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Spinodal decomposition of the γ-phase upon quenching in the Ti–Al–Nb ternary alloy system
Intermetallics 19 (2011) 93e98
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Spinodal decomposition of the g-phase upon quenching in the TieAleNb ternary alloy 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
a b s t r a c t
Article history: Received 27 July 2010 Received in revised form 15 September 2010 Accepted 22 September 2010
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 phase becomes unstable upon rapid cooling into a nano-scale two-phase microstructure. In this paper, using detailed compositional and microstructural analyses, we have demonstrated that this phase goes through a spinodal decomposition that results in the compositionally distinct phases identified as a lower-Nb g-phase and the h-phase, which is rich in Nb and forms by the ordering of this element in the g-phase. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A: Titanium aluminides, based on TiAl B: Phase transformation B: Phase identification F: Electron microscopy, transmission
1. Introduction
g-TiAl-based alloys exhibit promising properties as turbine blade materials in aerospace jet engines [1e5]. More recently, near-g alloys have been implemented as low pressure turbine blade materials in the new generation of GE engines [6]. The addition of Nb has been shown to enhance the mechanical properties of the g-TiAl phase [7e9] and to stabilize the solid solution b-phase for improved forgeability [10e12]. Furthermore, it has been demonstrated that TieAleNb alloys that have a g-TiAl þ s-Nb2Al microstructure exhibit excellent high-temperature mechanical properties [13,14]. Recently, we have investigated the TieAleNb system theoretically [15,16] as well as experimentally [17e21]. Experimental examinations of the invariant reaction involving the L, g-TiAl, sNb2Al, h-Al3Ti phases [18] revealed a ternary eutectic reaction, consistent with the theoretical optimization [15]. A series of heattreating experiments were conducted on two different alloys to establish the tie-triangles between the g-TiAl, s-Nb2Al, and h-Al3Ti phases equilibrated at 1510 and 1410 C [18]. The compositions of these alloys and the s, g and h phases in equilibrium at the heattreatment temperatures are listed in Table 1. Microstructural evaluations indicated that the high Nb g-phase in equilibrium at 1510 C was not stable upon quenching to room temperature, however
* Corresponding author. Tel.: þ1 352 846 3791; fax: þ1 352 846 3355. E-mail address: [email protected]fl.edu (F. Ebrahimi). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.09.014
quenching from 1410 C did not cause any phase transformation in the g-phase. Literature reviews did not disclose any reported structural transformations in ternary high-Nb g-phase alloys. In the TieAl binary system, however, several references reported two metastable structures that form upon quenching of the off-stoichiometry TiAl, namely the Al2Ti (h-phase) and the Al5Ti3 phase [22e24]. Both of these phases are superstructures that form by the reordering of Al atoms to form layers on particular lattice sites of the TiAl L10 crystal structure. The present study focuses on the transformational changes in the high Nb g-phase when the formation of equilibrium phases is kinetically inhibited. Alloy A2 (Table 1) was selected for further examination as it contained the greater volume fraction of the gphase. High resolution analytical transmission electron microscopy as well as detailed analysis of the electron diffraction patterns was employed to characterize this transformation. 2. Experimental procedures Alloy A2 was produced by non-consumable-arc melting (tungsten electrode) of high purity starting components (Ti, Al, Nb) using a water-cooled copper hearth in a gettered ultra high purity argon atmosphere. The buttons were melted, turned over and remelted 6 times in an attempt to insure homogeneity of the alloys. Alloy A2 was subjected to 1510 C and 1410 C heat-treatments and then water quenched in an attempt to preserve the high temperature equilibrium microstructure. The as-cast button was
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Table 1 Compositional analysis of the heat-treated bulk materials as well as the composition of the individual phases as evaluated by EPMA.
s (at%)
Bulk (at%) Alloy-HT A1-1410 A2-1410 A1-1510 A2-1510
C C C C
h (at%)
g (at%)
Al
Ti
Nb
Al
Ti
Nb
Al
Ti
Nb
Al
Ti
Nb
57.2 51.4 57.4 51.3
7.1 8.8 6.7 8.4
35.7 39.8 35.9 40.3
42.6 41.7 42.2 42.3
9.5 9.1 7.5 7.4
47.9 49.2 50.3 50.3
69.9 69.7 69.9 70.0
4.5 3.7 3.2 3.0
25.6 26.6 26.9 27.0
54.3 54.6 47.8 47.3
11.5 12.1 12.2 11.3
34.1 33.3 40 41.4
sectioned into 1.5 mm thick slices, which were wrapped with tantalum foil. A vertical alumina tube furnace equipped with drop quenching capability was used for heat treating. The samples were suspended within the hot zone by a Ta wire and the alumina tube was sealed by end caps. The tube furnace was purged three times with oxygen gettered high purity Ar and heat treatment was conducted under a flowing Ar gas. The samples were heated to 1510 C and 1410 C at 10 K/min and held isothermally for 4 h at the temperature, followed by drop quenching into a water bath. Direct immersion water quenching was performed by removing the bottom cap and nearly simultaneously dropping the sample into the water bath. The samples were released from the upper cap by a custom built rotating hook. Following the water quench, the samples were ground and polished to remove any environmentally affected zones that may have formed during heat treatment. The microstructure of heat-treated samples was studied using the BSE (back scattered electron) mode in SEM (scanning electron microscope). The large difference in the atomic number of the alloying elements provided sufficient z-contrast between phases thus facilitating microstructural analyses. Conventional TEM (transmission electron microscopy) was conducted on a JEOL 200CX, whereas STEM (scanning TEM) and EDS (energy dispersive spectroscopy) were performed on a JEOL 2010F. High Angle Annular Dark Field imaging (dark field STEM) was employed to facilitate the z-contrast based imaging by the acquisition of only the incoherently scattered electrons that are highly sensitive to the atomic mass. The site-specific TEM foils were prepared by focused ion beam (FIB) sectioning and thinning using a FEI Strata DB 235 instrument.
bright and dark contrasts were found within the prior g-phase. The two distinct phase contrasts confirms a compositional difference between these two phases and suggest a diffusion assisted solidstate phase transformation. The dark-contrast phase in the bright field was located at the s/g interface as well as throughout the transformed g-phase (marked with “1” in Fig. 2). Dark field imaging of this phase revealed that it is the phase with the higher atomic number (Fig. 2b). The bright contrast phase was seen only in the interior of the transformed g-phase region (marked with “2” in Fig. 2). The phase adjacent to the two-phase region was identified as the s-phase. The s-phase produced a uniform contrast throughout the region indicating little to no compositional variation in this phase.
3. Microstructural evaluation Fig. 1a,b present the microstructure of samples quenched from 1510 C to 1410 C, respectively. Obviously, the microstructure equilibriated at 1510 C is coarser than the one formed at 1410 C. The interesting difference between these two microstructures is the appearance of the fine two-phase microstructure within the gphase region in the sample quenched from 1510 C (Fig. 1a). Close observation of the prior g-phase boundaries in this sample revealed a dark contrast phase that seemed to grow into the g-phase from its interface with the s-phase. However, this interfacial phase was not observed at the g/h interface. It should be noted that the XRD analysis did not indicate the presence of another phase in addition to the g, h and s phases in this sample [18]. To further investigate the structure and composition of the micro-constituent phases within the transformed g-phase region, TEM analysis was performed. A 15 mm wide site-specific TEM foil was prepared from the sample heat treated at 1510 C as indicated by the marker shown on Fig. 1a. The section was selected to consist mostly of the transformed g-phase region while cutting through the s/g interface and to include a small region of the s-phase in order to evaluate the dark-contrast boundary phase. A typical bright field/dark field STEM image set is shown in Fig. 2. Consistent with the BSE/SEM image (Fig. 1a) two phases with
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.
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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.
4. Compositional analysis The STEM imaging provided supporting evidence that the two phases with different contrasts were compositionally distinct. TEM/ EDS analysis was used to resolve the compositions of the individual micro-constituents. A bright field STEM image of one of the regions examined by TEM/EDS is shown in Fig. 3. The EDS analysis also revealed two statistically distinct compositions. The composition of region “2” was determined to be 31.2Nb 57.2Al 11.6Ti (at%) while the region “1” contained 45.2Nb 43.9Al 10.9Ti (at%). The compositions corresponding to regions 1 and 2 are marked on a ternary plot, which also shows the tie-triangles consisting of s, h and g phases at 1510 C and 1410 C (Fig. 3) that were measured from both alloys A1 and A2 as listed in Table 1. The compositional analysis elucidates that the bright phase in Fig. 3 (region “2”) rejects sufficient Nb for Al such that its composition is brought near to the equilibrium composition of the g-phase at 1410 C, whereas the lower-z phase marked as “1” is compositionally located near the sphase corner of the 1410 C tie-triangle. Prior work within our research group has shown that the nucleation of the s-phase is kinetically difficult therefore it is unlikely that it nucleated out of the g-phase upon quenching [10,15]. Thus the dark-contrast phase was suspected to be a metastable phase that formed upon quenching. 5. Structural analysis A bright field TEM image of the transformed g-phase is shown in Fig. 4, in which the two phases are marked as “a” and “b”. Structural analysis was performed by recording SAD (selected area diffraction) patterns of phases individually as well as from the two adjacent phases simultaneously. The analysis of the diffraction patterns revealed that the two phases are structurally distinct. The phase “a” was identified through the analysis of the diffraction patterns to be the g-phase. A diffraction pattern near the [110]g zone axis is shown in Fig. 4b. The second phase was determined not to be the g-phase. Inspection of the diffraction patterns of the second phase near its zone axes revealed that the lattice parameter of this phase is close to triple that of the g-phase (Fig. 4c). Among many phases considered, as mentioned previously, there is evidence in the literature of a metastable orthorhombic h-Al2Ti phase that forms in the binary TieAl system upon quenching the g-phase [22e24]. Morphological similarities were found between the h-phase within the TiAl matrix and the g-phase transformation observed in this study.
Fig. 3. A depiction of the compositions of the two-phase region that forms upon quenching from 1510 C demonstrating their relationship to the 1410 C equilibrium tie-triangle.
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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.
A structural model was generated using the Crystal Maker software package from the reported structural data for both the gTiAl [25,26] and the h-Al2Ti phases [22] and are reproduced in Fig. 5. The h-Al2Ti phase is a superlattice that forms by ordering of the Al atoms on the (001) planes of the g-TiAl L10 structure, thus its lattice parameter is inherent to that of parent g-phase. Essentially
the h-phase contains two lattice parameters that are approximately equal to those of the g-phase and the third parameter is about three times of the L10’s lattice parameter. Investigation of the diffraction patterns confirmed that the “b” phase in Fig. 4a is the h-phase. A SAD pattern that covers both phases in the prior g-phase is shown in Fig. 6a. This pattern is taken near the [110]g zone axis, which is also slightly off the [001]h zone axis. The combined diffraction patterns elucidates that the spacing between the (030)h planes is larger than the spacing between the (001)g planes and there is a 35 rotation between these two planes. On the other hand the spacing of the (030)h planes is very close to the spacing of the {110)g planes. The combined recorded and simulated patterns are shown in Fig. 6b, which shows an excellent match between the structures. The foil was tilted slightly in order to reach a two beam condition off the [110]g zone axis. This diffraction pattern was recorded and is shown in Fig. 6c. The ð110Þg and the ð110Þh diffracted beams shown on this figure were used for dark field imaging of each phase. The bright field and dark field images are shown in Fig. 7. The dark field image presented in Fig. 7a was obtained using the ð110Þg reflection and identified the “a” phase as the g-phase. The dark field presented in Fig. 7b was generated using the ð110Þh diffracted beam, which identified the “b” phase as the h-phase. This analysis was combined with the STEM imaging and compositional analyses previously discussed in order to correlate the phase identification
Fig. 6. (a) DPs of the g-phase (near [110] zone) and h-phase (near [001] zone) demonstrating that [110]g//[001]h with a rotation of approximately 35 between (001)g and (030)h planes. (b) A comparison of the simulated diffraction patterns with the actual ones. (c) The diffraction spots used for the bright field/dark field images shown in Fig. 7.
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Fig. 7. BF-TEM image of the two-phase region with the associated dark fields using the ð110Þg and the ð110Þh diffracted beams, respectively.
with the measured composition of each phase. The g-phase shown in Fig. 7 was determined to be the dark phase in the bright field STEM image that is marked with “1” in Figs. 2 and 3. The composition of the g-phase is then identified as 31.2Nb 57.2Al 11.6Ti (at%). The darker-contrast phase in STEM image marked as “2” is found to be the h-phase that has a composition of 45.2Nb 43.9Al 10.9Ti (at%). The composition of this phase suggests that approximately 25% of Al sites and 75% of the Ti sites in the Al2Ti structure are occupied by Nb atoms. Since no extra superlattice spots in addition to the diffraction spots from the h-phase were found, it is suggested that the Nb atoms are distributed randomly. The morphology of the g and h phases and their distinct compositional differences suggests that the high temperature g-phase goes through a spinodal decomposition upon fast cooling. The significant change in the solubility of Nb in the g-phase upon cooling is a key component driving this transformation. An abrupt increase in the chemical potential gradient, driving the diffusion of Nb out of the gphase and into the s-phase is anticipated. However; the long range diffusion is kinetically limited by quenching. Apparently, the enthalpy associated with the mixing of Ti, Al and Nb in the g-phase results in a miscibility gap in the Gibbs free energy surface during the cooling of this thermodynamically unstable phase. Localized short range diffusion driven by the negative curvature in the free energy surface leads to partitioning of Nb and Al within the g-phase. Therefore, the subsequent short range diffusion renders the h-phase formation through a spinodal decomposition of the g-phase. The combined action of the chemical driving forces and limited diffusion drive the nano scaled wormy appearing spinodal microstructure and the formation of the g-phase superstructure (h-phase). The characteristic spacing between the phases of under 200 nm scale suggests that nucleation was not the rate limiting mechanism. This spacing is also consistent with a slightly coarsened spinodal structure [27,28].
6. Summary and conclusions Quenching an alloy with nominal composition of 51.5Al 8.5Ti 40Nb at% (alloy A2) from 1510 C caused a solid-state transformation in the g-phase, whereas a 1410 C heat-treatment and quenching of the same alloy resulted in no detectable structural changes in the g-phase. Prior studies disclosed significant heat evolution or absorption between these two temperatures indicating a compositional and volume fraction change [18]. TEM/STEM analyses revealed that during the transformation upon quenching from 1510 C two phases identified as g and h phases evolve. The composition the g-phase obtained upon quenching was 31.2Nb 57.2Al 11.6Ti at% and it was close to the equilibrium composition of this phase found in the microstructure equilibrated at 1410 C. This is shown in Fig. 3b where now region “1” has been identified as the g-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 does suggesting that Nb is preferentially located in Al lattice positions. Observation of the interface between the s-phase and the transformed g-phase region revealed that the h-phase covered the boundary (Fig. 2). Compositional analysis revealed that the h-phase is a Nb rich phase therefore it is logical that it is found adjacent to the high Nb s-phase. In contrast near the Al rich h-phase boundaries the h-phase was not observed (Fig. 1). This further confirms that this phase forms by the diffusion of Al and Nb. These results suggest that at high temperatures the g-phase has a high solubility for Nb, which decreases drastically with reducing temperature. When the formation of s and h phases is inhibited by water quenching, the metastable g-phase releases its Nb by undergoing a spinodal decomposition.
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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- Physical Metallurgy 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 Processing 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, Muthukumar M. Physical Review Letters 1995;74:2034.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Spinodal decomposition of the g-phase upon quenching in the TieAleNb ternary alloy 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
a b s t r a c t
Article history: Received 27 July 2010 Received in revised form 15 September 2010 Accepted 22 September 2010
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 phase becomes unstable upon rapid cooling into a nano-scale two-phase microstructure. In this paper, using detailed compositional and microstructural analyses, we have demonstrated that this phase goes through a spinodal decomposition that results in the compositionally distinct phases identified as a lower-Nb g-phase and the h-phase, which is rich in Nb and forms by the ordering of this element in the g-phase. Ó 2010 Elsevier Ltd. All rights reserved.
Keywords: A: Titanium aluminides, based on TiAl B: Phase transformation B: Phase identification F: Electron microscopy, transmission
1. Introduction
g-TiAl-based alloys exhibit promising properties as turbine blade materials in aerospace jet engines [1e5]. More recently, near-g alloys have been implemented as low pressure turbine blade materials in the new generation of GE engines [6]. The addition of Nb has been shown to enhance the mechanical properties of the g-TiAl phase [7e9] and to stabilize the solid solution b-phase for improved forgeability [10e12]. Furthermore, it has been demonstrated that TieAleNb alloys that have a g-TiAl þ s-Nb2Al microstructure exhibit excellent high-temperature mechanical properties [13,14]. Recently, we have investigated the TieAleNb system theoretically [15,16] as well as experimentally [17e21]. Experimental examinations of the invariant reaction involving the L, g-TiAl, sNb2Al, h-Al3Ti phases [18] revealed a ternary eutectic reaction, consistent with the theoretical optimization [15]. A series of heattreating experiments were conducted on two different alloys to establish the tie-triangles between the g-TiAl, s-Nb2Al, and h-Al3Ti phases equilibrated at 1510 and 1410 C [18]. The compositions of these alloys and the s, g and h phases in equilibrium at the heattreatment temperatures are listed in Table 1. Microstructural evaluations indicated that the high Nb g-phase in equilibrium at 1510 C was not stable upon quenching to room temperature, however
* Corresponding author. Tel.: þ1 352 846 3791; fax: þ1 352 846 3355. E-mail address: [email protected]fl.edu (F. Ebrahimi). 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.09.014
quenching from 1410 C did not cause any phase transformation in the g-phase. Literature reviews did not disclose any reported structural transformations in ternary high-Nb g-phase alloys. In the TieAl binary system, however, several references reported two metastable structures that form upon quenching of the off-stoichiometry TiAl, namely the Al2Ti (h-phase) and the Al5Ti3 phase [22e24]. Both of these phases are superstructures that form by the reordering of Al atoms to form layers on particular lattice sites of the TiAl L10 crystal structure. The present study focuses on the transformational changes in the high Nb g-phase when the formation of equilibrium phases is kinetically inhibited. Alloy A2 (Table 1) was selected for further examination as it contained the greater volume fraction of the gphase. High resolution analytical transmission electron microscopy as well as detailed analysis of the electron diffraction patterns was employed to characterize this transformation. 2. Experimental procedures Alloy A2 was produced by non-consumable-arc melting (tungsten electrode) of high purity starting components (Ti, Al, Nb) using a water-cooled copper hearth in a gettered ultra high purity argon atmosphere. The buttons were melted, turned over and remelted 6 times in an attempt to insure homogeneity of the alloys. Alloy A2 was subjected to 1510 C and 1410 C heat-treatments and then water quenched in an attempt to preserve the high temperature equilibrium microstructure. The as-cast button was
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Table 1 Compositional analysis of the heat-treated bulk materials as well as the composition of the individual phases as evaluated by EPMA.
s (at%)
Bulk (at%) Alloy-HT A1-1410 A2-1410 A1-1510 A2-1510
C C C C
h (at%)
g (at%)
Al
Ti
Nb
Al
Ti
Nb
Al
Ti
Nb
Al
Ti
Nb
57.2 51.4 57.4 51.3
7.1 8.8 6.7 8.4
35.7 39.8 35.9 40.3
42.6 41.7 42.2 42.3
9.5 9.1 7.5 7.4
47.9 49.2 50.3 50.3
69.9 69.7 69.9 70.0
4.5 3.7 3.2 3.0
25.6 26.6 26.9 27.0
54.3 54.6 47.8 47.3
11.5 12.1 12.2 11.3
34.1 33.3 40 41.4
sectioned into 1.5 mm thick slices, which were wrapped with tantalum foil. A vertical alumina tube furnace equipped with drop quenching capability was used for heat treating. The samples were suspended within the hot zone by a Ta wire and the alumina tube was sealed by end caps. The tube furnace was purged three times with oxygen gettered high purity Ar and heat treatment was conducted under a flowing Ar gas. The samples were heated to 1510 C and 1410 C at 10 K/min and held isothermally for 4 h at the temperature, followed by drop quenching into a water bath. Direct immersion water quenching was performed by removing the bottom cap and nearly simultaneously dropping the sample into the water bath. The samples were released from the upper cap by a custom built rotating hook. Following the water quench, the samples were ground and polished to remove any environmentally affected zones that may have formed during heat treatment. The microstructure of heat-treated samples was studied using the BSE (back scattered electron) mode in SEM (scanning electron microscope). The large difference in the atomic number of the alloying elements provided sufficient z-contrast between phases thus facilitating microstructural analyses. Conventional TEM (transmission electron microscopy) was conducted on a JEOL 200CX, whereas STEM (scanning TEM) and EDS (energy dispersive spectroscopy) were performed on a JEOL 2010F. High Angle Annular Dark Field imaging (dark field STEM) was employed to facilitate the z-contrast based imaging by the acquisition of only the incoherently scattered electrons that are highly sensitive to the atomic mass. The site-specific TEM foils were prepared by focused ion beam (FIB) sectioning and thinning using a FEI Strata DB 235 instrument.
bright and dark contrasts were found within the prior g-phase. The two distinct phase contrasts confirms a compositional difference between these two phases and suggest a diffusion assisted solidstate phase transformation. The dark-contrast phase in the bright field was located at the s/g interface as well as throughout the transformed g-phase (marked with “1” in Fig. 2). Dark field imaging of this phase revealed that it is the phase with the higher atomic number (Fig. 2b). The bright contrast phase was seen only in the interior of the transformed g-phase region (marked with “2” in Fig. 2). The phase adjacent to the two-phase region was identified as the s-phase. The s-phase produced a uniform contrast throughout the region indicating little to no compositional variation in this phase.
3. Microstructural evaluation Fig. 1a,b present the microstructure of samples quenched from 1510 C to 1410 C, respectively. Obviously, the microstructure equilibriated at 1510 C is coarser than the one formed at 1410 C. The interesting difference between these two microstructures is the appearance of the fine two-phase microstructure within the gphase region in the sample quenched from 1510 C (Fig. 1a). Close observation of the prior g-phase boundaries in this sample revealed a dark contrast phase that seemed to grow into the g-phase from its interface with the s-phase. However, this interfacial phase was not observed at the g/h interface. It should be noted that the XRD analysis did not indicate the presence of another phase in addition to the g, h and s phases in this sample [18]. To further investigate the structure and composition of the micro-constituent phases within the transformed g-phase region, TEM analysis was performed. A 15 mm wide site-specific TEM foil was prepared from the sample heat treated at 1510 C as indicated by the marker shown on Fig. 1a. The section was selected to consist mostly of the transformed g-phase region while cutting through the s/g interface and to include a small region of the s-phase in order to evaluate the dark-contrast boundary phase. A typical bright field/dark field STEM image set is shown in Fig. 2. Consistent with the BSE/SEM image (Fig. 1a) two phases with
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.
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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.
4. Compositional analysis The STEM imaging provided supporting evidence that the two phases with different contrasts were compositionally distinct. TEM/ EDS analysis was used to resolve the compositions of the individual micro-constituents. A bright field STEM image of one of the regions examined by TEM/EDS is shown in Fig. 3. The EDS analysis also revealed two statistically distinct compositions. The composition of region “2” was determined to be 31.2Nb 57.2Al 11.6Ti (at%) while the region “1” contained 45.2Nb 43.9Al 10.9Ti (at%). The compositions corresponding to regions 1 and 2 are marked on a ternary plot, which also shows the tie-triangles consisting of s, h and g phases at 1510 C and 1410 C (Fig. 3) that were measured from both alloys A1 and A2 as listed in Table 1. The compositional analysis elucidates that the bright phase in Fig. 3 (region “2”) rejects sufficient Nb for Al such that its composition is brought near to the equilibrium composition of the g-phase at 1410 C, whereas the lower-z phase marked as “1” is compositionally located near the sphase corner of the 1410 C tie-triangle. Prior work within our research group has shown that the nucleation of the s-phase is kinetically difficult therefore it is unlikely that it nucleated out of the g-phase upon quenching [10,15]. Thus the dark-contrast phase was suspected to be a metastable phase that formed upon quenching. 5. Structural analysis A bright field TEM image of the transformed g-phase is shown in Fig. 4, in which the two phases are marked as “a” and “b”. Structural analysis was performed by recording SAD (selected area diffraction) patterns of phases individually as well as from the two adjacent phases simultaneously. The analysis of the diffraction patterns revealed that the two phases are structurally distinct. The phase “a” was identified through the analysis of the diffraction patterns to be the g-phase. A diffraction pattern near the [110]g zone axis is shown in Fig. 4b. The second phase was determined not to be the g-phase. Inspection of the diffraction patterns of the second phase near its zone axes revealed that the lattice parameter of this phase is close to triple that of the g-phase (Fig. 4c). Among many phases considered, as mentioned previously, there is evidence in the literature of a metastable orthorhombic h-Al2Ti phase that forms in the binary TieAl system upon quenching the g-phase [22e24]. Morphological similarities were found between the h-phase within the TiAl matrix and the g-phase transformation observed in this study.
Fig. 3. A depiction of the compositions of the two-phase region that forms upon quenching from 1510 C demonstrating their relationship to the 1410 C equilibrium tie-triangle.
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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.
A structural model was generated using the Crystal Maker software package from the reported structural data for both the gTiAl [25,26] and the h-Al2Ti phases [22] and are reproduced in Fig. 5. The h-Al2Ti phase is a superlattice that forms by ordering of the Al atoms on the (001) planes of the g-TiAl L10 structure, thus its lattice parameter is inherent to that of parent g-phase. Essentially
the h-phase contains two lattice parameters that are approximately equal to those of the g-phase and the third parameter is about three times of the L10’s lattice parameter. Investigation of the diffraction patterns confirmed that the “b” phase in Fig. 4a is the h-phase. A SAD pattern that covers both phases in the prior g-phase is shown in Fig. 6a. This pattern is taken near the [110]g zone axis, which is also slightly off the [001]h zone axis. The combined diffraction patterns elucidates that the spacing between the (030)h planes is larger than the spacing between the (001)g planes and there is a 35 rotation between these two planes. On the other hand the spacing of the (030)h planes is very close to the spacing of the {110)g planes. The combined recorded and simulated patterns are shown in Fig. 6b, which shows an excellent match between the structures. The foil was tilted slightly in order to reach a two beam condition off the [110]g zone axis. This diffraction pattern was recorded and is shown in Fig. 6c. The ð110Þg and the ð110Þh diffracted beams shown on this figure were used for dark field imaging of each phase. The bright field and dark field images are shown in Fig. 7. The dark field image presented in Fig. 7a was obtained using the ð110Þg reflection and identified the “a” phase as the g-phase. The dark field presented in Fig. 7b was generated using the ð110Þh diffracted beam, which identified the “b” phase as the h-phase. This analysis was combined with the STEM imaging and compositional analyses previously discussed in order to correlate the phase identification
Fig. 6. (a) DPs of the g-phase (near [110] zone) and h-phase (near [001] zone) demonstrating that [110]g//[001]h with a rotation of approximately 35 between (001)g and (030)h planes. (b) A comparison of the simulated diffraction patterns with the actual ones. (c) The diffraction spots used for the bright field/dark field images shown in Fig. 7.
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Fig. 7. BF-TEM image of the two-phase region with the associated dark fields using the ð110Þg and the ð110Þh diffracted beams, respectively.
with the measured composition of each phase. The g-phase shown in Fig. 7 was determined to be the dark phase in the bright field STEM image that is marked with “1” in Figs. 2 and 3. The composition of the g-phase is then identified as 31.2Nb 57.2Al 11.6Ti (at%). The darker-contrast phase in STEM image marked as “2” is found to be the h-phase that has a composition of 45.2Nb 43.9Al 10.9Ti (at%). The composition of this phase suggests that approximately 25% of Al sites and 75% of the Ti sites in the Al2Ti structure are occupied by Nb atoms. Since no extra superlattice spots in addition to the diffraction spots from the h-phase were found, it is suggested that the Nb atoms are distributed randomly. The morphology of the g and h phases and their distinct compositional differences suggests that the high temperature g-phase goes through a spinodal decomposition upon fast cooling. The significant change in the solubility of Nb in the g-phase upon cooling is a key component driving this transformation. An abrupt increase in the chemical potential gradient, driving the diffusion of Nb out of the gphase and into the s-phase is anticipated. However; the long range diffusion is kinetically limited by quenching. Apparently, the enthalpy associated with the mixing of Ti, Al and Nb in the g-phase results in a miscibility gap in the Gibbs free energy surface during the cooling of this thermodynamically unstable phase. Localized short range diffusion driven by the negative curvature in the free energy surface leads to partitioning of Nb and Al within the g-phase. Therefore, the subsequent short range diffusion renders the h-phase formation through a spinodal decomposition of the g-phase. The combined action of the chemical driving forces and limited diffusion drive the nano scaled wormy appearing spinodal microstructure and the formation of the g-phase superstructure (h-phase). The characteristic spacing between the phases of under 200 nm scale suggests that nucleation was not the rate limiting mechanism. This spacing is also consistent with a slightly coarsened spinodal structure [27,28].
6. Summary and conclusions Quenching an alloy with nominal composition of 51.5Al 8.5Ti 40Nb at% (alloy A2) from 1510 C caused a solid-state transformation in the g-phase, whereas a 1410 C heat-treatment and quenching of the same alloy resulted in no detectable structural changes in the g-phase. Prior studies disclosed significant heat evolution or absorption between these two temperatures indicating a compositional and volume fraction change [18]. TEM/STEM analyses revealed that during the transformation upon quenching from 1510 C two phases identified as g and h phases evolve. The composition the g-phase obtained upon quenching was 31.2Nb 57.2Al 11.6Ti at% and it was close to the equilibrium composition of this phase found in the microstructure equilibrated at 1410 C. This is shown in Fig. 3b where now region “1” has been identified as the g-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 does suggesting that Nb is preferentially located in Al lattice positions. Observation of the interface between the s-phase and the transformed g-phase region revealed that the h-phase covered the boundary (Fig. 2). Compositional analysis revealed that the h-phase is a Nb rich phase therefore it is logical that it is found adjacent to the high Nb s-phase. In contrast near the Al rich h-phase boundaries the h-phase was not observed (Fig. 1). This further confirms that this phase forms by the diffusion of Al and Nb. These results suggest that at high temperatures the g-phase has a high solubility for Nb, which decreases drastically with reducing temperature. When the formation of s and h phases is inhibited by water quenching, the metastable g-phase releases its Nb by undergoing a spinodal decomposition.
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Acknowledgements This research was supported by NSF/AFOSR under Grant No. DMR-0605702 and DMR-0856622.
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