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Evolution of B2(ω) region in high-Nb containing TiAl alloy in intermediate temperature range
Intermetallics 82 (2017) 32e39
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Evolution of B2(u) region in high-Nb containing TiAl alloy in intermediate temperature range Xuyang Wang, Jieren Yang*, Lin Song, Hongchao Kou, Jinshan Li, Hengzhi Fu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an 710072, Shaanxi, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 April 2016 Received in revised form 9 August 2016 Accepted 24 November 2016 Available online 1 December 2016
The thermal stability of the B2(u) region within an intermediate temperature range in a Ti-45Al-8.5Nb(W, B, Y) (at.%) alloy was investigated. First, the “pure B2 region” without the u phase was obtained after 1100 C/1 h followed by water quenching. Then, the samples were heat treated at 700 C, 800 C and 850 C for 1 h and 96 h respectively. The results show that B2 decomposes into four types of uo variants and g phase at 700 C; but only one uo variant would remain as the ageing time increased. The size of the uo particles obviously increases with the increase of heating temperature and treating time. It is also found that the phase transformation of B2 / g occurred at 800 C and the g phase prefers to nucleate at B2 region boundaries. gp would further coarsen to micron size after 96 h. In contrast, no g phase in the B2 matrix is observed when the annealing temperature increases to 850 C, which could be attributed to the influence of W element on the formation of g phase in B2 phase. Phase transformation mechanisms in the B2(u) region, involving B2 / uo and B2 / g, were discussed. © 2016 Elsevier Ltd. All rights reserved.
Keywords: Intermetallics Annealing Diffusion Thermal stability Microstructure
1. Introduction TiAl-based intermetallics have been considered as promising structural materials for the aerospace and automotive industries due to low densities and excellent performance at elevated temperatures [1,2]. The addition of transition metal Nb can increase their elevated temperature properties, so high Nb-containing gTiAl alloys have attracted increasing attention. However, Nb is a b phase stabilizer and would introduce b-segregation into TiAl ascast ingots, which usually exists at grain boundaries and triple junctions [3,4]. The disordered bcc b phase would transform into the ordered bo/B2 phase during cooling process. The bo/B2 phase is physically brittle, harmful to the room-temperature mechanical properties of TiAl alloys. Studies have also indicated that the bo/B2 phase is metastable and readily decomposes into u-related and g phases [5,6]. Trigonal u00 , hexagonal B82-u and D88-u phases are u-related phases existed in the B2 area in TiAl alloys with similar crystal structure but different occupations [7]. The nucleation of the u00 structure is inevitable even under fast cooling rate such as water quenching [8], which can be attributed to the short distance
* Corresponding author. E-mail address: [email protected] (J. Yang). http://dx.doi.org/10.1016/j.intermet.2016.11.007 0966-9795/© 2016 Elsevier Ltd. All rights reserved.
diffusion of atoms and displacive progress in which two adjacent B2 planes collapse to an intermediate position along one 〈111〉 direction [9]. The B82-u phase can transform into the D88-u phase with nanoscale during long time aging due to the surplus vacancies in the crystal structure [10,11]. The u-related phases including micron-level uo particles and nanoscale u00 particles tend to occupy most of the B2 area. Huang [7] found that D88-u and B82-u can be detrimental to ductility and fatigue strength in Ti-44Al-4Zr-4Nb0.2Si-1B alloy. Moreover, within the temperature range 700e900 C (i.e. intermediate temperature), the B2 phase can transform into the g phase through the growth of an adjacent lamellar g phase and the nucleation of g directly [12e14]. In the service state, the microstructures of TiAl-based alloys are expected to be stable, which is beneficial to the stability of mechanical properties and the safety of components. Therefore, the phase transformation of the B2 phase at evaluated temperatures should be fundamentally understood. However, the nucleation and evolution of u-related phases, as well as the B2 phase, are not clear. The study focuses on the thermal stability of the B2(u) phase in high Nb-containing TiAl alloy with the nominal composition of Ti45Al-8.5Nb-0.2W-0.2B-0.02Y (at.%). The microstructure evolution and phase transformation within the potential service temperature range of 700e850 C have been studied.
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Fig. 1. Schematic diagram of heat treatment processes starting from as-cast state.
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electro-discharge machining. A pre-heat treatment, annealing at 1100 C for 1 h followed by water quenching (WQ), was carried out on samples to eliminate the u-related phases in the B2 phase. Then these quenched samples were kept at 700, 800, 850 C for 1 h and 96 h respectively in a preheated resistance furnace. Fig. 1 shows a schematic diagram of heat treatment parameters. The furnace was calibrated to ensure the accuracy of temperature. The precision of temperature is estimated to be about 5 C. In order to retain the high-temperature microstructure of the B2 region, the heated samples were water-quenched. The microstructures of treated samples were investigated by scanning electron microscopy (SEM) in back scattered electron (BSE) mode, using a VEGA TESCAN operating at 20 kV. Transmission electron microscopy (TEM) analysis was carried out on a Tecnai G2 F30 TEM operating at 300 kV. Thin foils used for TEM analysis were prepared by twin-jet polishing with an electrolyte of 5 vol% perchloric acid, 30 vol% butan-1-ol and 65 vol% methanol operating at 30 V and 30 C.
2. Experimental TiAl alloy used in this work was produced by a plasma cold hearth melting furnace with an ingot size of 730 340 900 mm. The measured composition is Ti-44.93Al-8.52Nb-0.20W-0.19B0.02Y, which is very close to the nominal composition. Samples with the size of 8 mm 8 mm 8 mm were cut from the ingot by
3. Results 3.1. B2 (u) phase morphologies in as-cast and quenched samples A typical as-cast microstructure of Ti-45Al-8.5Nb-0.2W-0.2B0.02Y alloy is shown in Fig. 2a. The as-cast alloy exhibits a near-
Fig. 2. (a) SEM-BSE image of the as-cast Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy. (b) magnified image of the B2 area. (c) TEM observation of B2(u) region. (d) SADP taken from the B2 matrix.
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lamellar microstructure with the average lamellar colony size of 120 mm, mixed with retained B2 phase (bright contrast) and adjacent equiaxed g phase (dark contrast), which are located at lamellar colony boundaries. These g grains are transformed from the neighbouring B2 phase during the ingot cooling from high temperature. As shown in Fig. 2b and c, amounts of elliptical micron scale uo particles which occupy almost half of the B2 area (about 46 vol%) are distributed within the B2 area. The diffraction pattern (Fig. 2d) with the beam direction is parallel to 〈110〉B2, which indicates that four variants exist within B2 phase including two 〈1120〉uo and two 〈1102〉uo zone axes, but the 〈1102〉uo can not be distinguished from the 〈110〉B2 pattern of the B2 phase [9]. In high-Nb containing TiAl alloys, Nb-segregation would result in the formation of B2 phase. Because the composition of Nbsegregation areas in this alloy is very close to that of Ti4Al3Nb [9], the transformation between the B2 phase and the uo phase is reversible within an intermediate temperature range. After the heat treatment at 1100 C for 1 h, the microstructure has still remained stable, comprised of the a2/g lamellar colony, B2 and massive g phases, as shown in Fig. 3a. The SEM-BSE (Fig. 3b) and TEM (Fig. 3d) show that no uo particles are residual in the B2 phase. However, under the 〈110〉B2 zone axis in the selected area diffraction pattern (SADP) of the area, a weak u-related reflection can be found. The diffuse maxima in the diffraction pattern is typical of the diffuse u structure [15]. The structure of the u0 phase
has four distinct Wyckoff sites, but only two occupancies. This structure is inherently unstable and would transform to the u00 phase through short-range chemical ordering [8], so this kind of omega phase should be u00 . Here the question arises when the formation of u00 phase occurs. On the one hand, it is speculated that the displacive process and short-range chemical ordering could occur during rapidly cooling [16], which can lead to the formation of u-embryos. From this point of view, it should be pointed out that even high cooling rate such as water quenching is not sufficient to prevent the nucleation of the u-related phase from B2 phase in TiAl alloys [17]. On the other hand, it is proposed that u-related embryos already exist in the bcc structure such as Ti, Zr and Hf [18]. These embryos are probably stabilized by defects. Studies of bcc Zr and Zr-Nb alloys have proved that the existence temperature of these u-embryos is much higher than that of the u-start temperature (Tu) [19,20]. Moreover, the uembryos are present as dynamical u fluctuations with very short life-times. It is rational that the dynamical u fluctuations may also exist in the B2 phase at 1100 C. 3.2. Phase composition of the B2 (u) area after annealing at 700 C Fig. 4a shows the morphology of the B2(uo) region after annealing at 700 C for 1 h. It can be found that the morphology of the B2 area has no apparent change. uo particles with gray contrast
Fig. 3. (a) SEM-BSE image of the water quenched sample, (b) magnified image of the retained B2 phase. (c) TEM image showing the pure B2 phase. (d) SADP taken from the B2 phase.
X. Wang et al. / Intermetallics 82 (2017) 32e39
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Fig. 4. Reprecipitate of the uo phase after annealing at 700 C for 1 h (a) SEM-BSE image showing the morphology of the microstructure after annealing for 1 h. The corresponding image of the B2(u) region obtained by TEM is shown in (b).
Fig. 5. (a) SEM-BSE image showing the microstructure after annealing at 700 C for 96 h. The corresponding images of the B2(u) regions obtained by TEM are shown in (b) and (c). (d) SADP taken from the B2(u) matrix, revealing that the crystal orientation of uo particles is same.
morphology are hardly found within the B2 phase, as the as-cast microstructure. The TEM pattern in Fig. 4b indicates that the uo phase has already nucleated and grown within the B2 phase after
1 h of annealing. The morphology of the B2 phase with very fine uo particles exhibits a misty feature. These uo particles can not be clearly resolved in SEM-BSE observation due to their small sizes, so
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Fig. 4a just displays a uniform background contrast. The morphology of the B2(uo) region indicates that the uo phase transforms from the B2 phase by nucleation within the B2 phase. The reappearance of the uo phase after annealing at high temperature seems somewhat related to the existence of u embryos in the as-quenched B2 structure [8]. At 700 C, the u-embryos may grow simultaneously. When the annealing time reaches 96 h, the SEM-BSE image also shows a uniform background contrast within B2(u) region as annealing for 1 h. The TEM pattern in Fig. 5b shows the feature of B2 region. Multi-complex phases exist within the B2 matrix. Besides the uo phase, g phase also appears within the B2 matrix when the annealing time is prolonged. Studies on Tie44Ale4Nbe4Zre0.2Si and high-Nb containing TiAl alloy have proved that the direct transformation of B2 / g is feasible [11,13]. From Fig. 5b, several g particles distribute at the boundary between two B2(u) matrixes can be observed, indicating that the boundaries of B2 grains can be the prior nucleation sites of the g phase. In Cheng's study [12], there is no W in the alloy and the thermal stability of B2 phase is lower, so the decomposition of B2 / g can be observed even if rapidly cooled. However, in our study, the addition of W leads to higher stability of B2, so the B2 / g transformation delays and the degree of the decomposition is smaller. The size of the uo particles is also very small as shown in Fig. 5c (100e400 nm). This may be due to the low diffusion at 700 C that restricts the coarsening of the nano-
sized uo particles. Moreover, unlike annealing for 1 h, uo particles are composed of one orientation variant rather than four orientations, as shown in Fig. 5d. Three of the four uo phase variants have gradually dissolved into the B2 phase and only one variant is left. Generally speaking, the four variants are identical, so the volume fraction should be the same. In shock-loaded zirconium, only two of the six variants for the reverted a phase from the u phase are observed due to the pressure configuration [21]. However, there are no such motivational factors here, phase transformation residual stress or the free energy of the system may be the reason and this phenomenon should be studied further. 3.3. Phase composition of the B2 (u) area after annealing at 800 C After annealing at 800 C for 1 h, the SEM-BSE image shows no apparent change of contrast within the B2 matrix annealing for 1 h, as shown in Fig. 6a. Fig. 6b shows that abundant particles with a size of about 50e100 nm are homogeneously distributed in the B2(u) area when the sample is annealed at 800 C for 1 h. SADP confirms that these small particles are uo phase with four rotational orientations. Because their sizes are beyond the resolution, the SEM-BSE image shows no contrast within the B2 matrix. The size of uo particle is larger when compared to Fig. 4b due to the higher diffusion rate of elements at 800 C. Increasing the annealing time to 96 h, a large number of g
Fig. 6. Phase composition of B2(u) area after annealing at 800 C. (a) and (c) SEM-BSE images showing the microstructures of the B2 matrix for 1 h and 96 h, respectively. (b) and (d) corresponding images of the B2(u) regions from (a) and (c) obtained by TEM.
X. Wang et al. / Intermetallics 82 (2017) 32e39
precipitates (dark contrast) are observed in the B2(u) area, coexisting with the uo particles (Fig. 6c), which have been transformed from the B2 phase [22]. Fig. 6d shows the detailed morphology of the uo particles and g precipitates. The inset at the high-right corner in Fig. 6d is the SADP taken from a single g plate with 〈011〉g direction. The g phase has coarsened evidently, and the size has reached micron scale. The B2(u) region is occupied by lathshaped uo particles. In a certain area, these uo particles have the same length direction. The inset in the low-right corner of Fig. 6d shows the SADP of this area, revealing that the crystallographic orientation of these uo particles is the same and only one orientation variant remains. Interface energy minimization may be responsible for the phenomenon and promotes the growth of a certain crystal plane. It is difficult to check whether there are remainders of the B2 phase from Fig. 6d, but due to the addition of W element, it is impossible that the B2 phase is decomposed completely [11]. 3.4. Phase composition of the B2(u) area after annealing at 850 C As annealing at temperature of 700 C and 800 C, there is no apparent contrast change occurring in the B2 area after annealing at 850 C for 1 h according to the SEM-BSE image of Fig. 7a. Fig. 7b shows the bright field image of B2 matrix, indicating that particles with nanoscale have already formed during the annealing. The
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SADP confirms that these particles are the uo phase. It should be pointed out that these uo particles are in the same orientation as shown in the inset in the bottom-right corner of Fig. 7b. This is different to the microstructure of annealing at 700 and 800 C: there is only one uo variant when annealing at 850 C for 1 h, and the nucleation mechanism of the uo phase may be different to that at 700 and 800 C. A more precise analysis will be undertaken in a future study. 4. Discussion In high Nb-containing TiAl alloys, Nb-segregation leads to the formation of B2 phase. Because the composition of Nb-segregation areas in this alloy is very close to that of Ti4Al3Nb, the B2 phase and the uo phase can interconvert between each other within the intermediate temperature range. The nucleation process of u00 structure can not be prevented even by water quenching from 1100 C. The B2 phase is unstable in the temperature range of 700e850 C, and readily transform into the uo and g phases. The phase transformation mechanism and types are different at 700 C, 800 C and 850 C in the B2(u) matrix although the temperature difference is small. When annealing at 700 C and 800 C for 1 h, four uo variants with different orientations have occupied the B2 matrix. However, with the ageing time increasing, it is interesting to find that three of
Fig. 7. Reprecipitate of uo phase with the increase of annealing time at 850 C. (a), (c) SEM-BSE image showing the microstructure after annealing for 1 h and 96 h, respectively. The corresponding images of the B2(u) regions obtained by TEM are shown in (b) and (d).
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the four variants would dissolve back into the B2 phase and only one variant is left; whereas the four uo variants are thought to be completely the same. It is quite clear that the nucleation rates of the four uo variants are the same at 700 and 800 C. With the ageing time extension, three of the four variants would fade away. This may be related to the minimization of interface energy, and the growth rate of uo particles is crystallographically anisotropic. The element diffusion rate is lower at 700 C than 800 C and the growth of uo grains is limited, so the sizes of uo grains annealed at 700 C is smaller than that at 800 C. The enhanced long-range diffusion induced by higher annealing temperature at 850 C makes the phase transformation process more quickly [23]. In addition, to reduce the interface energy, the adjacent uo particles would aggregate to a larger size, leading to a more stable structure [24]. Moreover, coarsening would occur more easily as the uo particles have a consistent orientation relationship with the matrix [25]. For the aforementioned reasons, the uo grains can grow to the micrometre scale, as shown in Fig. 7c. The B2 / g transformation would occur when aging at 700 C. Besides being inside the B2 matrix, the g grains are inclined to nucleate at the boundaries between B2 grains, because there are more g grains that exist there (Fig. 5b). Boundaries can act as the concentration sites of defects, providing more nucleation sites for the g phase [26]. Small g grains are present along the B2 phase boundaries, which is similar to the TNM alloys when annealing at high temperature [13]. However, the sizes of the g precipitates in this study are much smaller than that in TNM alloys [27,28]. Taking the above into consideration, the addition of W must play an important role in the restriction of the g precipitate. No g phase has been found to exist in the B2(u) matrix when the alloy is annealed at 850 C for 96 h from the SEM-BSE and TEM image in Fig. 7. Annealing at 800 C for 96 h, the g plates coarsen to micron-size, similar to the results of Tang's study [29]. W content evidently decreases in the uo phase and increases in the B2 phase [23]. With the temperature increasing, the diffusion rate of the W element becomes higher. With the growth of the uo phase, W is expelled from the uo phase and concentrated in the residual B2 phase, so the B2 phase becomes quickly stable when ageing at 850 C. As a result, the B2 / g transformation would be restricted. When annealing at 700 and 800 C, the diffusion rate of the element is lower. The concentration of W in the B2 phase is sluggish, so the B2 / g transformation can take place.
5. Conclusions Within intermediate temperature range, the phase transformations in the B2(u) region of a high Nb-TiAl alloy (Ti-45Al8.5Nb-0.2W-0.2B-0.02Y) are sensitive to heating temperature and time. The main results are summarized as follows: (1) After 1100 C/1 h þ WQ, the uo phase almost disappears and only weak u-related structures remain in the B2 region. Annealing at intermediate temperatures, the B2 phase is thermodynamically unstable and transforms to the uo and g phases. (2) Holding at 700 C and 800 C, four uo variants with different orientations nucleate. With the prolonging of annealing time, three of the four variants fade away and only one variant remains. Increasing the annealing temperature from 700 C to 850 C, the sizes of the uo particles coarsen from nano-to micron-level. (3) The g phase prefers to nucleate at the boundaries between B2 grains. The transformation of B2 / g would be promoted by increasing the annealing temperature from 700 C to
800 C. In contrast, it is more difficult to form the g phase when the temperature further increased to 850 C. Acknowledgments The financial support from the National Natural Science Foundation of China (No. 51401168), the Aeronautical Science Foundation of China (No. 2014ZF53065) and the Fundamental Research Funds for the Central Universities (No. 3102014JCQ01026) is gratefully acknowleged. References [1] X. Wu, Review of alloy and process development of TiAl alloys, Intermetallics 14 (2006) 1114e1122. [2] K. Kothari, R. Radhakrishnan, N.M. Wereley, Advances in gamma titanium aluminides and their manufacturing techniques, Prog. Aero. Sci. 55 (2012) 1e16. [3] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Microsegregation in high Nb containing TiAl alloy ingots beyond laboratory scale, Intermetallics 15 (2007) 625e631. [4] X.J. Xu, J.P. Lin, Y.L. Wang, Z. Lin, G.L. Chen, Deformability and microstructure transformation of pilot ingot of Tie45Ale(8e9)Nbe(W, B,Y) alloy, Mater. Sci. Eng. A 416 (2006) 98e103. [5] K.J. Leonard, R. Tewari, A. Arya, J.C. Mishurda, G.K. Dey, V.K. Vasudevan, Decomposition of the bo phase to the P63/mcm, hP18 structure in Nbe(24e36)Tie40Al alloys, Acta Mater. 57 (2009) 4440e4453. [6] T.H. Yu, C.H. Koo, Microstructural evolution of a hot-rolled Tie40Ale10Nb alloy, Mater. Sci. Eng. A 239e240 (1997) 694e701. [7] Z.W. Huang, Ordered u phases in a 4Zre4Nb-containing TiAl-based alloy, Acta Mater. 56 (2008) 1689e1700. [8] L. Song, X.J. Xu, J. Sun, J.P. Lin, Cooling rate effects on the microstructure evolution in the bo zones of cast Tie45Ale8.5Nbe(W, B,Y) alloy, Mater. Charact. 93 (2014) 62e67. [9] L.A. Bendersky, W.J. Boettinger, B.P. Burton, F.S. Biancaniello, C.B. Shoemaker, The formation of ordered u-related phases in alloys of composition Ti4Al3Nb, Acta Metall. Mater. 38 (1990) 931e943. [10] Z.W. Huang, Thermal stability of Tie44Ale4Nbe4Zre0.2Sie1B alloy, Intermetallics 42 (2013) 170e179. [11] L. Song, X.J. Xu, L. You, Y.F. Liang, J.P. Lin, Phase transformation and decomposition mechanisms of the bo(u) phase in cast high Nb containing TiAl alloy, J. Alloy. Compd. 616 (2014) 483e491. [12] T.T. Cheng, M.H. Loretto, The decomposition of the beta phase in Tie44Ale8Nb and Tie44Ale4Nbe4Zre0.2Si alloys, Acta Mater. 46 (1998) 4801e4819. € berl, E. Schwaighofer, Z. Zhang, H. Clemens, [13] M. Schloffer, B. Rashkova, T. Scho S. Mayer, Evolution of the uo phase in a b-stabilized multi-phase TiAl alloy and its effect on hardness, Acta Mater. 64 (2014) 241e252. [14] D. Daloz, U. Hecht, J. Zollinger, H. Combeau, A. Hazotte, M. Zalo znik, Microsegregation, macrosegregation and related phase transformations in TiAl alloys, Intermetallics 19 (2011) 749e756. [15] G. Shao, P. Tsakiropoulos, Prediction of u phase formation in TieAleX alloys, Mater. Sci. Eng. A 329e331 (2002) 914e919. [16] A. Stark, M. Oehring, F. Pyczak, A. Schreyer, Situ observation of various phase transformation paths in Nb-Rich TiAl alloys during quenching with different rates, Adv. Eng. Mater. 13 (2011) 700e704. [17] C.W. Dawson, S.L. Sass, The as-quenched form of the omega phase in Zr-Nb alloys, Metall. Trans. 1 (1970) 2225e2233. [18] R. Tewari, P. Mukhopadhyay, S. Banerjee, L.A. Bendersky, Evolution of ordered u phases in (Zr3Al)eNb alloys, Acta Mater. 47 (1999) 1307e1323. [19] T.S. Kuan, S.L. Sass, The structure of a linear omega-like vacancy defect in ZrNb B.C.C. solid solutions, Acta Metall. 24 (1976) 1053e1059. [20] S.K. Sikka, Y.K. Vohra, R. Chidambaram, Omega phase in materials, Prog. Mater. Sci. 27 (1982) 245e310. [21] S.G. Song, G.T. Gray III, Microscopic and crystallographic aspects of retained omega phase in shock-loaded zirconium and its formation mechanism, Philos. Mag. A 71 (1995) 275e290. [22] L. Song, X. Xu, L. You, Y. Liang, J. Lin, Ordered u phase transformations in Ti45Al-8.5Nb-0.2B alloy, Intermetallics 65 (2015) 22e28. [23] R. Kainuma, M. Palm, G. Inden, Solid-phase equilibria in the Ti-rich part of the TiAl system, Intermetallics 2 (1994) 321e332. [24] B. Shao, Y. Zong, D. Wen, Y. Tian, D. Shan, Investigation of the phase transformations in Ti22Al25Nb alloy, Mater. Charact. 114 (2016) 75e78. [25] D. Ping, Review on u phase in body-centered cubic metals and alloys, Acta Metall. Sin. Engl. Lett. 27 (2014) 1e11. [26] E. Sukedai, D. Yoshimitsu, H. Matsumoto, H. Hashimoto, M. Kiritani, b to u phase transformation due to aging in a TieMo alloy deformed in impact compression, Mater. Sci. Eng. A 350 (2003) 133e138. [27] E. Schwaighofer, H. Clemens, S. Mayer, J. Lindemann, J. Klose, W. Smarsly, V. Güther, Microstructural design and mechanical properties of a cast and heat-treated intermetallic multi-phase g-TiAl based alloy, Intermetallics 44
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Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Evolution of B2(u) region in high-Nb containing TiAl alloy in intermediate temperature range Xuyang Wang, Jieren Yang*, Lin Song, Hongchao Kou, Jinshan Li, Hengzhi Fu State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an 710072, Shaanxi, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 April 2016 Received in revised form 9 August 2016 Accepted 24 November 2016 Available online 1 December 2016
The thermal stability of the B2(u) region within an intermediate temperature range in a Ti-45Al-8.5Nb(W, B, Y) (at.%) alloy was investigated. First, the “pure B2 region” without the u phase was obtained after 1100 C/1 h followed by water quenching. Then, the samples were heat treated at 700 C, 800 C and 850 C for 1 h and 96 h respectively. The results show that B2 decomposes into four types of uo variants and g phase at 700 C; but only one uo variant would remain as the ageing time increased. The size of the uo particles obviously increases with the increase of heating temperature and treating time. It is also found that the phase transformation of B2 / g occurred at 800 C and the g phase prefers to nucleate at B2 region boundaries. gp would further coarsen to micron size after 96 h. In contrast, no g phase in the B2 matrix is observed when the annealing temperature increases to 850 C, which could be attributed to the influence of W element on the formation of g phase in B2 phase. Phase transformation mechanisms in the B2(u) region, involving B2 / uo and B2 / g, were discussed. © 2016 Elsevier Ltd. All rights reserved.
Keywords: Intermetallics Annealing Diffusion Thermal stability Microstructure
1. Introduction TiAl-based intermetallics have been considered as promising structural materials for the aerospace and automotive industries due to low densities and excellent performance at elevated temperatures [1,2]. The addition of transition metal Nb can increase their elevated temperature properties, so high Nb-containing gTiAl alloys have attracted increasing attention. However, Nb is a b phase stabilizer and would introduce b-segregation into TiAl ascast ingots, which usually exists at grain boundaries and triple junctions [3,4]. The disordered bcc b phase would transform into the ordered bo/B2 phase during cooling process. The bo/B2 phase is physically brittle, harmful to the room-temperature mechanical properties of TiAl alloys. Studies have also indicated that the bo/B2 phase is metastable and readily decomposes into u-related and g phases [5,6]. Trigonal u00 , hexagonal B82-u and D88-u phases are u-related phases existed in the B2 area in TiAl alloys with similar crystal structure but different occupations [7]. The nucleation of the u00 structure is inevitable even under fast cooling rate such as water quenching [8], which can be attributed to the short distance
* Corresponding author. E-mail address: [email protected] (J. Yang). http://dx.doi.org/10.1016/j.intermet.2016.11.007 0966-9795/© 2016 Elsevier Ltd. All rights reserved.
diffusion of atoms and displacive progress in which two adjacent B2 planes collapse to an intermediate position along one 〈111〉 direction [9]. The B82-u phase can transform into the D88-u phase with nanoscale during long time aging due to the surplus vacancies in the crystal structure [10,11]. The u-related phases including micron-level uo particles and nanoscale u00 particles tend to occupy most of the B2 area. Huang [7] found that D88-u and B82-u can be detrimental to ductility and fatigue strength in Ti-44Al-4Zr-4Nb0.2Si-1B alloy. Moreover, within the temperature range 700e900 C (i.e. intermediate temperature), the B2 phase can transform into the g phase through the growth of an adjacent lamellar g phase and the nucleation of g directly [12e14]. In the service state, the microstructures of TiAl-based alloys are expected to be stable, which is beneficial to the stability of mechanical properties and the safety of components. Therefore, the phase transformation of the B2 phase at evaluated temperatures should be fundamentally understood. However, the nucleation and evolution of u-related phases, as well as the B2 phase, are not clear. The study focuses on the thermal stability of the B2(u) phase in high Nb-containing TiAl alloy with the nominal composition of Ti45Al-8.5Nb-0.2W-0.2B-0.02Y (at.%). The microstructure evolution and phase transformation within the potential service temperature range of 700e850 C have been studied.
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Fig. 1. Schematic diagram of heat treatment processes starting from as-cast state.
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electro-discharge machining. A pre-heat treatment, annealing at 1100 C for 1 h followed by water quenching (WQ), was carried out on samples to eliminate the u-related phases in the B2 phase. Then these quenched samples were kept at 700, 800, 850 C for 1 h and 96 h respectively in a preheated resistance furnace. Fig. 1 shows a schematic diagram of heat treatment parameters. The furnace was calibrated to ensure the accuracy of temperature. The precision of temperature is estimated to be about 5 C. In order to retain the high-temperature microstructure of the B2 region, the heated samples were water-quenched. The microstructures of treated samples were investigated by scanning electron microscopy (SEM) in back scattered electron (BSE) mode, using a VEGA TESCAN operating at 20 kV. Transmission electron microscopy (TEM) analysis was carried out on a Tecnai G2 F30 TEM operating at 300 kV. Thin foils used for TEM analysis were prepared by twin-jet polishing with an electrolyte of 5 vol% perchloric acid, 30 vol% butan-1-ol and 65 vol% methanol operating at 30 V and 30 C.
2. Experimental TiAl alloy used in this work was produced by a plasma cold hearth melting furnace with an ingot size of 730 340 900 mm. The measured composition is Ti-44.93Al-8.52Nb-0.20W-0.19B0.02Y, which is very close to the nominal composition. Samples with the size of 8 mm 8 mm 8 mm were cut from the ingot by
3. Results 3.1. B2 (u) phase morphologies in as-cast and quenched samples A typical as-cast microstructure of Ti-45Al-8.5Nb-0.2W-0.2B0.02Y alloy is shown in Fig. 2a. The as-cast alloy exhibits a near-
Fig. 2. (a) SEM-BSE image of the as-cast Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy. (b) magnified image of the B2 area. (c) TEM observation of B2(u) region. (d) SADP taken from the B2 matrix.
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lamellar microstructure with the average lamellar colony size of 120 mm, mixed with retained B2 phase (bright contrast) and adjacent equiaxed g phase (dark contrast), which are located at lamellar colony boundaries. These g grains are transformed from the neighbouring B2 phase during the ingot cooling from high temperature. As shown in Fig. 2b and c, amounts of elliptical micron scale uo particles which occupy almost half of the B2 area (about 46 vol%) are distributed within the B2 area. The diffraction pattern (Fig. 2d) with the beam direction is parallel to 〈110〉B2, which indicates that four variants exist within B2 phase including two 〈1120〉uo and two 〈1102〉uo zone axes, but the 〈1102〉uo can not be distinguished from the 〈110〉B2 pattern of the B2 phase [9]. In high-Nb containing TiAl alloys, Nb-segregation would result in the formation of B2 phase. Because the composition of Nbsegregation areas in this alloy is very close to that of Ti4Al3Nb [9], the transformation between the B2 phase and the uo phase is reversible within an intermediate temperature range. After the heat treatment at 1100 C for 1 h, the microstructure has still remained stable, comprised of the a2/g lamellar colony, B2 and massive g phases, as shown in Fig. 3a. The SEM-BSE (Fig. 3b) and TEM (Fig. 3d) show that no uo particles are residual in the B2 phase. However, under the 〈110〉B2 zone axis in the selected area diffraction pattern (SADP) of the area, a weak u-related reflection can be found. The diffuse maxima in the diffraction pattern is typical of the diffuse u structure [15]. The structure of the u0 phase
has four distinct Wyckoff sites, but only two occupancies. This structure is inherently unstable and would transform to the u00 phase through short-range chemical ordering [8], so this kind of omega phase should be u00 . Here the question arises when the formation of u00 phase occurs. On the one hand, it is speculated that the displacive process and short-range chemical ordering could occur during rapidly cooling [16], which can lead to the formation of u-embryos. From this point of view, it should be pointed out that even high cooling rate such as water quenching is not sufficient to prevent the nucleation of the u-related phase from B2 phase in TiAl alloys [17]. On the other hand, it is proposed that u-related embryos already exist in the bcc structure such as Ti, Zr and Hf [18]. These embryos are probably stabilized by defects. Studies of bcc Zr and Zr-Nb alloys have proved that the existence temperature of these u-embryos is much higher than that of the u-start temperature (Tu) [19,20]. Moreover, the uembryos are present as dynamical u fluctuations with very short life-times. It is rational that the dynamical u fluctuations may also exist in the B2 phase at 1100 C. 3.2. Phase composition of the B2 (u) area after annealing at 700 C Fig. 4a shows the morphology of the B2(uo) region after annealing at 700 C for 1 h. It can be found that the morphology of the B2 area has no apparent change. uo particles with gray contrast
Fig. 3. (a) SEM-BSE image of the water quenched sample, (b) magnified image of the retained B2 phase. (c) TEM image showing the pure B2 phase. (d) SADP taken from the B2 phase.
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Fig. 4. Reprecipitate of the uo phase after annealing at 700 C for 1 h (a) SEM-BSE image showing the morphology of the microstructure after annealing for 1 h. The corresponding image of the B2(u) region obtained by TEM is shown in (b).
Fig. 5. (a) SEM-BSE image showing the microstructure after annealing at 700 C for 96 h. The corresponding images of the B2(u) regions obtained by TEM are shown in (b) and (c). (d) SADP taken from the B2(u) matrix, revealing that the crystal orientation of uo particles is same.
morphology are hardly found within the B2 phase, as the as-cast microstructure. The TEM pattern in Fig. 4b indicates that the uo phase has already nucleated and grown within the B2 phase after
1 h of annealing. The morphology of the B2 phase with very fine uo particles exhibits a misty feature. These uo particles can not be clearly resolved in SEM-BSE observation due to their small sizes, so
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Fig. 4a just displays a uniform background contrast. The morphology of the B2(uo) region indicates that the uo phase transforms from the B2 phase by nucleation within the B2 phase. The reappearance of the uo phase after annealing at high temperature seems somewhat related to the existence of u embryos in the as-quenched B2 structure [8]. At 700 C, the u-embryos may grow simultaneously. When the annealing time reaches 96 h, the SEM-BSE image also shows a uniform background contrast within B2(u) region as annealing for 1 h. The TEM pattern in Fig. 5b shows the feature of B2 region. Multi-complex phases exist within the B2 matrix. Besides the uo phase, g phase also appears within the B2 matrix when the annealing time is prolonged. Studies on Tie44Ale4Nbe4Zre0.2Si and high-Nb containing TiAl alloy have proved that the direct transformation of B2 / g is feasible [11,13]. From Fig. 5b, several g particles distribute at the boundary between two B2(u) matrixes can be observed, indicating that the boundaries of B2 grains can be the prior nucleation sites of the g phase. In Cheng's study [12], there is no W in the alloy and the thermal stability of B2 phase is lower, so the decomposition of B2 / g can be observed even if rapidly cooled. However, in our study, the addition of W leads to higher stability of B2, so the B2 / g transformation delays and the degree of the decomposition is smaller. The size of the uo particles is also very small as shown in Fig. 5c (100e400 nm). This may be due to the low diffusion at 700 C that restricts the coarsening of the nano-
sized uo particles. Moreover, unlike annealing for 1 h, uo particles are composed of one orientation variant rather than four orientations, as shown in Fig. 5d. Three of the four uo phase variants have gradually dissolved into the B2 phase and only one variant is left. Generally speaking, the four variants are identical, so the volume fraction should be the same. In shock-loaded zirconium, only two of the six variants for the reverted a phase from the u phase are observed due to the pressure configuration [21]. However, there are no such motivational factors here, phase transformation residual stress or the free energy of the system may be the reason and this phenomenon should be studied further. 3.3. Phase composition of the B2 (u) area after annealing at 800 C After annealing at 800 C for 1 h, the SEM-BSE image shows no apparent change of contrast within the B2 matrix annealing for 1 h, as shown in Fig. 6a. Fig. 6b shows that abundant particles with a size of about 50e100 nm are homogeneously distributed in the B2(u) area when the sample is annealed at 800 C for 1 h. SADP confirms that these small particles are uo phase with four rotational orientations. Because their sizes are beyond the resolution, the SEM-BSE image shows no contrast within the B2 matrix. The size of uo particle is larger when compared to Fig. 4b due to the higher diffusion rate of elements at 800 C. Increasing the annealing time to 96 h, a large number of g
Fig. 6. Phase composition of B2(u) area after annealing at 800 C. (a) and (c) SEM-BSE images showing the microstructures of the B2 matrix for 1 h and 96 h, respectively. (b) and (d) corresponding images of the B2(u) regions from (a) and (c) obtained by TEM.
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precipitates (dark contrast) are observed in the B2(u) area, coexisting with the uo particles (Fig. 6c), which have been transformed from the B2 phase [22]. Fig. 6d shows the detailed morphology of the uo particles and g precipitates. The inset at the high-right corner in Fig. 6d is the SADP taken from a single g plate with 〈011〉g direction. The g phase has coarsened evidently, and the size has reached micron scale. The B2(u) region is occupied by lathshaped uo particles. In a certain area, these uo particles have the same length direction. The inset in the low-right corner of Fig. 6d shows the SADP of this area, revealing that the crystallographic orientation of these uo particles is the same and only one orientation variant remains. Interface energy minimization may be responsible for the phenomenon and promotes the growth of a certain crystal plane. It is difficult to check whether there are remainders of the B2 phase from Fig. 6d, but due to the addition of W element, it is impossible that the B2 phase is decomposed completely [11]. 3.4. Phase composition of the B2(u) area after annealing at 850 C As annealing at temperature of 700 C and 800 C, there is no apparent contrast change occurring in the B2 area after annealing at 850 C for 1 h according to the SEM-BSE image of Fig. 7a. Fig. 7b shows the bright field image of B2 matrix, indicating that particles with nanoscale have already formed during the annealing. The
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SADP confirms that these particles are the uo phase. It should be pointed out that these uo particles are in the same orientation as shown in the inset in the bottom-right corner of Fig. 7b. This is different to the microstructure of annealing at 700 and 800 C: there is only one uo variant when annealing at 850 C for 1 h, and the nucleation mechanism of the uo phase may be different to that at 700 and 800 C. A more precise analysis will be undertaken in a future study. 4. Discussion In high Nb-containing TiAl alloys, Nb-segregation leads to the formation of B2 phase. Because the composition of Nb-segregation areas in this alloy is very close to that of Ti4Al3Nb, the B2 phase and the uo phase can interconvert between each other within the intermediate temperature range. The nucleation process of u00 structure can not be prevented even by water quenching from 1100 C. The B2 phase is unstable in the temperature range of 700e850 C, and readily transform into the uo and g phases. The phase transformation mechanism and types are different at 700 C, 800 C and 850 C in the B2(u) matrix although the temperature difference is small. When annealing at 700 C and 800 C for 1 h, four uo variants with different orientations have occupied the B2 matrix. However, with the ageing time increasing, it is interesting to find that three of
Fig. 7. Reprecipitate of uo phase with the increase of annealing time at 850 C. (a), (c) SEM-BSE image showing the microstructure after annealing for 1 h and 96 h, respectively. The corresponding images of the B2(u) regions obtained by TEM are shown in (b) and (d).
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the four variants would dissolve back into the B2 phase and only one variant is left; whereas the four uo variants are thought to be completely the same. It is quite clear that the nucleation rates of the four uo variants are the same at 700 and 800 C. With the ageing time extension, three of the four variants would fade away. This may be related to the minimization of interface energy, and the growth rate of uo particles is crystallographically anisotropic. The element diffusion rate is lower at 700 C than 800 C and the growth of uo grains is limited, so the sizes of uo grains annealed at 700 C is smaller than that at 800 C. The enhanced long-range diffusion induced by higher annealing temperature at 850 C makes the phase transformation process more quickly [23]. In addition, to reduce the interface energy, the adjacent uo particles would aggregate to a larger size, leading to a more stable structure [24]. Moreover, coarsening would occur more easily as the uo particles have a consistent orientation relationship with the matrix [25]. For the aforementioned reasons, the uo grains can grow to the micrometre scale, as shown in Fig. 7c. The B2 / g transformation would occur when aging at 700 C. Besides being inside the B2 matrix, the g grains are inclined to nucleate at the boundaries between B2 grains, because there are more g grains that exist there (Fig. 5b). Boundaries can act as the concentration sites of defects, providing more nucleation sites for the g phase [26]. Small g grains are present along the B2 phase boundaries, which is similar to the TNM alloys when annealing at high temperature [13]. However, the sizes of the g precipitates in this study are much smaller than that in TNM alloys [27,28]. Taking the above into consideration, the addition of W must play an important role in the restriction of the g precipitate. No g phase has been found to exist in the B2(u) matrix when the alloy is annealed at 850 C for 96 h from the SEM-BSE and TEM image in Fig. 7. Annealing at 800 C for 96 h, the g plates coarsen to micron-size, similar to the results of Tang's study [29]. W content evidently decreases in the uo phase and increases in the B2 phase [23]. With the temperature increasing, the diffusion rate of the W element becomes higher. With the growth of the uo phase, W is expelled from the uo phase and concentrated in the residual B2 phase, so the B2 phase becomes quickly stable when ageing at 850 C. As a result, the B2 / g transformation would be restricted. When annealing at 700 and 800 C, the diffusion rate of the element is lower. The concentration of W in the B2 phase is sluggish, so the B2 / g transformation can take place.
5. Conclusions Within intermediate temperature range, the phase transformations in the B2(u) region of a high Nb-TiAl alloy (Ti-45Al8.5Nb-0.2W-0.2B-0.02Y) are sensitive to heating temperature and time. The main results are summarized as follows: (1) After 1100 C/1 h þ WQ, the uo phase almost disappears and only weak u-related structures remain in the B2 region. Annealing at intermediate temperatures, the B2 phase is thermodynamically unstable and transforms to the uo and g phases. (2) Holding at 700 C and 800 C, four uo variants with different orientations nucleate. With the prolonging of annealing time, three of the four variants fade away and only one variant remains. Increasing the annealing temperature from 700 C to 850 C, the sizes of the uo particles coarsen from nano-to micron-level. (3) The g phase prefers to nucleate at the boundaries between B2 grains. The transformation of B2 / g would be promoted by increasing the annealing temperature from 700 C to
800 C. In contrast, it is more difficult to form the g phase when the temperature further increased to 850 C. Acknowledgments The financial support from the National Natural Science Foundation of China (No. 51401168), the Aeronautical Science Foundation of China (No. 2014ZF53065) and the Fundamental Research Funds for the Central Universities (No. 3102014JCQ01026) is gratefully acknowleged. References [1] X. Wu, Review of alloy and process development of TiAl alloys, Intermetallics 14 (2006) 1114e1122. [2] K. Kothari, R. Radhakrishnan, N.M. Wereley, Advances in gamma titanium aluminides and their manufacturing techniques, Prog. Aero. Sci. 55 (2012) 1e16. [3] G.L. Chen, X.J. Xu, Z.K. Teng, Y.L. Wang, J.P. Lin, Microsegregation in high Nb containing TiAl alloy ingots beyond laboratory scale, Intermetallics 15 (2007) 625e631. [4] X.J. Xu, J.P. Lin, Y.L. Wang, Z. Lin, G.L. Chen, Deformability and microstructure transformation of pilot ingot of Tie45Ale(8e9)Nbe(W, B,Y) alloy, Mater. Sci. Eng. A 416 (2006) 98e103. [5] K.J. Leonard, R. Tewari, A. Arya, J.C. Mishurda, G.K. Dey, V.K. Vasudevan, Decomposition of the bo phase to the P63/mcm, hP18 structure in Nbe(24e36)Tie40Al alloys, Acta Mater. 57 (2009) 4440e4453. [6] T.H. Yu, C.H. Koo, Microstructural evolution of a hot-rolled Tie40Ale10Nb alloy, Mater. Sci. Eng. A 239e240 (1997) 694e701. [7] Z.W. Huang, Ordered u phases in a 4Zre4Nb-containing TiAl-based alloy, Acta Mater. 56 (2008) 1689e1700. [8] L. Song, X.J. Xu, J. Sun, J.P. Lin, Cooling rate effects on the microstructure evolution in the bo zones of cast Tie45Ale8.5Nbe(W, B,Y) alloy, Mater. Charact. 93 (2014) 62e67. [9] L.A. Bendersky, W.J. Boettinger, B.P. Burton, F.S. Biancaniello, C.B. Shoemaker, The formation of ordered u-related phases in alloys of composition Ti4Al3Nb, Acta Metall. Mater. 38 (1990) 931e943. [10] Z.W. Huang, Thermal stability of Tie44Ale4Nbe4Zre0.2Sie1B alloy, Intermetallics 42 (2013) 170e179. [11] L. Song, X.J. Xu, L. You, Y.F. Liang, J.P. Lin, Phase transformation and decomposition mechanisms of the bo(u) phase in cast high Nb containing TiAl alloy, J. Alloy. Compd. 616 (2014) 483e491. [12] T.T. Cheng, M.H. Loretto, The decomposition of the beta phase in Tie44Ale8Nb and Tie44Ale4Nbe4Zre0.2Si alloys, Acta Mater. 46 (1998) 4801e4819. € berl, E. Schwaighofer, Z. Zhang, H. Clemens, [13] M. Schloffer, B. Rashkova, T. Scho S. Mayer, Evolution of the uo phase in a b-stabilized multi-phase TiAl alloy and its effect on hardness, Acta Mater. 64 (2014) 241e252. [14] D. Daloz, U. Hecht, J. Zollinger, H. Combeau, A. Hazotte, M. Zalo znik, Microsegregation, macrosegregation and related phase transformations in TiAl alloys, Intermetallics 19 (2011) 749e756. [15] G. Shao, P. Tsakiropoulos, Prediction of u phase formation in TieAleX alloys, Mater. Sci. Eng. A 329e331 (2002) 914e919. [16] A. Stark, M. Oehring, F. Pyczak, A. Schreyer, Situ observation of various phase transformation paths in Nb-Rich TiAl alloys during quenching with different rates, Adv. Eng. Mater. 13 (2011) 700e704. [17] C.W. Dawson, S.L. Sass, The as-quenched form of the omega phase in Zr-Nb alloys, Metall. Trans. 1 (1970) 2225e2233. [18] R. Tewari, P. Mukhopadhyay, S. Banerjee, L.A. Bendersky, Evolution of ordered u phases in (Zr3Al)eNb alloys, Acta Mater. 47 (1999) 1307e1323. [19] T.S. Kuan, S.L. Sass, The structure of a linear omega-like vacancy defect in ZrNb B.C.C. solid solutions, Acta Metall. 24 (1976) 1053e1059. [20] S.K. Sikka, Y.K. Vohra, R. Chidambaram, Omega phase in materials, Prog. Mater. Sci. 27 (1982) 245e310. [21] S.G. Song, G.T. Gray III, Microscopic and crystallographic aspects of retained omega phase in shock-loaded zirconium and its formation mechanism, Philos. Mag. A 71 (1995) 275e290. [22] L. Song, X. Xu, L. You, Y. Liang, J. Lin, Ordered u phase transformations in Ti45Al-8.5Nb-0.2B alloy, Intermetallics 65 (2015) 22e28. [23] R. Kainuma, M. Palm, G. Inden, Solid-phase equilibria in the Ti-rich part of the TiAl system, Intermetallics 2 (1994) 321e332. [24] B. Shao, Y. Zong, D. Wen, Y. Tian, D. Shan, Investigation of the phase transformations in Ti22Al25Nb alloy, Mater. Charact. 114 (2016) 75e78. [25] D. Ping, Review on u phase in body-centered cubic metals and alloys, Acta Metall. Sin. Engl. Lett. 27 (2014) 1e11. [26] E. Sukedai, D. Yoshimitsu, H. Matsumoto, H. Hashimoto, M. Kiritani, b to u phase transformation due to aging in a TieMo alloy deformed in impact compression, Mater. Sci. Eng. A 350 (2003) 133e138. [27] E. Schwaighofer, H. Clemens, S. Mayer, J. Lindemann, J. Klose, W. Smarsly, V. Güther, Microstructural design and mechanical properties of a cast and heat-treated intermetallic multi-phase g-TiAl based alloy, Intermetallics 44
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