M A T E R I A L S C H A R A C T E R I Z A T I O N 6 1 ( 2 0 1 0 ) 1 0 2 9 – 1 0 3 4

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Deformation-induced α2→γ phase transformation inTiAl alloys

C.L. Chena, W. Lub, Dai Sunc, L.L. Hea,⁎, H.Q. Yea

aShenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road,Shenyang 110016, PR ChinabThe Hong Kong Polytechnic University, Hung Hom, Kowloon, Hong Kong, PR ChinacInstitute of Information Science and Engineering, Central South University, Changsha 410083, PR China

A R T I C L E D A T A

⁎ Corresponding author. Institute of Metal ReE-mail address: llhe@imr.ac.cn (L.L. He).

1044-5803/$ – see front matter © 2010 Publisdoi:10.1016/j.matchar.2010.06.009

A B S T R A C T

Article history:Received 29 January 2010Received in revised form13 April 2010Accepted 10 June 2010

Deformation-induced α2→γ phase transformation in high Nb containing TiAl alloys wasinvestigated using high-resolution transmission electron microscopy (HREM) and energydispersive X-ray spectroscopy (EDS). The dislocations appearing at the tip of deformation-induced γ plate (DI-γ) and the stacking sequence change of the α2 matrix were two keyevidences for determining the occurrence of the deformation-induced α2→γ phasetransformation. Compositional analysis revealed that the product phase of the room-temperature transformation was not standard γ phase; on the contrary, the product phaseof the high-temperature transformation was standard γ phase.

© 2010 Published by Elsevier Inc.

Keywords:Titanium aluminidesPhase transformationsTransmission electron microscopyEDS

1. Introduction

Titanium aluminides are one of the promising candidates forthe next generation of high-temperature structural materialsdue to lowdensity and goodhigh-temperature properties [1–3].Since their potential service conditions are usually under hightemperature and high stress, the microstructural stability ofTiAl alloys is extremely important for their safe applications[4,5]. It is well known that TiAl alloys usually are comprised oftwo phases: one is the γ phase with a face-centered tetragonalL10 structure; the other is the α2 phase with a hexagonal DO19

structure. The unit cells and atomic positions of these twophases were described in detail in the literature [6].

Deformation-induced phase transformation is interestingbecause the strong stress concentration during deformationplays an important role in the occurrence of this kind of phasetransformation [7–11]. In recent years, the deformation-induced

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phase transformations inTiAl alloyshave been reported severaltimes during plastic deformation, some even in room-temper-ature deformation process [12–22]. Thanks for extensive HREMinvestigations, the mechanisms of structural transition duringthe γ↔HCP (α and α2) phase transformation come to obvious:a/6[112

P] Shockley partial dislocations (including a/6[112

P], a/6

[1P12

P], a/6[11

P2P] dislocations) sweep on alternate {111}γ planes to

fulfill the conversion of the stacking sequence from ABCABCtype to ABAB type in the process of γ→HCP (α and α2) phasetransformation. In the reverse transformation, a/3<11

P00>

Shockley dislocations sweep on successive (0002)α2planes

which change the stacking sequence from ABAB type toABCABC type. However, there is still controversy as to whetherthere is compositional diffusion or not during the deformation-induced γ↔HCP (α and α2) phase transformations. Because thedeformation-induced γ↔HCP phase transformations normallyoccur upon thermal treatments, a major difficulty in this study

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1030 M A T E R I A L S C H A R A C T E R I Z A T I O N 6 1 ( 2 0 1 0 ) 1 0 2 9 – 1 0 3 4

is how to distinguish the deformation-induced phase transfor-mations from the thermally induced phase transformations.Previous studies lacked good methods to distinguish these twokinds of phase transformations.

In the present work, the deformation-induced α2→γ phasetransformation in high Nb containing TiAl alloys has beeninvestigated using high-resolution transmission electronmicroscopy (HREM) and energy dispersive X-ray spectroscopy(EDS). The methods for distinguishing the deformation-induced phase transformations from the thermally inducedphase transformations were discussed.

Fig. 1 – Thenearly-lamellarmicrostructures of the alloys priorto plastic deformation. (a) Ti–45Al–8Nb (at.%);(b) Ti–45Al–8Nb–0.2W–0.2B–0.02Y (at.%).

2. Experimental

The alloy with a nominal composition of Ti–45Al–8Nb (at.%)was melted from pure metals (>99.9%) in an arc furnace withmagnetic agitation under argon protection. Button ingots of∼45 g were obtained. The ingots were re-melted three times toensure completemixing of the constituents. In order to reducethe composition segregation, annealing at 1200 °C for 24 h wasconducted. Cylindrical samples for compression test, withdimensions Ф10×10 mm, were spark-machined from theingots. Room temperature compression tests were performedto a nominal strain of ∼10% with a strain rate ∼10−3/s.

The alloy with nominal composition of Ti–45Al–8Nb–0.2W–0.2B–0.02Y (at.%) was prepared using consumable electrode arcmelting technique under argon protection, and re-melted in avacuum-skull melting furnace in order to reduce compositionheterogeneity. An ingotwith a diameter of 120mmand a heightof 420 mm was cast and then hot isostatic pressed (HIP) toeliminate castingporosity.Chemical analyses indicated that thecomposition of the ingot matched the nominal compositionwithin the experimental error. The forging billet was machinedfrom the ingot, canned and forged at α2+γ phase region to aheight reduction of more than 75%.

For the TEM observation, slices about 0.4 mm in thicknesswere cut from the samples after deformation. The slices weremechanically ground from both sides to ∼30 μm and thenthinned by standard twin jet polishing or ion milling. The thinfoils were examined on a JEOL-2010 high-resolution transmis-sion electron microscope with a point resolution of 0.19 nm.Compositional analysis was carried out on a Tecnai-F30 G2

microscopewhich can provide a highly coherent electron beamwith the minimum size of ∼2 nm for compositional analysis.

3. Experimental Results

Fig. 1 shows the nearly-lamellar microstructures of the twoalloys prior to plastic deformation. Obviously, the microstruc-tures of these two alloys are similar, which are comprised ofprevailing γ/α2 lamellar colonies and a small quantity ofequiaxed γ grains. The deformation-induced α2→γ transfor-mation observed in the present study occurred within the α2

phase of the γ/α2 lamellar colonies. Thus, it was inherentlydifficult to distinguish the deformation-induced transforma-tion product (DI-γ phase) from the pre-existing one (standardγ phase). However, detailed HREM investigations and compo-sitional analyses of the precipitated γ phase will find various

evidences to distinguish the DI-γ phase from the pre-existingγ phase.

3.1. Deformation-induced α2→γ Phase Transformationduring Room-temperature Deformation

Fig. 2 shows a thin γ plate precipitated froma piece of α2 lath ofa lamellar colony after room-temperature compression. Asshown in Fig. 2(a), the precipitated γ plate has only severalatomic planes. The thickness of this precipitated γ plate is toothin to obtain any of its diffraction patterns by using micro-diffraction technique. Fig. 2(b) shows the composite diffrac-tion pattern of the precipitated γ phase and surrounding α2

matrix, which was obtained by Fourier transform of the HREMimage in Fig. 2(a). In order to reveal the differences andrelations between the deformation-induced γ phase (DI-γphase) and the pre-existing γ phase, the composite diffractionpattern of the pre-existing γ and α2 phases prior to deforma-tion is shown in Fig. 2(c) which was also obtained by Fouriertransform of a HREM image. Comparison of these twocomposite diffraction patterns reveals that the deformation-induced γ phase and the pre-existing γ phase have similardiffraction patterns, suggesting that they have similar crystalstructures. However, there indeed exist distinctions betweenthe deformation-induced γphase and the pre-existing γphase. As indicated in Fig. 2(b), the line between the 11

P1γ

and 02P0γ diffraction spots of the deformation-induced γ phase

does not pass through the 22P00α2

diffraction spot of the α2

matrix, which locates on the inner side of the 22P00α2

diffraction; on the contrary, as shown in Fig. 2(c), the line

Fig. 2 – The deformation-induced α2→γ phase transformation in the α2 lath during room temperature compression. Theincident beamparallel to the axis [1120

_]α2 | | [1

_01]γ. (a) HREM image showing the formation of DI-γ plate in a pre-existingα2 lath.

(b) The composite diffraction pattern of the DI-γ plate and the α2 matrix obtained by Fourier transform. (c) The compositediffraction pattern of the pre-existing γ and α2 phases obtained by Fourier transform.

Table 1 – Chemical composition of the DI-γ phase and thesurrounding α2 matrix.

Al (at.%) Ti (at.%) Nb (at.%) (Ti+Nb)/Al

DI-γ phase 39.7±0.4 50.9±0.6 9.4±0.7 1.52Surrounding α2

phase39.1±0.5 53.1±0.8 7.8±0.9 1.56

The symbol±means uncertainty.

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between the 11P1γ and 02

P0γ diffraction spots of the pre-existing γ

phase does pass through the 22P00α2

diffraction spot of the α2

matrix. This fine but important distinction between theirdiffraction patterns reveals that the deformation-induced γphase has a slightly larger crystal cell than the pre-existing γphase. According to the diffraction patterns in Fig. 2, the latticeparameters of the DI-γ phase formed during room-temperaturedeformation are as follows: a=0.402 nm, c=0.408 nm.

In order to determine the compositional change during thedeformation-induced α2→γ phase transformation, composi-tional analysis using EDS was performed in the precipitated γphase and the surrounding α2 matrix. The results are listed inTable 1. It is clear that the precipitated γ phase has the samecomposition as the α2 matrix. There is no compositionaldiffusion in the process of the deformation-induced α2→γphase transformation during room-temperature deformation.Considering the fact that Nb atoms tend to replace Ti atoms[23–25], the ratios (Ti+Nb)/Al in both phases are very close to1.5. In the case of thermally induced α2→γ phase transforma-tions, compositional diffusion must take place and thecompositional difference between the α2 matrix and the γproduct phase will be obvious and detectable. Therefore, thedifferences on structures (indicated by diffraction pattern) andcomposition between the present precipitated γ phase and thepre-existing γ phase suggested that the precipitated γ phase

shown in Fig. 2(a) was the product of deformation-inducedphase transition.

3.2. Deformation-induced α2→γ Phase Transformationduring High-temperature Deformation

Fig. 3(a) shows three thin γ lamellae (labeled Lath1, Lath2 andLath3, respectively), which apparently precipitated from a pre-existingα2 lathofa lamellar colony.DetailedHREMinvestigationsof the microstructural evolution around these precipitated γlamellae are necessary to determine whether the present α2→γphase transformation is induced by deformation or not.

As shown in Fig. 3(b), the γ plate Lath1 only has four layersof atomic planes. The stacking sequences of the α2 matrix atboth sides of the Lath1 γ plate have changed: the α2 matrix at

Fig. 3 – (a) TEM image showing three pieces of γ thin plates (labeled Lath1, Lath2 and Lath3, respectively) precipitated from thepre-existing α2 phase during high-temperature deformation. The incident beam parallels to the crystal direction[1120

_]α2 | | [1

_01]γ. (b) HREM image of Lath1, (c) HREM image of Lath3, (d) HREM image of the tip of Lath3. (e) The result of EDS linear

scanning of the Lath3 γ plate.

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the underside has the stacking sequence of ABAB type and theα2 matrix at the upside has the stacking sequence of CACAtype. At the same time, a Shockley partial dislocation withBurger vector of a/3[11

P00]α2

can be found at the tip of the Lath1γ plate. The change of stacking sequence of the α2 matrixreveals that the Lath1 γ plate results from the glide ofa= 3½1100�α2

Shockley dislocation on the (0001)α2crystal

plane. That is to say, the Lath1 γ plate is the product of thedeformation-induced α2→γ phase transformation, instead ofthe thermal phase transformation. The a= 3½1100�α2

Shockleydislocation at the tip and the change of stacking sequence ofthe α2 matrix are two direct and convincing evidences for thedeformation-induced α2→γ phase transformation.

Fig. 3(c) shows the HREM image of the precipitated γ plateLath3, which has a thickness of six {111} planes (two layersthicker than the Lath1 and Lath2 γ plates). Similar to the case oftheLath1γplate, the stacking sequencesof theα2matrixat bothsides of the Lath3 γ plate have also changed. The stacking

sequence of the α2 matrix changes from ABAB type at theunderside to BCBC type at the upside. The relationship betweenthe thickness of the precipitated γ plate and the change of thestacking sequence of the α2 matrix will be further discussed inthe Discussion section.

Fig. 3(d) shows the HREM image of the tip of the Lath3 γplate. Though this HREM image is not so well-resolved due tothe existence of strong stress concentration in the tip area, itshows a very important microstructural information of theprecipitated γ phase: A growth ledge of two close-packedplanes (111)γ. This growth ledge is the migration trace of onea= 3½1100�α2

Shockley partial dislocation. The precipitated γplate proceeds by the migration of each Shockley partialdislocation on alternate (0001)α2

crystal plane, which results inthe structural transition of the hcp–fcc phase transformation.

Inorder todetermine thecompositional changesduring thesephase transformations, the experiment of EDS linear scanningacross theDI-γphasewas carried out on the STEMmode. Fig. 3(e)

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shows the compositional scanning result of the Lath3 γ plate, inwhich the composition at the place of 12 nm corresponds to thatof the Lath3 γ plate. It is obvious that the Al content increasedand theTi content decreased in the Lath3γplate comparingwiththose of the α2 matrix. Compositional changes and atomicdiffusions took place in the process of deformation-inducedα2→γ phase transformation under high-temperature deforma-tion. Obviously, the DI-γ phase formed during high-temperaturedeformation has the same structure and composition as the pre-existing γ phase. Therefore, the lattice parameters of the DI-γphase formed during high-temperature deformation are asfollows: a=0.396 nm, c=0.402 nm.

4. Discussion

α2-Ti3Al has an ordered D019 structure which derives from thehcp structure. γ-TiAl has an ordered L10 structure whose c/aratio is∼1.02.Thec/a ratioofγ-TiAl is so close to1 that theα2→γphase transformation can be treated as a transformation fromhcp to fcc [26]. Although the mechanism of the hcp→ fcc phasetransformation is well known, detailed analysis of this mech-anism will help us understand the methods for distinguishingthe DI-γ phase from the pre-existing γ phase.

Fig. 4 shows a schematicmodel for explaining the structuralchange from α2 to γ by the glide of a = 3½1100�α2

Shockley partialdislocationsonalternate (0001)α2 planes.Asshown in this figure,if one a= 3½1100�α2

Shockley partial dislocation sweeps acrossone (0001)α2 plane, a nucleus of γ phase with four atomic planeswill be formed in the α2 matrix; at the same time, the stackingsequenceof theα2matrixwill be changed fromABAB type at theunderside to CACA type at the upside. After that, if anothera= 3½1100�α2

Shockley partial dislocation glides on alternate(0001)α2 plane of the α2 matrix, another two layers of α2 phasewill be converted into the γ phase and the γ nucleus willincrease two atomic planes; at the same time, the stackingsequence of the α2 matrix at the upside will be further changedfrom previous CACA type to BCBC type. The stacking sequencechanges of the α2 matrix, accompanying with the occurrence ofthe α2→γ phase transformation, are convincing evidences toprove this phase transformation to be a deformation-induced

Fig. 4 – Schematic illustration of the deformation-inducedα2→γ phase transformation.

phase transformation because the stacking sequence of the α2

matrix will not change during a thermally induced phasetransformation. Also according to this mechanism, the initialstep of the deformation-induced α2→γ phase transformation isto form a γ nucleus with four layers of atomic planes. This γnucleus can also be considered as a stacking fault of the α2

matrix. It is impossible for the γ nucleus to have three or fivelayers of atomic planes. After the nucleation, the γ phase canonly increase two crystal planes after each migration ofShockley dislocation. Therefore, the four-layer γ nucleus andthe two-layer-height ledge at the tip of γ plate are also twoimportant evidences to prove the present α2→γ phase trans-formation to be adeformation-inducedphase transformation. Itis because the γ nucleus can have random layers of atomicplanes and the γ nucleus can increase random layers of atomicplanes in a thermally induced phase transformation. The α2→γphase transformations observed in this study are consistentwell with all the evidences of the deformation-induced phasetransformation. Therefore, the α2→γ phase transformations inthis study are induced by deformation, instead of thermaleffects.

According to the experimental results, it can be concludedthat the structural transition and the compositional changeare two aspects of the deformation-induced α2→γ phasetransformation. These two aspects are not only related butalso independent of each other. On one hand, the occurrenceof a true deformation-induced α2→γ phase transformationrequires both the structural transition and the compositionalchange, just as those observed in the high-temperature phasetransformation; the products of this kind of phase transfor-mations are standard γ phase. On the other hand, thestructural transition and the compositional change can takeplace independently. The structural transition can take placewithout any compositional change, just as those observed inthe room-temperature phase transformation. The products ofthis kind of phase transformations are not standard γ phase.

5. Conclusions

Deformation-induced α2→γ phase transformation took placeduring room-temperature and high-temperature deforma-tions of high Nb containing TiAl alloys. The dislocationsappearing at the tip of deformation-induced γ plate (DI-γ) andthe stacking sequence change of the α2 matrix were two keyevidences for determining the occurrence of the deformation-induced α2→γ phase transformation. There was no composi-tional diffusion during the room-temperature transformationand the product phase was not standard γ phase; on thecontrary, the compositional diffusion took place during thehigh-temperature transformation and the product phase wasstandard γ phase.

Acknowledgments

This research was carried out with supports from NSFC grantNo. 50271073 and theMajor State Basic Research Projections ofChina grant No. 2006CB600905, which the authors gratefullyacknowledge.

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