Elongation and Vickers Hardness of Ti–Al binary alloys measured over a

range of Al concentrations (at%). The Vickers hardness is shown for as-cast,

thermo-mechanically processed, and annealed material at room temperature as

well as at 1000 1C [19].


By varying the Nb concentration in the a2 phase, the roomtemperature ductility for a2 phase was increased to as high as4–5%. The ternary a2 alloys, that have been researched widelyinclude Ti–24%Al–11%Nb and Ti–25%Al–10%Nb–3%V–1%Mo (at%).Due to large additions of Nb, the density corrected strength for thesealloys and similar ternary alloys was equivalent to that of conven-tional titanium alloys and Ni-based superalloys [21]. Hence, furtherdevelopment of these alloys has not received much attention.

The alloy composition with Al concentration between 40% and48.5% yield dual (a2þg) phase alloys. As discussed earlier, thesealloys exhibit a range of microstructural morphologies dependingon the thermo-mechanical processing. Several researchers haveshown that ductility of these dual phase alloys increases with theincrease in Al concentration up to 48% and then decreases withincrease in Al concentration beyond 50% [20,38,39]. As shown inFig. 9, Ti–48Al exhibits the maximum ductility in a Ti–Al inter-metallic binary alloy. The specific ductility of the alloy furtherdepends on the microstructure type and grain size, which in turnis governed by thermo-mechanical processing. Hence, the actualductility of the alloys is not listed in Fig. 9.

2. Factors affecting mechanical properties

2.1. Ductility

The duplex microstructure formed with fine lamellar coloniesand equiaxed gamma grains exhibit the best ductility in dualphase alloys. The presence of the lamellar structure aids thedeformation mechanisms of the gamma phase [40]. The presenceof lamellar structure is usually characterized by a ratio of lamellargrains to that of the gamma grains (L/g). In general, the ductilityin the dual phase alloys is dictated by four main factors, (1) grainsize, (2) L/g ratio, (3) changes in lattice dimensions, and (4) impur-ity level [29,38]. As discussed earlier, decrease in grain size leadsto increase in ductility. This is due to the fact that as the grain

size decreases, the volume of defects such as grain boundariesincrease, which in turn aids in the deformation mechanisms. TheL/g ratio of 0.3 to 0.4 produces the maximum ductility in the dualphase alloy. The L/g ratio is further dependent on the a2/g ratio.The a2/g phase ratio of 3% to 15% exhibits maximum ductility.Above this range the grain growth becomes pronounced, whenheat-treated in the aþg phase field and below it the brittle a2

eliminates the beneficial effect of refined microstructure. All thesevalues are mainly governed by the Al concentration and are in theoptimum range, when the Al concentration is 48%. Therefore, Ti–48Al alloy exhibits the best ductility.

The lattice dimension as controlled by the tetragonality ratio(c/a) and unit cell volume has a significant impact on ductility.The g phase crystallizes into the L10 type face centered cubicstructure and has a small tetragonality ratio. This ratio decreasesas the Al concentration is decreased. A decrease in the c/a ratio oran increase in the symmetry of the crystal structure increasesductility in the g phase. Also, decreasing the unit cell volume ofthe lattice increases ductility. This can be achieved by addingsmall ternary or quaternary alloying constituents. An in-depthdiscussion of ternary and quaternary alloying additions is pre-sented in a later section. Finally, by reducing of impurities such asoxygen and nitrogen, the ductility has been observed to increasesignificantly. For example, reduction of oxygen from 800 ppm to370 ppm increased the ductility from 2% to 2.7%.

2.2. Creep resistance

The creep resistance in dual phase titanium aluminides is mainlycontrolled by the microstructural morphology and aluminum con-tent. An increase in Al content usually increases creep resistance.Fully lamellar structures with coarse grain size show better creepresistance over fine-grained duplex microstructures. The increasedcreep resistance for a lamellar structure is attributed to the a2 lathsacting as reinforcements [38]. The creep-rupture strength is higherfor duplex microstructures at temperatures up to 650 1C, abovewhich the lamellar structure shows higher rupture strength. Ingeneral, the fully lamellar microstructure has been found to be moreconducive to creep resistance. But the problem associated with fullylamellar structure is its low room temperature ductility due to itslarge coarse grains. Therefore, much research effort is directedtowards decreasing grain size of the fully-lamellar microstructure[38,40].

2.3. Fatigue life

The factors affecting fatigue are similar to that for creep. Theduplex microstructure shows high fatigue crack growth rates,while the crack growth in fully-lamellar microstructure is a slowprocess [41,42]. Upto 800 1C, the fine-grained duplex microstruc-ture increases fatigue life, but at higher temperatures, the fully-lamellar microstructure exhibits longer fatigue life [11].

2.4. Fracture toughness

The fracture toughness for the fine-grained duplex microstruc-ture is found to be in the range of 10 MPa m½ to 16 MPa m½.These values are much higher for fully-lamellar microstructuresand are found to be as high as 30 MPa m½. It has been observedthat the duplex microstructure exhibits little plastic strain nearthe onset of crack extension and no resistance to crack propaga-tion, whereas the lamellar structure yields large plastic strainsnear the crack tip and increased resistance to crack propagationwith crack length [39].


2.5. Tensile strength

The tensile strength of dual phase titanium aluminides isinversely proportional to the grain size as expected from theHall–Petch relationship [43]. Due to the coarse grain size in fullylamellar microstructures, the tensile strength is low compared tothat of fine grained duplex microstructure. The tensile strengthshows an anomalous behavior at high temperatures, where itshows an increase in strength with increase in temperature up toa certain point and then shows a decrease in strength withincreasing temperature. This behavior has been found to becharacteristic of intermetallics due to their long range ordering.The increase in yield strength with temperature in TiAl can beexplained by the anomalous hardening caused due to cross-slipof dislocations on to octahedral planes. This will be discussedin detail in the section describing the deformation mechanismsin TiAl.

2.6. General remarks

As discussed above, the mechanical properties of dual phasetitanium aluminides are very sensitive to microstructural mor-phology. The duplex microstructure exhibits good room tempera-ture ductility and strength, but for high temperature propertiessuch as creep and fatigue resistance, a fully lamellar microstruc-ture is desirable. The low room temperature ductility andstrength of fully lamellar microstructure is due to its coarsegrains. By decreasing the grain size in fully lamellar structure,the room temperature ductility and strength are expected toincrease. There have been several thermo-mechanical processingroutes suggested to obtain a fine grained fully lamellar structure.To further understand the evolution of microstructures afterthermo-mechanical processing, the knowledge of phase transfor-mations in the dual phase titanium aluminides is important. Thenext section discusses various phase transformation mechanismsin dual phase titanium aluminides.

3. Microstructural evolution and phase transformations

As discussed earlier, the microstructure of the dual phasetitanium aluminides can be broadly classified into four types,(1) near-gamma, (2) duplex, (3) nearly lamellar, and (4) fully-lamellar. In the following paragraphs, the evolution of thesemicrostructures is discussed.

3.1. Near gamma microstructure

The near-gamma microstructure is formed when the materialis heat-treated in the a2þg phase field. Heat-treatment in thisphase field results in the coarsening of the existing g grains. Themicrostructure as shown in Fig. 7d, is characterized by coarsegamma grain regions with fine gamma grain stringer regions withdispersed alpha-2 particles.

3.2. Duplex microstructure

The duplex microstructure is produced by heat treatment inthe aþg phase field. The temperature in the aþg phase field issuch that the a/g volume ratio is close to 1. In this case, the heat-treatment results in the dissociation of the existing a2 particles.Additional a precipitates are nucleated to grow into a plates inthe {1 1 1} habit planes at the expense of gamma phase. Theinitially predominant gamma phase is gradually reduced involume until the equilibrium volume fraction is reached andgrain growth occurs. The growth of gamma grains is limited by

the dispersed alpha phase, which also experiences growth. Thesecompetitive processes result in the formation of a fine-grainedstructure as shown in Fig. 7c.

3.3. Nearly-lamellar microstructure

At temperatures greater or lower than where the duplexmicrostructure forms (i.e., where the a/g volume ratio is closeto 1), coarsening of the predominant phase occurs. Hence, heat-treatment at temperatures below the duplex microstructuretemperature leads to coarsening of gamma grains and formationof near-gamma microstructure, while heat-treatment above theduplex microstructure temperature results in the coarsening ofalpha grains and formation of nearly-lamellar microstructure. Thenearly-lamellar microstructure is characterized by coarse lamellarstructure with fine gamma grains.

3.4. Fully lamellar microstructure

Finally, heat-treatment in the alpha phase field above thealpha-transus line (Ta) results in the formation of large grainedfully lamellar microstructure. The lamellar structure forms inthree different ways and hence is classified into three differenttypes: type I, II, and II [28,29].

Type I lamellar structure is formed by heat-treatment abovethe Ta line followed by air-cooling. It is formed via the followingreaction.

a-aþgp-Lða=gÞ-Lða2=gÞ

In this reaction, plate like gamma precipitates (gp) begin toprecipitate out of alpha matrix at the a/aþg line and growradially to result in the high temperature lamellar structure,L(a/g). This structure transforms at low temperatures to L(a2/g)structure simply by a-a2 ordering reaction. This occurs belowthe a/aþa2 line.

Type II lamellar structure is observed in the duplex micro-structure. In this structure, the alpha-2 plates contain anti-phaseboundaries (APB) which are continuous across the thin gammaplates. For this type of lamellar structure, the nucleation ofgamma precipitates (gpt) and their growth into plates (gp) ispreceded by the a-a2 ordering reaction. The whole process canbe expressed as

a-a2-a2þgpt-a2þgp-Lða2=gÞ

Type III lamellar structure is formed when the heat-treatmentis done well below the duplex microstructure temperature. Herethe predominant phase is g with minor a2 particles (ap

2). Uponheating, the ap

2 in the gamma matrix (gm) disorders to ap andgrow into alpha plates (ap) to yield a lamellar structure L(g/a),which upon cooling transforms to lamellar structure L(g/a2) bysimple a-a2 reaction. The entire reaction can be expressed asfollows:

gmþap2��!

heatinggmþap-Lðg=aÞ��!

coolingLðg=a2Þ

4. Crystal structure and deformation mechanisms

A review of crystal structure and deformation mechanisms oftitanium aluminides gives a deeper insight into the effect ofmicrostructural morphology on its mechanical properties Fig. 10.

4.1. Crystal structure

The g-TiAl phase has (ordered face centered tetragonal) L10

type structure as shown in Fig. 11a. The tetragonality is due to


different atomic radii of Ti and Al. For the stoichiometric com-pound, the tetragonality ratio is 1.02. This tetragonality ratioincreases to 1.03, with increasing aluminum concentration anddecreases to 1.01 with decreasing aluminum concentration [20].The a2-Ti3Al phase has a DO19 (hexagonal closed packed) typestructure as shown in Fig. 11b.

4.2. Deformation mechanisms

It has been established that the main deformation modes ing-TiAl are slip and twinning, both operating on the close-packed{1 1 1} planes [44]. Due to the face-centered cubic type structure ofthe g-TiAl phase, slip occurs on the close-packed {1 1 1} planes. Thelayered arrangement of atoms on the successive (0 0 2) planes andthe tetragonality ratio of 1.02 results in two types of dislocationswith ½/1 1 0S type Burgers vectors on {1 1 1} planes [45].

More specifically these are ordinary dislocations with Burgersvector ½o1 1 0], and superdislocations with Burgers vectors

Fig. 11. Crystal structure for, (

Fig. 10. Dependence of mechanical properties such as fracture toughness,

strength, elongation, impact resistance (IR), creep resistance (CR), and grain size

on the microstructure type, where NG is near-gamma, NL is nearly lamellar and FL

is fully-lamellar [37].

o1 0 1] [46,47]. (For g-TiAl, the miller indices convention wasintroduced by Hug et al., , where, in the notation oh k l] or {h k l]only h and k are mutually permutable, while l is fixed and can bepositive or negative) [48]. In the single phase g-TiAl o1 0 1]superdislocations dominate at low temperatures while at hightemperatures (above about 800 1C) slip by ½o1 1 0] ordinarydislocations, and also twinning, become controlling deformationmodes. In contrast, in the dual phase TiAl with lamellar structure,twinning and glide of ½o1 1 0] ordinary dislocations prevail atlow temperatures and glide of superdislocations becomes signifi-cant only at high temperatures. The ½o1 1 2] type superdisloca-tions have also been reported in TiAl [20]. Fig. 12 shows theordinary ½o0 1 1] dislocations as well as the o0 1 1] and½o1 1 2] type superdislocations.

The L10 structure can be twinned by the {1 1 1}o1 1 2]variants of the normal fcc twinning mode. As shown in Fig. 12,the Burgers vector b3 ¼ 1=6½112� preserves the order of g-TiAl andthis twinning mode is therefore called true twinning. Partialdislocations by Burgers vectors b1 ¼ 1=6½211� and b2 ¼ 1=6½121�change the order of g-TiAl and are called pseudo-twinning.

At room temperature, the o0 1 1] type superdislocations havevery limited mobility due to the covalent nature of Ti–Ti, andTi–Al bonds. Their mobility is further inhibited due to theformation of extrinsically faulted dipoles. Segments of the trailingsuperpartials, 1/6[112] type, form faulted dipoles, which must beextended as deformation proceeds [49]. This explains the lowroom temperature ductility in single phase g-TiAl, where o0 1 1]superdislocations are the dominant deformation modes at roomtemperature. o0 1 1] type superdislocations can be dissociatedinto two ½o0 1 1] dislocations separated by an antiphaseboundary (APB). The o0 1 1] type superdislocation can be furtherdissociated by

101h i

¼1

6112h i

þ1

2101h i

þ1

6211h i

a) g-TiAl, and (b) a2-Ti3Al.

Fig. 12. Potential slip and twinning systems in the g-TiAl {1 1 1} planes (Circles of

varying sizes indicate atoms on different parallel {1 1 1} planes) [44].


with superlattice intrinsic stacking fault (SISF) between 1=6½112�and 1=2½101� partials, and the complex stacking fault (CSF)between 1=2½101� and 1=6½211� partials [50].

In dual phase TiAl, the presence of a2-Ti3Al as secondary phaseresults in ½o1 1 0] ordinary dislocations and twinning tobecome the dominant deformation modes. Twinning is mainlycaused by the presence of lamellar structure in dual phase TiAl[51,52]. As discussed earlier, the lamellar structure is formed byheat-treatment in the a or aþg phase field. The formation of gplates from a is described above. The nucleation from the a phaseproceeds from the stacking fault by the movement of 1/3/1 0 1 0S Shockley partials [21,29]. The resulting lamellar struc-ture has the crystallographic orientation relationships ½0 1 1�g==½1 1 2 0�a2

9ð1 1 1Þg==ð0 0 0 1Þa2between the alpha-2 plates

and gamma plates, with gamma plates being twin related. The[101], and [110] directions are not equivalent to [0 1 1] directionon {1 1 1} plane due to the tetragonality in the L10 structure of gphase. Directions of /1 1 2 0Son the basal plane in the a phase(hexagonal closed packed) and a2 phase (hexagonal D019) are allequivalent as shown in Fig. 12. Hence, six interfaces arise from thecombination of the three possible directions of the c-axis of the gphase and the ordered nature of the two phases [40].

The mechanical properties of the lamellar microstructures inTiAl depend on the lamellar orientation with respect to theloading axis and lamellar microstructure variables such as grainsize, thickness, and spacing of g and a2 lamellae and g domainsize. However, the lamellar orientation has far more influencethan lamellar microstructural variables. In order to fully under-stand the effect of lamellar orientation on its properties, TiAlcrystals where the entire crystal consists of only a single lamellargrain were produced and characterized. Since numerous thin twinrelated lamellae are contained in the major constituent g phase,these crystals are named polysynthetically twinned (PST) crystalsfrom analogy with the phenomenon ‘‘polysynthetic twinning’’,observed in mineral crystals.

4.3. Polysynthetically twinned (PST) crystals

As mentioned above, the PST crystals have a single lamellargrain. The mechanical properties of the PST crystals depend stronglyon the lamellar orientation relative to the loading axis. PST crystalsalso display anisotropic macroscopic flow behavior. When theorientation of the lamellae is perpendicular to that of the loadingaxis, the PST crystal exhibits excellent strength and poor ductility. Inthe parallel orientation of the lamellae with respect to the loadingaxis, some ductility and strength was observed. The strength in thiscase is not as high as the perpendicular orientation case. Thesedeformation modes in parallel and perpendicular directions arereferred to as hard modes.

At the intermediate orientations, such that the lamellae forman angle of 301–601 with respect to the loading axis (soft mode),the yield stress is much lower but the elongation is significantlyhigher than the hard modes. The orientation dependence of theyield stress and ductility of PST crystals is due to the fact thatshear occurs parallel to the lamellar boundaries when thelamellar orientation is such that it forms an angle of 301–601with the loading axis (deformation in soft mode) but it occursmostly on {1 1 1} planes intersecting the lamellar boundarieswhen the lamellae is either parallel or perpendicular (hard mode)to the loading axis.

4.4. Effect of lamellar structure on deformation modes

In dual phase TiAl with duplex and fully lamellar microstruc-ture morphologies, the presence of a2-Ti3Al phase is usually in theform of lamellar structure. In a lamellar structure, the orientation

of the lamellae with respect to loading axis plays a significantrole in the control of deformation mechanisms. As discussedabove, with more lamellae oriented in the soft mode, the ductilityincreases. Hence, it is desirable to have maximum lamellaeoriented in the soft mode. In different grains lamellae will beoriented at different angles relative to stress-axis, and when grainsize is large, it is likely that fewer grains are oriented favorably forslip. Hence, fully lamellar structures with coarse grains sufferfrom low ductility.

On the other hand, duplex microstructure with fine grainsexhibit good ductility. Due to the presence of more grains per unitvolume, the probability of grains with lamellae oriented in thesoft mode will increase. As a result, the ductility in duplexmicrostructure is higher than that in fully lamellar microstruc-ture. A decrease in grain size has been correlated to increase inductility by several researchers [39]. The duplex microstructuredoes not perform well in terms of creep and fatigue resistance athigh temperatures compared to that of fully lamellar microstruc-ture. Hence, it has been suggested by several researchers thatrefining the grain size in fully lamellar microstructures couldpotentially solve the problem of low ductility.

As discussed earlier, fully lamellar microstructures are formedby heat-treatment in the a phase field. The diffusion process inthe a phase field is very rapid due to thermal activation and theabsence of second phase barriers. Therefore the grain growth isalso rapid. In order to circumvent this, several thermo-mechanicaltreatments along with small alloy additions have been proposed.These will be discussed in the next sections.

4.5. Effect of a2-Ti3Al on deformation modes

As mentioned earlier, in single phase g-TiAl the deformationmodes are dominated by o1 0 1] type dislocations. Ordinarydislocations of ½o1 1 0], and twinning are hardly observed. Butin two phase (a2þg) TiAl, where the dominant phase is g-TiAl,the deformation modes are dominated by ordinary dislocations½o1 1 0] and twinning. This phenomenon has been explained interms of a2-Ti3Al acting as a gettering agent to improve the purityof g-TiAl phase [49].

In other independent studies it has been concluded that byreducing interstitials and impurities in nominally pure g-TiAlphase, the directionality of the bonds between the Ti atomsreduces. This in turn leads to increased mobility of dislocationswith ½o1 1 0], and 1/6o1 1 2] Burgers vectors. The 1/6o1 1 2]Burgers vectors are order twinning defects. It has been proposedthat in dual phase (a2þg) TiAl, where the dominant phase isg-TiAl, the a2-Ti3Al laths getter the interstitial impurities fromtheir adjacent g-TiAl laths. Hence, an increase in the mobility ofdislocations with ½o1 1 0], and 1/6o1 1 2] Burgers vectors isobserved in dual phase (a2þg) TiAl. The solubility of interstitialimpurities in a2-Ti3Al is greater than that in g-TiAl.

4.6. Anomalous yield strength behavior

In single phase g-TiAl, anomalous increase in flow stress isobserved with increase in temperature up to a certain point. Thisincrease in strength usually begins to occur around 200 1Cand peaks between 600 1C and 700 1C depending on the alloycomposition. As discussed earlier, in single phase g-TiAl roomtemperature deformation is controlled by o0 1 1] superdisloca-tions, which have limited mobility. As the temperature increases,ordinary dislocations ½o1 1 0] and twinning become active andare relatively mobile, but the o0 1 1] dislocations becomeincreasingly blocked as they adopt a thermally activated non-planar configuration and thereby give rise to the flow anomaly[53]. These o0 1 1] superdislocations dissociate into a primary


plane with stacking faults and APBs as discussed earlier. With theincrease in temperature, segments of the o0 1 1] cross-slip ontothe adjoining octahedral planes [54]. Since these cross-slippedsuperdislocations are highly immobile, an increase in flow stresswith increase in temperature occurs. But as temperatureincreases further, the o0 1 1] superdislocations dissociate intounit ½o1 1 0] and ½o1 1 2] superdislocations, which are suffi-ciently mobile such that the strength decreases following the flowstress peak.

In the dual phase TiAl, similar behavior is observed except thatthe temperature range of anomalous yield strength behavior is abit contracted and flow stress peaks appear at a bit lowertemperatures. The deformation mechanism in dual phase TiAlresponsible for such behavior is the interaction of ½o1 1 0]screw dislocations on adjoining {1 1 1} planes. These interactionsbecome frequent at higher temperatures and as a result anincrease in strain is followed by formation of jogs, dislocationdipoles and debris defects [55]. The ½o1 1 0] dislocations arepinned by these defects and the number of pinned dislocationsincreases with the temperature. At higher temperatures due tothermal activation, the jog-dislocation climbing occurs, whichresults in decreased strength with temperature [56].

4.7. General remarks

The deformation mechanisms in TiAl have been investigatedover years and have been helpful in explaining the effect ofmicrostructural morphology especially that of lamellar structureon the mechanical properties of TiAl. From the above discussions,several conclusions can be made regarding approaches to enhancemechanical properties in TiAl.

i.
The brittle nature of single phase g-TiAl is due to dominantnature of the o1 0 1] type dislocations, which are primarilysessile due to extrinsically faulted dipoles.
ii.
The glissile slip systems of twinning and ordinary dislocations,1/6o1 1 2] and ½o1 1 0], respectively become active due tothe presence of a2-Ti3Al. The secondary a2-Ti3Al phase acts asa gettering agent to improve the purity of g-TiAl phase andhence enables slip via ordinary dislocations and twinningdominant.
iii.

Table 3Effect of selected alloying elements on mechanical properties of TiAl.

Element Effect

The deformation modes in lamellar structures are dependenton its orientation (soft mode or hard mode) with respect tothe loading axis. Fully lamellar microstructures with goodroom temperature ductility and high strength can be devel-oped by refining its grain size.

The discussion so far has been limited to the binary TiAl alloys.Significant accomplishment has been made in enhancing themechanical properties of TiAl by small alloy additions. The nextsection will look into various alloy additions explored in TiAl sofar and its effect on mechanical properties. Along with this, threestate-of-the art alloys developed by several commercial vendorsare discussed.

Nb Increases oxidation and creep resistance in small amounts, also

increases high temperature strength if added between 5% and 10%

Ta Increases oxidation and creep resistance and tendency for hot

cracking

V Increases ductility

W Oxidation and creep resistance

B Grain refiner

C Increases creep and oxidation resistance

Cr Increases ductility if added in small amounts; increases oxidation

resistance if added in the range of 8%

Mn Increases ductility

Mo Increases strength, and creep and oxidation resistance.

5. Alloy development

The dual phase (a2þg) TiAl have shown good response tosmall alloying additions. In general, the alloying additions for TiAlcan be broadly divided into three different categories [44] withthe following compositions (in at%):

Ti45-52�Al45-48�X1-3�Y2-5�Zo1

where X, Y and Z represent the three categories of alloyingadditions. These alloying elements more or less affect the position

of the phase boundaries in the Ti–Al binary phase diagram[20,28,57,58].

The X additions are usually elements which increase theductility of TiAl. They lower the stacking fault energy and therebyincrease the propensity of twinning. Usually elements such as Cr,Mn, and V fall in this category.

The Y additions such as Nb, Ta, W, and Mo increase theoxidation and creep resistance at high temperatures in TiAl. Theeffect of Nb in varying levels is discussed in more depth in thenext section.

Finally, the Z alloying elements are usually B or C. Boron isprimarily used as a grain refiner to produce fully lamellarmicrostructures with fine microstructures. Since, B is not solublein TiAl, it acts as an impediment to grain growth in the a phasefield to form fully lamellar structures. Carbon additions furtherincrease the strength of TiAl due to Ti3AlC perovskite precipitates.C also increases the creep resistance of TiAl. The Ti3AlC perovs-kites act as glide obstacles with long-range stress fields, whichcannot be overcome with the aid of thermal activation [31].Table 3 lists the mechanical properties enhanced by the variousalloy additions.

5.1. Effect of Nb

Nb is especially used to increase oxidation and creep resistance.Over the years, alloys with Nb additions have been developed andcan be classified into four different types [31,58,59].

i.
Ti–48Al–2Nb ii. Ti–(46–47)Al–(2–3)Nb
iii.
Ti–45Al–(5–10)Nb iv. Ti–45Al–(5–7)Nb–RM (refractory metals)
Alloys with high Nb content with type III and IV are alsoknown to increase the high temperature strength by solid solu-tion strengthening. Recently large amounts of Nb in the order of5% to 10% have been used in TiAl alloys to form a new class of TiAlalloys known as TNB alloys [9,60,61,62,63,64] and TNMalloys [65]. In these types of alloys, Nb plays a critical role inincreasing high temperature strength along with creep andoxidation resistance.

The Nb atoms occupy Ti atomic sites and reduce the Al contentin the gamma phase. This shifts the a phase boundary to the leftand thereby lowers the processing temperatures to produce fullylamellar structures. With low processing temperatures, graingrowth can be controlled, which leads to the refinement of themicrostructure and increase in strength.

Nb has also been reported to increase the tetragonality ratio ofTiAl, which enhances its structural anisotropy and hence itsstrength [63]. Recently, Nb has also been reported to increase


the intrinsic strength of the gamma phase in Ti–45Al alloys [66].Twinning in TiAl has been found to be enhanced by Nb, which inturn suppresses pre-mature failure due to the inherent brittlenessof TiAl [67]. Along with strengthening effects, Nb is also known toincrease the diffusion activation energy in TiAl. This leads tocontrol of high creep rates in TiAl at temperatures above 700 1C.

Initially, Nb was used primarily to increase the oxidationresistance of TiAl alloys [68]. This was achieved in type I and IIalloys listed above by small Nb additions. The a2 phase has agreater solubility for oxygen than the g phase. At high tempera-tures, absorption of oxygen by a2 results in further embrittlementof TiAl. More recently it has been reported that Nb decreases thea2 phase in TiAl alloy and hence decreases its oxygen solubility[66]. Also, Nb promotes the formation of Al2O3 which in turnincreases the oxidation resistance of TiAl.

5.2. State of the art TiAl alloys

The alloy development in TiAl over the last 20 years has resultedin four state-of-the-art alloys with exceptional high temperaturemechanical properties. These four alloys are listed in Table 4 withtheir compositions and strengths [11,26,27,30,69,70,71,72].

Fig. 13. TiAl ingots with coarse columnar grains ready for homogenization

process.

6. Manufacturing techniques

Investment casting, ingot metallurgy (IM), and powder metal-lurgy (PM) are popular techniques that have been used to produceTiAl parts. More recently, advanced techniques such as directrolling [19,73,74], laser forming, and mechanical alloying havebeen investigated with good success. Several rapid sintering/consolidation techniques such as plasma pressure compaction,spark plasma sintering, pulse discharge sintering, high densityinfrared processing and explosive consolidation have also beensuccessful in forming TiAl with desired mechanical properties[75–82].

One of the biggest drawbacks of TiAl compared to that ofNi-based superalloys is its production cost. This is largely due tothe fact that processing techniques for materials with fairly lowductility do not exist. Further, due to the long-range ordering ofTiAl up to its melting point, their processing temperatures arefairly high. This requires capital investment in processing equip-ment with good high temperature characteristics.

Investment casting, ingot metallurgy and powder metallurgytechniques have been successful at producing TiAl parts withdesirable mechanical properties only after a series of post proces-sing steps, like hot-isostatic pressing, ageing, annealing, and hotworking. This further adds to the production costs of TiAl.

Advanced techniques such as direct rolling, laser forming, andspark plasma sintering have been aimed at reducing the post-processing steps for TiAl. These processes usually require lesstime compared to the traditional techniques and can formcomplicated shapes without much post-processing. But the pro-blems associated with these techniques involve porosity andscalability Fig. 13.

Table 4TiAl state-of-the-art alloys.

Alloy name Composition (at%)

General Electric, USA: 48-2-2 [11,27] Ti–48Al–2Cr–2Nb

Plansee, Austria: g–MET [30,69] Ti–45Al–(5–10)Nb

GKSS Research Center, Germany: TNB Alloy [26,70] Ti–(45–47)Al–10Nb

Martin Marietta Laboratories, USA: XDTM TiAl [71,72] Ti–45Al–2Mn–2Nb–0.8B

6.1. Ingot metallurgy and casting

Ingot metallurgy and casting routes involve producing TiAlingots and castings by skull melting [64]. The resulting micro-structure of the ingots is characterized by large columnar grainsconsisting of chemical inhomogeneities, and segregation asshown in Fig. 14 [20]. The differences in the melting points andthe densities of the constituent elements and peritectic solidifica-tion process leads to segregation and chemical inhomogeneities[83]. The chemical inhomogeneities are seen even on a micro-scopic level from one lamellar to the next within a single lamellargrain [84]. Further, the solidification process occurs at the purealpha phase field as shown in Fig. 6. Hence, the grain growth isvery rapid and leads to formation of large columnar grains [85].

Significant improvements in chemical homogeneity and micro-structure refinement can be achieved by thermo-mechanicalprocessing and the associated dynamic recrystallization [86]. Asdiscussed earlier, the thermo-mechanical processing conditions,alloy chemistry and microstructure evolution have an intimatecorrelation [87]. Hence, a tight optimization of the alloy compositionand thermo-mechanical processing parameters is required toachieve a more homogeneous crystallization and refined micro-structure [88].

Several thermo-mechanical routes such as hot-rolling, forging,and extrusion have been successfully utilized on titanium alumi-nide ingots to produce parts with refined microstructure andimproved chemical homogeneity [30,32,89,90]. These thermo-mechanical treatments are performed in the (aþb) phase field(shown in Fig. 6) [91]. This is usually in the temperature range of

Alloy strengths

Ductility, fracture toughness, and oxidation resistance

High temperature strength, creep, fatigue, and oxidation resistance

High temperature strength, creep, and oxidation resistance

Ductility, high temperature strength, stiffness, creep, and oxidation resistance

Fig. 14. Processing routes required for the production of high-pressure aero-engine compressor blades from a TiAl ingot [30].


1200 1C to 1350 1C. The thermo-mechanical treatments con-ducted between 1200 1C and 1250 1C result in a fine-grainedduplex microstructure with the grain size as low as 5 mm [30].At temperatures close to 1350 1C a fully lamellar microstructurewith average colony size of 100 mm is obtained. The colony size isusually controlled by introducing small amounts of Boron asdiscussed earlier.

In order to achieve chemical homogeneity, the primary TiAlingots and castings produced by melting have to be free ofimpurities. Hence, as an alternative to skull melting, several othercleaner processes such as vacuum arc remelting, and cold-hearthplasma arc melting have been employed. In vacuum arc remelting[92], large rods consisting of given amount of Ti, Al and masteralloys are cold welded to each other in an inert atmosphere. Thecold welded electrode is then melted in a vacuum arc furnace. Theelectrode is melted to a primary ingot, which is then remeltedtwice to improve chemical homogeneity and reduce segregation.

Finally, in the cold-hearth plasma arc melting process, the metalalloy is melted and poured into a water-cooled copper hearthforming a solid skull. This acts like a secondary hearth of the samecomposition as the parent alloy. Without the solid skull, harmfulrefractories could be picked up and cause melt-related inclusions.This also allows enough time for inclusions to sink into the skull ordissolve into the melt before the alloy is poured into a water-cooledcopper crucible for solidification. For this reason, this process isreferred to as a ‘‘clean melting’’ process. Cold-hearth melting is analternative method to other melting techniques, such as vacuum arcremelting, where alloy cleanness is of great significance. Electronbeams and plasma-arc heating are both effective heat sources thatare suitable for this type of clean melting.

As mentioned above, most of the ingot metallurgy and castingtechniques result in the formation of coarse columnar grainsconsisting of chemical inhomogeneity and segregation. This issolved by further thermo-mechanical treatments. The ingots areeither extruded or forged to desired shapes with further heattreatments to relieve any residual stresses. Fig. 14 illustrates thepost-processing steps required for forming a TiAl compressorblade. Apart from thermo-mechanical treatments, the ingots andcastings are also subjected to hot-isostatic pressing (HIP) toimprove homogeneity and refine microstructure. The microstruc-ture is further refined by ageing treatments [28].

6.2. Powder metallurgy

Powder metallurgy (PM) offers the potential for minimizingmany of the problems associated with large ingot production and

reducing the overall cost of the final TiAl component. Many of theproblems associated with ingot metallurgy (IM), such as center-line porosity, chemical inhomogeneity, regions of varying densityand microstructure can be solved by powder metallurgy. Further,powder metallurgy enables the development of new alloys thatcannot be made by conventional ingot metallurgy [93].

Gas atomization of pre-alloyed powder [94] followed by hotisostatic pressing or consolidation by extrusion to full density hasbeen the general route taken in powder metallurgy. Mechanicalalloying (MA) of the elemental powders to form the alloy powdersof TiAl has also been investigated [95–97] with success inobtaining metastable to stable TiAl phases. One of the biggestproblems associated with mechanical alloying is contamination ofthe powders from the milling media and the container.

Hot isostatic pressing (HIP) of powder compact is a verypopular powder metallurgy based method to produce TiAl billets[59,98–104]. In this process, gas atomized powder of high purityis canned. The can is usually made of commercially pure Ti.The can is then degassed at �500 1C under high vacuum afterwhich the can is sealed gas tight. The canned powder compact issubjected to isostatic pressure in the range of 100MPa to 150 MPaat a temperature range of 1200 1C to 1400 1C depending on thedesired microstructure. This is done for about 2 h to 6 h. Theconsolidated powder compact is de-canned to produce billets.Some of the post processing steps after hot isostatic pressinginvolve hot extrusion, isothermal forging and hot rolling to sheets.In hot isostatic pressing, the porosity in the final part is limitedand the microstructure is homogeneous with little segregation.The drawback of hot isostatic pressing is that it requires exposureto high temperatures for several hours. This leads to grain growth.Hence, there is limited flexibility in hot isostatic pressing tocontrol the grain size of the final part.

In powder metallurgy, the presence of impurities in thepowders in the form of interstitial elements leads to porosity[103,105]. Gas atomization process has been successful in produ-cing high purity TiAl powders. In this process, the molten metal isproduced by either consumable or nonconsumable arc melting.The metal is subsequently atomized using an annular gas flownozzle. For gas atomization of TiAl alloys only inert gases such asArgon or Helium can be used because of the highly reactive melt.However, during atomization the liquid metal interacts with theinert gas environment, leading to the entrapment of gas withinthe particles in spherical pores [106–110].

For comparison purposes, sheets of TiAl alloys of similarchemical composition have been produced by powder metallurgyand ingot metallurgy routes [8]. In powder metallurgy, gas


atomized TiAl powder was consolidated into a prematerial rec-tangle. The prematerial was then canned, and hot-rolled into TiAlsheets. In the ingot metallurgy process, an ingot of TiAl wasforged into a pancake prematerial. This pancake was thenmachined into rectangle shape, canned, and hot rolled into sheets.The yield from the ingot metallurgy process was found to bemuch lower than that of powder metallurgy process. Also, thepowder metallurgy route has been found to be more cost-effective than IM. But, areas of microporosity have been foundin powder metallurgy processed TiAl sheets, and consequentlythis has been found to limit the strength of the material [8].Hence, it is very important that the process of powder synthesis isvery clean and inhibits the presence of impurities.

Most of the current powder metallurgy processes have beensuccessful in producing TiAl parts, which will require further hot-working or machining to produce the final component. Metalinjection molding (MIM) and spray forming are two powdermetallurgy methods, which strive to produce near-net shapeparts with minimal post processing [111]. The basic steps inmetal injection molding are the kneading of a feedstock bymixing the powders with a binder, usually consisting of waxand polymers; the molding of the part; the extraction of most ofthe binder components and final sintering in order to obtainhighly dense parts. Porosity and contamination are some ofthe challenges in using metal injection molding for producingTiAl parts.

In spray forming, a melt stream is atomized using low pressureargon gas. The droplets and partly solidified droplets are gatheredon a substrate which is positioned at a certain distance belowthe gas nozzle. In this way, high density deposits with highchemical and microstructural homogeneity as well as low puritylevels can be obtained. There has been very limited work doneon spray forming of titanium aluminides and hence the data onthe spray formed components is very limited to determine itssuccess [112].

6.3. State-of-the-art in TiAl sheet manufacturing

Plansee, Austria has been a pioneer in producing TiAl sheetmaterial using both ingot metallurgy and powder metallurgytechniques. For powder metallurgy based technique, gas atomizedpowders are consolidated by hot isostatic pressing and subse-quently rolled on a conventional hot-rolling mill using Plansee’spatented Advanced Sheet Rolling Process [113,114]. For ingotmetallurgy based technique, TiAl ingots produced by vacuum arcremelting are treated by hot isostatic pressing and rolled byAdvanced Sheet Rolling Process. The microstructural homogene-ity and mechanical properties of these sheets have been found tobe consistent and superior to most TiAl based alloys. The cost forthese sheets is currently at $1300/lb, but is expected to go as lowas $150/lb over the years.

An alternate approach has been used to produce thin g-TiAlsheets at a lower cost [3]. In this approach, cast g-TiAl is directlyrolled into thin sheets, thereby eliminating costly and wastefulintermediate steps. Direct rolling of cast plates into thin sheetshas been demonstrated for a number of g-TiAl alloys. This processhas been developed by NASA Glenn research center in Cleveland,Ohio [19]. For direct-rolling, the as-cast TiAl alloy does notundergo additional process steps such as atomizing, hot isostati-cally pressing, extruding or conditioning, prior to being encapsu-lated. Once the TiAl alloy is cast as a preform, the as-cast TiAlalloy preform is encapsulated and directly rolled to form TiAlsheets [115]. Limited microstructural and mechanical propertyevaluations on these sheets have produced encouraging results. Ithas been estimated that the direct rolling may lead to a cost

reduction of �35% over the conventional powder metallurgy andingot metallurgy routes.

6.4. Manufacturing challenges and opportunities

The manufacturing techniques for TiAl via both ingot metal-lurgy and powder metallurgy routes have been advanced enoughto produce TiAl parts with desirable mechanical properties. As aresult, these processes have found use in producing TiAl parts forseveral high temperature applications in the aerospace andautomobile industry. But the cost associated with TiAl, especiallydue to its post processing steps after ingot production or powderconsolidation, has outweighed its benefits of weight savings.Hence, advanced near-net shaping manufacturing methods withminimal post-processing methods are being investigated for TiAl.

Also, methods to refine fully-lamellar microstructures are beingactively investigated. As discussed earlier, fully-lamellar micro-structures with fine grain sizes and lamellar spacing may possessmechanical properties with acceptable room temperature ductilityand toughness and excellent high temperature characteristics suchas creep and fatigue resistance. Several thermo-mechanicaltreatments have been employed to achieve this. These thermo-mechanical treatments are quite involved due to the complexthermodynamic Ti–Al system. Achieving thermo-dynamic equili-brium in fine spaced lamellae with small grain sizes requirescarefully designed heat-treatments. Fine grained microstructuresin TiAl have a tendency to deteriorate into Widmanstatten struc-tures [70] even before reaching service temperature of the TiAlpart. Therefore, a lot research has concentrated on reducing thegrain size of the material by powder metallurgy techniques. Herethe starting powder particle size is reduced to a considerableextent so that the grain size in the consolidated sample is alsosmall. With this technique the grain size can be lowered into thenanometer range [116]. A decrease in the grain size up to thenanometer range has the potential of increasing the strength andductility of the material significantly [21].

High energy ball milling and attrition milling have been usedto reduce powder particle size of the TiAl powders [116–118].Pre-alloyed TiAl powders have been ball-milled for up to 50 h toform amorphous TiAl powders with particle size in the sub-micrometer to nanometer range. One of the advantages of millingis that it disorders the lattice to improve the dislocation motion.This is especially important in TiAl, where due to its long rangeordering its ductility is limited by immobility of the superdisloca-tions as discussed earlier [119].

Mechanical alloying (MA) route has also been used to reducepowder particle size [74,120–122]. For mechanical alloying, pureTi and Al powders are ball-milled for several hours to form a Ti–Alsolution, which is partly disordered and its particle size is reducedconsiderably from the starting elemental powder particle size.

Several studies have been conducted to consolidate the milledpowders to produce TiAl with ultra-fine grains in the sub-micronto nanometer range [123–125]. The consolidation processes usedfor this were typical sintering or hot isostatic pressing processes.Due to the small particle size, diffusion in the milled powders isvery rapid and hence grain growth is very rapid too. The graingrowth can be controlled by rapid consolidation of the powders.

Finally, metal matrix composites (MMCs) based on TiAl havealso been considered as potential solution to enhance its hightemperature service range. TiAl based metal matrix compositesreinforced with TiC, TiB2, TiB, Al2O3 and TiN [121,126–131] havebeen considered. TiAl metal matrix composites reinforced withtitanium borides have been very popular due to the grain refiningcharacteristics of Boron. XDTM TiAl alloy with needle-shaped TiB2

particulates have been produced via casting process and havedisplayed excellent room temperature ductility and high


temperature characteristics. But several studies have focused onusing more economical approaches such as mechanical alloying.In mechanical alloying, the reinforcement and matrix powders aremechanically alloyed and consolidated to produce the desiredTiAl metal matrix composite.

7. Conclusion

Titanium aluminides have made strong gains in terms ofacceptance in the aerospace and automobile industries. Although,a wide acceptance of titanium aluminides in various industries isstill lacking, several improvements in manufacturing and proces-sing techniques have made titanium aluminide materialspotential candidates for several high temperature applications.Research over the past two decades has led to a deeper under-standing of the inter-relationships between microstructural mor-phology, phase evolution, and deformation mechanisms. This inturn has led to the development of advanced powder metallurgybased (PM) manufacturing techniques, which negotiate the lowworkability of titanium aluminides. Advanced powder metallurgybased novel manufacturing opportunities exist to produce near-net shape parts of TiAl with enhanced mechanical properties.Along with this, several new challenges emerge in terms ofcharacterization of a whole new set of TiAl based alloys producedby these novel powder metallurgy based techniques. Also, chal-lenges lay ahead in the implementation of these techniques forthe industrial production of near-net shape TiAl parts.

Besides solving manufacturing issues, it is also important thatnew processing techniques address the issues of developingtitanium aluminides with a balance of desirable mechanicalproperties. The new processing techniques can probably achievethis balance by incorporating some of the mechanical propertyenhancement mechanisms such as, (1) Control of microstructuralmorphology, (2) Small alloying additions, such as Nb, Cr, and B,(3) Reduction of grain size to sub-micrometer to nanometerrange, and (4) Reinforcement with discontinuous particulates orfibers.

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Advances in gamma titanium aluminides and their manufacturing techniques

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