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Intermetallics 19 (2011) 93e98

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Intermetallics

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

Spinodal decomposition of the g-phase upon quenching in the TieAleNb ternaryalloy system

Orlando Rios, Fereshteh Ebrahimi*

University of Florida, Materials Science and Engineering, PO Box 116400, Gainesville, FL 32611, USA

a r t i c l e i n f o

Article history:Received 27 July 2010Received in revised form15 September 2010Accepted 22 September 2010

Keywords:A: Titanium aluminides, based on TiAlB: Phase transformationB: Phase identificationF: Electron microscopy, transmission

* Corresponding author. Tel.: þ1 352 846 3791; faxE-mail address: febra@mse.ufl.edu (F. Ebrahimi).

0966-9795/$ e see front matter � 2010 Elsevier Ltd.doi:10.1016/j.intermet.2010.09.014

a b s t r a c t

The g-TiAl with L10 crystal structure shows extensive solubility for Nb at elevated temperatures. Recently(Rios et al., Acta materialia 2009; 57:6243), we have demonstrated that the high-Nb g-TiAl phasebecomes unstable upon rapid cooling into a nano-scale two-phase microstructure. In this paper, usingdetailed compositional and microstructural analyses, we have demonstrated that this phase goesthrough a spinodal decomposition that results in the compositionally distinct phases identified asa lower-Nb g-phase and the h-phase, which is rich in Nb and forms by the ordering of this element in theg-phase.

� 2010 Elsevier Ltd. All rights reserved.

1. Introduction

g-TiAl-based alloys exhibit promising properties as turbine bladematerials in aerospace jet engines [1e5]. More recently, near-g alloyshave been implemented as low pressure turbine blade materials inthe new generation of GE engines [6]. The addition of Nb has beenshown to enhance the mechanical properties of the g-TiAl phase[7e9] and to stabilize the solid solution b-phase for improved forge-ability [10e12]. Furthermore, it has been demonstrated thatTieAleNb alloys that have a g-TiAlþ s-Nb2Al microstructure exhibitexcellent high-temperature mechanical properties [13,14].

Recently, we have investigated the TieAleNb system theoreti-cally [15,16] as well as experimentally [17e21]. Experimentalexaminations of the invariant reaction involving the L, g-TiAl, s-Nb2Al, h-Al3Ti phases [18] revealed a ternary eutectic reaction,consistent with the theoretical optimization [15]. A series of heat-treating experiments were conducted on two different alloys toestablish the tie-triangles between the g-TiAl, s-Nb2Al, and h-Al3Tiphases equilibrated at 1510 and 1410 �C [18]. The compositions ofthese alloys and the s, g and h phases in equilibrium at the heat-treatment temperatures are listed in Table 1. Microstructural evalu-ations indicated that the high Nb g-phase in equilibrium at 1510 �Cwas not stable upon quenching to room temperature, however

: þ1 352 846 3355.

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quenching from 1410 �C did not cause any phase transformation inthe g-phase. Literature reviews did not disclose any reportedstructural transformations in ternary high-Nb g-phase alloys. In theTieAl binary system, however, several references reported twometastable structures that form upon quenching of the off-stoichi-ometry TiAl, namely the Al2Ti (h-phase) and the Al5Ti3 phase[22e24]. Both of these phases are superstructures that form by thereordering of Al atoms to form layers on particular lattice sites of theTiAl L10 crystal structure.

The present study focuses on the transformational changes inthe high Nb g-phase when the formation of equilibrium phases iskinetically inhibited. Alloy A2 (Table 1) was selected for furtherexamination as it contained the greater volume fraction of the g-phase. High resolution analytical transmission electron microscopyas well as detailed analysis of the electron diffraction patterns wasemployed to characterize this transformation.

2. Experimental procedures

Alloy A2 was produced by non-consumable-arc melting (tung-sten electrode) of high purity starting components (Ti, Al, Nb) usinga water-cooled copper hearth in a gettered ultra high purity argonatmosphere. The buttons were melted, turned over and remelted 6times in an attempt to insure homogeneity of the alloys.

Alloy A2 was subjected to 1510 �C and 1410 �C heat-treatmentsand then water quenched in an attempt to preserve the hightemperature equilibrium microstructure. The as-cast button was

Table 1Compositional analysis of the heat-treated bulk materials as well as the composition of the individual phases as evaluated by EPMA.

Bulk (at%) s (at%) h (at%) g (at%)

Alloy-HT Al Ti Nb Al Ti Nb Al Ti Nb Al Ti Nb

A1-1410 �C 57.2 7.1 35.7 42.6 9.5 47.9 69.9 4.5 25.6 54.3 11.5 34.1A2-1410 �C 51.4 8.8 39.8 41.7 9.1 49.2 69.7 3.7 26.6 54.6 12.1 33.3A1-1510 �C 57.4 6.7 35.9 42.2 7.5 50.3 69.9 3.2 26.9 47.8 12.2 40A2-1510 �C 51.3 8.4 40.3 42.3 7.4 50.3 70.0 3.0 27.0 47.3 11.3 41.4

O. Rios, F. Ebrahimi / Intermetallics 19 (2011) 93e9894

sectioned into 1.5 mm thick slices, which were wrapped withtantalum foil. A vertical alumina tube furnace equipped with dropquenching capability was used for heat treating. The samples weresuspended within the hot zone by a Ta wire and the alumina tubewas sealed by end caps. The tube furnace was purged three timeswith oxygen gettered high purity Ar and heat treatment was con-ducted under a flowing Ar gas. The samples were heated to 1510 �Cand 1410 �C at 10 K/min and held isothermally for 4 h at thetemperature, followed by drop quenching into a water bath. Directimmersion water quenching was performed by removing thebottom cap and nearly simultaneously dropping the sample intothe water bath. The samples were released from the upper cap bya custom built rotating hook. Following the water quench, thesamples were ground and polished to remove any environmentallyaffected zones that may have formed during heat treatment.

The microstructure of heat-treated samples was studied usingthe BSE (back scattered electron) mode in SEM (scanning electronmicroscope). The large difference in the atomic number of thealloying elements provided sufficient z-contrast between phasesthus facilitating microstructural analyses. Conventional TEM(transmission electron microscopy) was conducted on a JEOL200CX, whereas STEM (scanning TEM) and EDS (energy dispersivespectroscopy) were performed on a JEOL 2010F. High Angle AnnularDark Field imaging (dark field STEM) was employed to facilitate thez-contrast based imaging by the acquisition of only the incoher-ently scattered electrons that are highly sensitive to the atomicmass. The site-specific TEM foils were prepared by focused ionbeam (FIB) sectioning and thinning using a FEI Strata DB 235instrument.

Fig. 1. (a) and (b) present SEM/BSE micrographs showing the formation of the g-TiAl,s-Nb2Al and h-Nb3Al phases in the samples heat treated at 1510 �C and 1410 �C,respectively.

3. Microstructural evaluation

Fig. 1a,b present the microstructure of samples quenched from1510 �C to 1410 �C, respectively. Obviously, the microstructureequilibriated at 1510 �C is coarser than the one formed at 1410 �C.The interesting difference between these two microstructures isthe appearance of the fine two-phase microstructure within the g-phase region in the sample quenched from 1510 �C (Fig. 1a). Closeobservation of the prior g-phase boundaries in this sample revealeda dark contrast phase that seemed to grow into the g-phase from itsinterface with the s-phase. However, this interfacial phase was notobserved at the g/h interface. It should be noted that the XRDanalysis did not indicate the presence of another phase in additionto the g, h and s phases in this sample [18].

To further investigate the structure and composition of themicro-constituent phases within the transformed g-phase region,TEM analysis was performed. A 15 mm wide site-specific TEM foilwas prepared from the sample heat treated at 1510 �C as indicatedby the marker shown on Fig. 1a. The section was selected to consistmostly of the transformed g-phase region while cutting throughthe s/g interface and to include a small region of the s-phase inorder to evaluate the dark-contrast boundary phase.

A typical bright field/dark field STEM image set is shown inFig. 2. Consistent with the BSE/SEM image (Fig. 1a) two phases with

bright and dark contrasts were found within the prior g-phase. Thetwo distinct phase contrasts confirms a compositional differencebetween these two phases and suggest a diffusion assisted solid-state phase transformation. The dark-contrast phase in the brightfield was located at the s/g interface as well as throughout thetransformed g-phase (markedwith “1” in Fig. 2). Dark field imagingof this phase revealed that it is the phase with the higher atomicnumber (Fig. 2b). The bright contrast phase was seen only in theinterior of the transformed g-phase region (marked with “2” inFig. 2). The phase adjacent to the two-phase region was identifiedas the s-phase. The s-phase produced a uniform contrastthroughout the region indicating little to no compositional varia-tion in this phase.

Fig. 2. (a) Bright field and (b) dark field STEM images demonstrating the compositional difference between the two phases within the transformed g-TiAl region.

Fig. 3. A depiction of the compositions of the two-phase region that forms uponquenching from 1510 �C demonstrating their relationship to the 1410 �C equilibriumtie-triangle.

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4. Compositional analysis

The STEM imaging provided supporting evidence that the twophases with different contrasts were compositionally distinct. TEM/EDS analysis was used to resolve the compositions of the individualmicro-constituents. A bright field STEM image of one of the regionsexamined by TEM/EDS is shown in Fig. 3. The EDS analysis alsorevealed two statistically distinct compositions. The composition ofregion “2” was determined to be 31.2Nb 57.2Al 11.6Ti (at%) whilethe region “1” contained 45.2Nb 43.9Al 10.9Ti (at%).

The compositions corresponding to regions 1 and 2 are markedona ternaryplot,which also shows the tie-triangles consistingofs,handgphases at 1510 �C and1410 �C (Fig. 3) thatweremeasured fromboth alloys A1 and A2 as listed in Table 1. The compositional analysiselucidates that the bright phase in Fig. 3 (region “2”) rejects suffi-cient Nb for Al such that its composition is brought near to theequilibrium composition of the g-phase at 1410 �C, whereas thelower-z phase marked as “1” is compositionally located near the s-phase corner of the 1410 �C tie-triangle. Prior work within ourresearch group has shown that the nucleation of the s-phase iskinetically difficult therefore it is unlikely that it nucleated out of theg-phase upon quenching [10,15]. Thus the dark-contrast phase wassuspected to be a metastable phase that formed upon quenching.

5. Structural analysis

A bright field TEM image of the transformed g-phase is shown inFig. 4, inwhich the two phases are marked as “a” and “b”. Structuralanalysis was performed by recording SAD (selected area diffraction)patterns of phases individually as well as from the two adjacentphases simultaneously. The analysis of the diffraction patternsrevealed that the two phases are structurally distinct. The phase “a”was identified through the analysis of the diffraction patterns to bethe g-phase. A diffraction pattern near the [110]g zone axis is shownin Fig. 4b. The second phase was determined not to be the g-phase.

Inspection of the diffraction patterns of the second phase near itszone axes revealed that the lattice parameter of this phase is close totriple that of the g-phase (Fig. 4c). Amongmany phases considered,as mentioned previously, there is evidence in the literature ofa metastable orthorhombic h-Al2Ti phase that forms in the binaryTieAl system upon quenching the g-phase [22e24]. Morphologicalsimilarities were found between the h-phasewithin the TiAl matrixand the g-phase transformation observed in this study.

Fig. 4. (a) Bright field TEM image of the two-phase region. (b) The diffraction pattern (DP) identifying the “a” phase as the g-phase. (c) The DP identifying the “b”phase as the h-phase.

Fig. 5. Crystal structures of the g-phase and the h-phase.

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A structural model was generated using the Crystal Makersoftware package from the reported structural data for both the g-TiAl [25,26] and the h-Al2Ti phases [22] and are reproduced inFig. 5. The h-Al2Ti phase is a superlattice that forms by ordering ofthe Al atoms on the (001) planes of the g-TiAl L10 structure, thus itslattice parameter is inherent to that of parent g-phase. Essentially

Fig. 6. (a) DPs of the g-phase (near [110] zone) and h-phase (near [001] zone) demonstratiplanes. (b) A comparison of the simulated diffraction patterns with the actual ones. (c) The

the h-phase contains two lattice parameters that are approximatelyequal to those of the g-phase and the third parameter is about threetimes of the L10’s lattice parameter. Investigation of the diffractionpatterns confirmed that the “b” phase in Fig. 4a is the h-phase.

A SAD pattern that covers both phases in the prior g-phase isshown in Fig. 6a. This pattern is taken near the [110]g zone axis,which is also slightly off the [001]h zone axis. The combineddiffraction patterns elucidates that the spacing between the (030)hplanes is larger than the spacing between the (001)g planes andthere is a 35� rotation between these two planes. On the other handthe spacing of the (030)h planes is very close to the spacing of the{110)g planes. The combined recorded and simulated patterns areshown in Fig. 6b, which shows an excellent match between thestructures.

The foil was tilted slightly in order to reach a two beam condi-tion off the [110]g zone axis. This diffraction pattern was recordedand is shown in Fig. 6c. The ð110Þg and the ð110Þh diffracted beamsshown on this figurewere used for dark field imaging of each phase.The bright field and dark field images are shown in Fig. 7. The darkfield image presented in Fig. 7a was obtained using the ð110Þgreflection and identified the “a” phase as the g-phase. The dark fieldpresented in Fig. 7b was generated using the ð110Þh diffractedbeam, which identified the “b” phase as the h-phase. This analysiswas combined with the STEM imaging and compositional analysespreviously discussed in order to correlate the phase identification

ng that [110]g//[001]h with a rotation of approximately 35� between (001)g and (030)hdiffraction spots used for the bright field/dark field images shown in Fig. 7.

Fig. 7. BF-TEM image of the two-phase region with the associated dark fields using the ð110Þg and the ð110Þh diffracted beams, respectively.

O. Rios, F. Ebrahimi / Intermetallics 19 (2011) 93e98 97

with the measured composition of each phase. The g-phase shownin Fig. 7 was determined to be the dark phase in the bright fieldSTEM image that is marked with “1” in Figs. 2 and 3. The compo-sition of the g-phase is then identified as 31.2Nb 57.2Al 11.6Ti (at%).The darker-contrast phase in STEM imagemarked as “2” is found tobe the h-phase that has a composition of 45.2Nb 43.9Al 10.9Ti (at%).The composition of this phase suggests that approximately 25% ofAl sites and 75% of the Ti sites in the Al2Ti structure are occupied byNb atoms. Since no extra superlattice spots in addition to thediffraction spots from the h-phase were found, it is suggested thatthe Nb atoms are distributed randomly.

The morphology of the g and h phases and their distinct compo-sitional differences suggests that the high temperature g-phase goesthrough a spinodal decomposition upon fast cooling. The significantchange in the solubility of Nb in the g-phase upon cooling is a keycomponent driving this transformation. An abrupt increase in thechemical potential gradient, driving the diffusion of Nb out of the g-phase and into the s-phase is anticipated. However; the long rangediffusion is kinetically limitedbyquenching. Apparently, the enthalpyassociated with the mixing of Ti, Al and Nb in the g-phase results inamiscibility gap in theGibbs free energy surface during the cooling ofthis thermodynamically unstable phase. Localized short range diffu-sion driven by the negative curvature in the free energy surface leadsto partitioning of Nb and Al within the g-phase. Therefore, thesubsequent short range diffusion renders the h-phase formationthrough a spinodal decomposition of the g-phase. The combinedaction of the chemical driving forces and limited diffusion drive thenano scaled wormy appearing spinodal microstructure and theformation of theg-phase superstructure (h-phase). The characteristicspacing between the phases of under 200 nm scale suggests thatnucleation was not the rate limiting mechanism. This spacing is alsoconsistent with a slightly coarsened spinodal structure [27,28].

6. Summary and conclusions

Quenching an alloy with nominal composition of 51.5Al 8.5Ti40Nb at% (alloy A2) from 1510 �C caused a solid-state trans-formation in the g-phase, whereas a 1410 �C heat-treatment andquenching of the same alloy resulted in no detectable structuralchanges in the g-phase. Prior studies disclosed significant heatevolution or absorption between these two temperatures indi-cating a compositional and volume fraction change [18]. TEM/STEManalyses revealed that during the transformation upon quenchingfrom 1510 �C two phases identified as g and h phases evolve. Thecomposition the g-phase obtained upon quenching was 31.2Nb57.2Al 11.6Ti at% and it was close to the equilibrium composition ofthis phase found in the microstructure equilibrated at 1410 �C. Thisis shown in Fig. 3b where now region “1” has been identified as theg-phase.

The h-phase, which showed a dark contrast in the STEM image,exhibited a composition of 45.2Nb 43.9Al 10.9Ti at%. Apparently,this phase has a higher solubility for Nb than the g-phase doessuggesting that Nb is preferentially located in Al lattice positions.Observation of the interface between the s-phase and the trans-formed g-phase region revealed that the h-phase covered theboundary (Fig. 2). Compositional analysis revealed that the h-phaseis a Nb rich phase therefore it is logical that it is found adjacent tothe high Nb s-phase. In contrast near the Al rich h-phase bound-aries the h-phase was not observed (Fig. 1). This further confirmsthat this phase forms by the diffusion of Al and Nb.

These results suggest that at high temperatures the g-phase hasa high solubility for Nb, which decreases drastically with reducingtemperature. When the formation of s and h phases is inhibited bywater quenching, the metastable g-phase releases its Nb byundergoing a spinodal decomposition.

O. Rios, F. Ebrahimi / Intermetallics 19 (2011) 93e9898

Acknowledgements

This research was supported by NSF/AFOSR under Grant No.DMR-0605702 and DMR-0856622.

References

[1] Martin PL, Mendiratta MG, Lipsitt HA. Metallurgical Transactions A- PhysicalMetallurgy and Materials Science 1983;14:2170.

[2] Hayes RW, London B. Acta Metallurgica et Materialia 1992;40:2167.[3] Apple F, Wagner R. Materials Science and Engineering R-Report 1998;22:187.[4] Kim Y-W. Intermetallics 1998;6:623.[5] Wallgram W, Schomtzer T, Cha LM, Das G, Guther V, Clemens H. International

Journal of Materials Research 2009;100:1021.[6] G.E. Aviation, http://www.geae.com/education/theater/genx/.[7] Kandra JT, Lee EW. Metallurgical and Materials Transactions A-Physical

Metallurgy and Materials Science 1994;25:1667.[8] Tetsui T. Intermetallics 2002;10:239.[9] Xu XJ, Xu LH, Lin JP, Wang YL, Lin Z, Chen GL. Intermetallics 2005;13:337.

[10] Masahashi N, Mizuhara Y, Matsuo M, Hanamura T, Kimura M, Hashimoto K.Isij International 1991;31:728.

[11] Kawabata T, Fukai H, Izumi O. Acta Materialia 1998;46:2185.

[12] Clemens H, Chaldif HF, Wallgram W, Zickler GA, Gerling R, Liss KD, et al.Intermetallics 2008;16:827.

[13] Ebrahimi F, Castillo-Gomez JR. Acta Metallurgica et Materialia 1991;25:1737.[14] DeAragao BJG, Ebrahimi F. Materials Science and Engineering A-Structural

Materials Properties Microstructure and Processing 1996;208:37.[15] D.M. Cupid, Thermodynamic assessment of the TiAlNb, TiAlCr, and TiAlMo

systems. Ph.D. thesis, University of Florida; 2009.[16] Cupid DM, Fabrichnaya O, Rios O, Ebrahimi F, Seifert HJ. International Journal

of Materials Research 2009;100:218.[17] O. Rios. Examination of the high temperature equilibrium in the TiAlNb

ternary alloy systems. PhD thesis, University of Florida; 2009.[18] Rios O, Cupid DM, Seifert HJ, Ebrahimi F. Acta Materialia 2009;57:6243.[19] Rios O, Goyel S, Kesler MS, Cupid DM, Seifert HJ, Ebrahimi F. Scripta Materialia

2009;60:156.[20] Kesler MS, Goyel S, Rios O, Cupid DM, Seifert HJ, Ebrahimi F. Materials Science

and Engineering A-Structural Materials Properties Microstructure and Pro-cessing 2010;527:2857.

[21] Goyel S, Rios O, Kesler MS, Ebrahimi F. Intermetallics 2010;18:1491.[22] Nakano T, Negishi A, Hayashi K, Umakoshi Y. Acta Materialia 1999;47:1091.[23] Stein F, Zhang LC, Sauthoff G, Palm M. Acta Materialia 2001;49:2919.[24] Hayashi K, Nakano T, Umakoshi Y. Intermetallics 2002;10:771.[25] Konitzer DG, Jones IP, Fraser HL. Scripta Metallurgica 1986;20:265.[26] Jackson AG. Scripta Metallurgica et Materialia 1993;28:673.[27] Cahn JW. Acta Metallurgica 1961;9:795.[28] Glotzer SC, Dimarzio EA, MuthukumarM. Physical Review Letters 1995;74:2034.