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Atomic-scale observations of B2 → ω-related phases transition in high-Nb containing TiAl alloy
Materials Characterization 130 (2017) 135–138
Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Atomic-scale observations of B2 → ω-related phases transition in high-Nb containing TiAl alloy
MARK
Xuyang Wang, Jieren Yang⁎, Keren Zhang, Rui Hu, Lin Song, 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
Keywords: TiAl alloy Transmission electron microscopy (TEM) Diffusion Phase transformation Omega phase
The nucleation of ω-related phases within B2-phase region in Ti-45Al-8.5Nb-(W, B, Y) (at.%) alloy were investigated by transmission electron microscopy (TEM) and high resolution electron microscopy (HREM). At 850 °C, the metastable B2 phase decomposes to ω particles quickly. Experimental evidence from high-resolution images revealed that first the atoms redistribute in the B2 structure, and then the original (111)B2 planes collapse during B2 → ω″, which represents a “diffusion-collapse” pathway. The further collapse of (0001)ω planes in [0001]ω zone, that is ω″ → ωo, is a sequential and uneven process rather than an instantaneous transition, which may be attributed to that an uneven structure transition could decrease the energy barrier.
1. Introduction β/B2 phase widely exist in high temperature TiAl alloys (due to the addition of β-stabilizers) that worked beyond 800 °C [1,2], which can improve the hot workability at high temperature distinctly [3,4]. It is noted that the β/B2 phase is metastable and would decompose to trigonal ω' and ω″ phase (P3m1) , hexagonal B82-ω phase (P63/mmc) after cooling from high temperature or during isothermal treatment in intermediate temperature range [5,6]. The room temperature ductility and fatigue strength would deteriorate with the precipitation of the hexagonal ωo phase in the B2 phase [7–9]. In titanium alloys, ω forms from β lattice by the collapse of two {111}β planes among every three {111}β planes [10–12]. In contrast, β and ω phases exist in TiAl alloys as the ordered structures, and the chemical order among different Wykoff sites is inevitable in the process of B2 → ωo transition, which is different to that in titanium alloys. Previous works indicated that, in cooling process of TiAl alloys, the transition of B2 → ωo occurs via the collapse and diffusion mechanism [13–15], that is, every second and third (111)B2 layers of the B2 phase “collapse” toward each other along [111]B2 (Fig. 1a) and form (0001)h layers in the trigonal or hexagonal structure [10]. It was believed that if the product directly inherits the chemical order of B2 phase, which is a pure displacive transition, it is referred to as ω′ transition (Fig. 1b). ω′ phase is inherently unstable, and the ω″ phase will form via chemical order, as seen in Fig. 1c. It should be emphasized that the double collapsed layers in ω' and ω″ phases are rumpled and the ω-collapse is not complete at this moment. Further, the (0001)h collapse and the atom reconstruction in ω″ structure will continue [16]. Finally, the formation ⁎
Corresponding author. E-mail address: [email protected] (J. Yang).
http://dx.doi.org/10.1016/j.matchar.2017.06.003 Received 14 November 2016; Received in revised form 15 May 2017; Accepted 1 June 2017 Available online 01 June 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.
of hexagonal B82-ω phase (which is termed as ωo) illustrated in Fig. 1d is complete. However, there is less known about the formation mechanism of ωo phase in B2 region at intermediate temperature range in aging process. In particular, conclusive experimental evidence for the postulated mechanism is still lacking. Two mechanisms about the formation of ω″ phase were proposed: the lattice collapse promotes the atoms redistribution (collapse-diffusion mechanism) or the redistribution triggers the collapse (diffusion-collapse mechanism) [17]. Which mechanism dominates the phase transition of B2 → ω″ has not been confirmed. In the present study, the B2 → ωo transition in Ti-45Al-8.5Nb-(W, B, Y) (at.%) was investigated. The processes of atom redistribution and lattice collapse in B2 → ωo transition were characterized using transmission electron microscopy (TEM) combined with high resolution electron microscopy (HREM). The transition mechanism was discussed on the basis of experimental evidence. 2. Materials and Experimental Procedures The nominal composition of Ti-45Al-8.5Nb-0.2 W-0.2B-0.02Y (at. %) ingot was produced by a plasma cold hearth melting furnace. Samples in size of 8 mm × 8 mm × 8 mm were wire-cut from the ingot. A pre-heat treatment, annealing at 1100 °C for 1 h followed by water quenching, in order to eliminate the ωo phase in the B2 matrix. Then these samples were annealed at 850 °C for 20 min in a preheated resistance furnace and then were quenched in water. To eliminate the influence of the high-temperature oxidation on the transition of B2 → ω, in this study, the TME specimens were cut from the center of these
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Fig. 1. Schematic illustrations of crystal structures and Wyckoff positions of (a) B2, (b) ω′, (c) ω″ and (d) ωo.
heat-treated samples, further grinded to 50 μm in thickness and finally thinned using twin-jet electro-polisher in an electrolyte of 5 vol% perchloric acid, 30 vol% butan-1-ol and 65 vol% methanol with operating at 30 V and − 30 °C. TEM observations were performed on Tecnai G2 F30 operating at 300 kV.
the [110]B2 zone axis, indicating that four variants exist in B2 phase region including two < 1120>ωo and two < 1102>ωo zone axes, as also revealed in Ref. [16]. After quenching from 1100 °C, no ωo phase particles can be found in the B2 area (Fig. 2b). However, from the selected area electron diffraction pattern along the [110]B2 zone axis, in addition to the {110} B2 , {001} B2 and {111}B2, the diffuse maxima related to ω phase (marked by blue circles) are visible, which indicates that some ω-related structures exist in the B2 matrix [10,18]. These ω structures with the habit plane {0001}ω//{111}B2 in a range of several nanometers distribute within the B2 matrix uniformly [5]. After annealed at 850 °C for 20 min, granular ωo particles precipitate along the B2 grain boundaries, as presented in Fig. 2c. Fig. 2d indicates that the average diameter of the precipitated ω particles is
3. Results and Discussion 3.1. The Reprecipitation of ω Phase Fig.2a shows the microstructure of the as-cast Ti-45Al-8.5Nb-(W, B, Y) alloy. A large number of ωo particles with the size of 0.3–1.0 μm distribute within B2 matrix. The inset in the corner of Fig. 2a shows the selected area diffraction pattern (SADP) of this area that viewed along
Fig. 2. TEM observation of B2(ω) region: (a) as-cast Ti-45Al-8.5Nb-(W, B, Y) alloy; (b) water-quenched sample; (c) annealing at 850 °C for 20 min; (d) magnified image of the blue dotted box in Fig. 2c. (For interpretation of the references to colour in this figure, the reader is referred to the web version of this article.)
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about 80 nm. The growth rate of ω particles is high, which is because the composition of Nb-segregation areas in this alloy is very close to that of ωo phase [16] and the free energy of ωo (Gωo) phase is much lower than that of B2 phase (GB2) at 850 °C [17]. It is worthy of noting that no ωo reflections arising in the SADP inserted in the top-right corner of Fig. 2c, which suggests that no ω-related phases form in the center region of B2 phase. It can be speculated that the difference between Gωo and GB2 is smaller than the energy barrier of B2 → ωo transition at 850 °C. The higher defect density can provide additional nucleation energy for the ωo phase [19], so the ωo particles precipitate first at the grain boundaries at 850 °C. Results in [20] revealed that the point defects (interstitial atoms) could induce the nucleation of the ω phase from β phase. In this study, interstitial atoms are likely to concentrate at grain boundaries, acting as the point defects. These point defects would induce local distortions in B2 matrix that containing the ωo structure fluctuations domains. On the other hand, the stress arising from quenching easily concentrates at the grain boundaries, which could promote the B2 → ωo transition when annealing at 850 °C [21]. Therefore, the B2 grain boundaries are the preferential area for the nucleation of ωo phase.
disordered mixture of A1 and Nb. Moreover, the contrast in (0001)ω layers (correspond to layers 1 and 7 in Fig. 1) gradually becomes brighter from left to right in Fig. 3c, compared with that Nb content in these layers gradually increases during the transition (Fig. 1c), indicating that the Nb atoms exhibits as a high brightness in this IFFT graph. It is generally believed that two main steps (the lattice collapse and the atoms diffusion) are involved in the B2 → ωo transition. However, it is not clear which step occurs first, that is, whether the diffusion induces the collapse or the opposite path is still in debate. From Fig. 3c, on the one hand, the Ti atoms (where the blue arrows point) are on the yellow dotted line and no displacement/collapse take places, so there is no partial collapse of the lattice during this phase transition stage; On the other hand, it can be distinctly observed that the brightness of the 1b sites (occupied by Al and Nb atoms) gradually becomes darker from the left side to the right side (white arrows point), which indicates that the Nb atoms gradually transfer out of 1b sites. A reasonable speculation can be made: Nb atoms diffuse from the 1b sites in the second and sixth layers to the sites in the first, fourth and seventh layers, and Ti atoms are diffusing from the first and seventh layers to the fourth layer. Similarly, the calculation in Ref. [17] indicated that during the B2 → ω″ transition, the atoms exchange first and then the lattice collapses to an intermediate structure, unlike Ti-V-Cu and Ti-Mo alloys in which there is no such short distance atoms reconstruction during β → ω [23,24]. This transition occurs without an energy barrier except the activation energy of the atoms rearrangement. When the B2 phase transfers to the ω″ type structure, the layers 2 and 3 as well as 5 and 6 of the unit cell begin to move toward each other, and reach an intermediate stable structure with the minimum free energy. Therefore, it can be confirmed that the mechanism of the B2 → ω″ transition is the “diffusion-collapse” controlled pathway. Along the A-B line marked in Fig. 3a, the intensity of an array of atoms was plotted according to the contrast and shown in Fig. 3d. Notably, the seventh intensity peak appears earlier about 0.156 nm. In order to explain this phenomenon, to get started, it can be assumed that this is because of the mismatch of crystal lattice in the phase transition. If there is no crystal lattice mismatch, the interplanar spacing of (1010)ω planes is equal to three times that of (112)B2 . In this research, the3d (112)B2 = 0.393 nm and the interplanar spacing of (1010)ω is 0.395 nm based on the HRTEM image. Therefore, the size of precipitated ω particle can be roughly estimated as:
3.2. Process of B2 → ω″ Transition To understand the process of B2 → ωo phase transition, the boundary between B2 and ω phase was characterized. Fig. 3a shows the inverse fast Fourier transform (IFFT) graph of the B2 → ω″ area along the [110]B2 zone axis. The left and the right side represent the B2 and ω phase respectively. Importantly, the middle zone reveals the transition of B2 → ω. Fig. 3b shows the magnified view of the B2 phase. Lattice fringes with the spacing of 0.322 and 0.228 nm corresponding to the (001)B2 and the (110)B2 plane of the matrix B2 phase structure, respectively. The lattice parameter of B2 in the left area of Fig. 3a is a = 0.322 nm and that of ω-related phase in the right area of Fig. 3a are a = b = 0.456 nm, c = 0.565 nm. It can be seen that there are two different levels of brightness along the (110)B2 planes. In the present work, Fig. 3a was obtained at a perfect Scherzer defocus state, the atoms attributed to a same element are represented by a specific contrast [22]. As presented in Fig. 1, the structure of layers 1 and 7 in ω phase completely inherit that of (111)B2 layers which are occupied by Ti atoms. Considering that the (0001)h layers in ω phase are connected with the (111)B2 layers that occupied by bright atoms, it can be deduced from Fig. 3a that the bright sites are Ti and the dark sites are by a
Fig. 3. (a) IFFT graph of the HRTEM graph in B2 → ω transition area; (b) and (c) are the magnified image of area I and II, respectively; (d) the line profile taken from the line AB marked in Fig. 3a. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
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Fig. 4. (a) IFFT graph of the HRTEM graph in ω″ → ωo transition area, (b) three parallel intensity profiles along the [0001]ω direction from the rows marked as A, B and C in (a). The three intensity profiles start from the same (0001)ω plane.
dω = 2⋅L⋅(d (1010)ω − 3d (112)B2)−1⋅d (1010)ω
of ωo particles is very fast. (2) The first stage of B2 → ωo transition is B2 → ω″, in which the atoms redistribute first and then the lattice partially collapses. It is a “diffusion-collapse” controlled pathway. (3) The second stage of the phase transition is ω″ → ωo. During this stage, the collapse process among different layers is out of sync. The collapse of (0001)ω along the [0001]ω is sequential.
(1)
Where, dω, L, d (1010)ω and d (112)B2 are the size of ω particle, the total atomic displacement (as marked in Fig.3d), the interplanar spacings of the (1010)ω and the (112)B2 , respectively. Based on the Formula (1), the dω in this work was calculated and the value is 87.9 nm which agrees well with the result in Fig. 2d. So this phenomenon can be responsible for the mismatch of crystal lattic. 3.3. The Early Stage of ω″ → ωo Transition
Acknowledgements
For the ω″ → ωo transition, the atomic structure transition was also identified by HRTEM. Further, the IFFT graph is obtained and presented in Fig. 4a. The layers 2 and 3 (5 and 6) collapse toward each other and form a double-atom layer, which is similar to the typical ω″ structure (displayed in Fig. 1c). To evaluate the collapse along [0001]ω, three representative intensity profiles along [0001]ω (marked as A, B and C in Fig. 4a) beginning at a same (0001)ω plane were plotted. The collapsed displacements (l) of layers 2–3 (l23) and 5–6 (l56) are measured based on Fig. 4a. It should be noted that the l23 is evidently different with the l56. l56 is a constant that is 0.07 nm, which indicates that layers 5–6 do not collapse in the early stage of ω″ → ωo transition. In contrast, l23 gradually decreases from 0.05 nm in the left to 0.02 nm in the right, as revealed in Fig. 4b. So it is seems that the collapses of the layers 2–3 and the layers 5–6 are two unsynchronized processes. Therefore, the collapse of (0001)ω planes along [0001]ω zone axis is a sequential (similar to the near-β titanium alloys [12]) and uneven process. This may be attributed to the high energy barrier during ω″ → ωo transition [17], this uneven structure transition could decrease the energy barrier. It is worth to note that the stage of ω″ → ωo may be a mixed-mode transition, which is similar with the study in [25]. For further study, some thermodynamic calculations will be performed to understand the structure transition in a quantitative way. The change of the lattice parameters of the ω phase during this stage is difficult to be detected owing to the resolution limitation. Here, only the average lattice parameters in this area are given: a = b = 0.458 nm, c = 0.570 nm.
This work was supported by the National Natural Science Foundation of China (No. 51401168), the Fundamental Research Funds for the Central Universities (No.3102014JCQ01026) and the Shaanxi Provincial Natural Science Foundation (No. 2016JQ5087). References [1] T. Voisin, J.P. Monchoux, M. Hantcherli, S. Mayer, H. Clemens, A. Couret, Acta Mater. 73 (2014) 107–115. [2] G. Yang, H. Kou, J. Yang, J. Li, H. Fu, Acta Mater. 112 (2016) 121–131. [3] A.V. Kartavykh, E.A. Asnis, N.V. Piskun, I.I. Statkevich, M.V. Gorshenkov, A.V. Korotitskiy, Mater. Lett. 162 (2016) 180–184. [4] Y. Chen, H. Niu, F. Kong, S. Xiao, Intermetallics 19 (2011) 1405–1410. [5] L. Song, X.J. Xu, J. Sun, J.P. Lin, Mater. Charact. 93 (2014) 62–67. [6] L. Song, L.Q. Zhang, X.J. Xu, J. Sun, J.P. Lin, Scr. Mater. 68 (2013) 929–932. [7] E. Schwaighofer, H. Clemens, S. Mayer, J. Lindemann, J. Klose, W. Smarsly, V. Güther, Intermetallics 44 (2014) 128–140. [8] M. Schloffer, B. Rashkova, T. Schöberl, E. Schwaighofer, Z. Zhang, H. Clemens, S. Mayer, Acta Mater. 64 (2014) 241–252. [9] Z.W. Huang, Acta Mater. 56 (2008) 1689–1700. [10] D. De Fontaine, N.E. Paton, J.C. Williams, Acta Metall. 19 (1971) 1153–1162. [11] M.J. Lai, C.C. Tasan, J. Zhang, B. Grabowski, L.F. Huang, D. Raabe, Acta Mater. 92 (2015) 55–63. [12] T. Li, D. Kent, G. Sha, L.T. Stephenson, A.V. Ceguerra, S.P. Ringer, M.S. Dargusch, J.M. Cairney, Acta Mater. 106 (2016) 353–366. [13] A. Stark, M. Oehring, F. Pyczak, A. Schreyer, Adv. Eng. Mater. 13 (2011) 700–704. [14] S.K. Sikka, Y.K. Vohra, R. Chidambaram, Prog. Mater. Sci. 27 (1982) 245–310. [15] G. Shao, P. Tsakiropoulos, Acta Mater. 48 (2000) 3671–3685. [16] L.A. Bendersky, W.J. Boettinger, B.P. Burton, F.S. Biancaniello, C.B. Shoemaker, Acta Metall. Mater. 38 (1990) 931–943. [17] Q.M. Hu, L. Vitos, R. Yang, Phys. Rev. B 90 (2014) (054109). [18] G. Shao, P. Tsakiropoulos, Mater. Sci. Eng. A 329–331 (2002) 914–919. [19] B. Shao, Y. Zong, D. Wen, Y. Tian, D. Shan, Mater. Charact. 114 (2016) 75–78. [20] H. Dosch, A.V. Schwerin, J. Peisl, Phys. Rev. B 34 (1986) 1654–1661. [21] B. Cao, J. Yang, X. Wang, R. Hu, Adv. Eng. Mater. (2017). [22] H. Zhang, H. Kou, J. Yang, J. Wang, J. Li, X. Xue, Intermetallics 60 (2015) 13–18. [23] A. Devaraj, S. Nag, R. Srinivasan, R.E.A. Williams, S. Banerjee, R. Banerjee, H.L. Fraser, Acta Mater. 60 (2012) 596–609. [24] H.P. Ng, A. Devaraj, S. Nag, C.J. Bettles, M. Gibson, H.L. Fraser, B.C. Muddle, R. Banerjee, Acta Mater. 59 (2011) 2981–2991. [25] S. Nag, A. Devaraj, R. Srinivasan, R.E.A. Williams, N. Gupta, G.B. Viswanathan, J.S. Tiley, S. Banerjee, Phys. Rev. Lett. 106 (2011) 245701.
4. Conclusions In this work, the process of B2 → ωo phase transformation in Ti45Al-8.5Nb-0.2 W-0.2B-0.02Y alloy was investigated by TEM. Following are the main results: (1) At 850 °C, the B2 phase is thermodynamic unstable, granular ωo particles precipitate along the B2 grain boundaries and the growth
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Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Atomic-scale observations of B2 → ω-related phases transition in high-Nb containing TiAl alloy
MARK
Xuyang Wang, Jieren Yang⁎, Keren Zhang, Rui Hu, Lin Song, 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
Keywords: TiAl alloy Transmission electron microscopy (TEM) Diffusion Phase transformation Omega phase
The nucleation of ω-related phases within B2-phase region in Ti-45Al-8.5Nb-(W, B, Y) (at.%) alloy were investigated by transmission electron microscopy (TEM) and high resolution electron microscopy (HREM). At 850 °C, the metastable B2 phase decomposes to ω particles quickly. Experimental evidence from high-resolution images revealed that first the atoms redistribute in the B2 structure, and then the original (111)B2 planes collapse during B2 → ω″, which represents a “diffusion-collapse” pathway. The further collapse of (0001)ω planes in [0001]ω zone, that is ω″ → ωo, is a sequential and uneven process rather than an instantaneous transition, which may be attributed to that an uneven structure transition could decrease the energy barrier.
1. Introduction β/B2 phase widely exist in high temperature TiAl alloys (due to the addition of β-stabilizers) that worked beyond 800 °C [1,2], which can improve the hot workability at high temperature distinctly [3,4]. It is noted that the β/B2 phase is metastable and would decompose to trigonal ω' and ω″ phase (P3m1) , hexagonal B82-ω phase (P63/mmc) after cooling from high temperature or during isothermal treatment in intermediate temperature range [5,6]. The room temperature ductility and fatigue strength would deteriorate with the precipitation of the hexagonal ωo phase in the B2 phase [7–9]. In titanium alloys, ω forms from β lattice by the collapse of two {111}β planes among every three {111}β planes [10–12]. In contrast, β and ω phases exist in TiAl alloys as the ordered structures, and the chemical order among different Wykoff sites is inevitable in the process of B2 → ωo transition, which is different to that in titanium alloys. Previous works indicated that, in cooling process of TiAl alloys, the transition of B2 → ωo occurs via the collapse and diffusion mechanism [13–15], that is, every second and third (111)B2 layers of the B2 phase “collapse” toward each other along [111]B2 (Fig. 1a) and form (0001)h layers in the trigonal or hexagonal structure [10]. It was believed that if the product directly inherits the chemical order of B2 phase, which is a pure displacive transition, it is referred to as ω′ transition (Fig. 1b). ω′ phase is inherently unstable, and the ω″ phase will form via chemical order, as seen in Fig. 1c. It should be emphasized that the double collapsed layers in ω' and ω″ phases are rumpled and the ω-collapse is not complete at this moment. Further, the (0001)h collapse and the atom reconstruction in ω″ structure will continue [16]. Finally, the formation ⁎
Corresponding author. E-mail address: [email protected] (J. Yang).
http://dx.doi.org/10.1016/j.matchar.2017.06.003 Received 14 November 2016; Received in revised form 15 May 2017; Accepted 1 June 2017 Available online 01 June 2017 1044-5803/ © 2017 Elsevier Inc. All rights reserved.
of hexagonal B82-ω phase (which is termed as ωo) illustrated in Fig. 1d is complete. However, there is less known about the formation mechanism of ωo phase in B2 region at intermediate temperature range in aging process. In particular, conclusive experimental evidence for the postulated mechanism is still lacking. Two mechanisms about the formation of ω″ phase were proposed: the lattice collapse promotes the atoms redistribution (collapse-diffusion mechanism) or the redistribution triggers the collapse (diffusion-collapse mechanism) [17]. Which mechanism dominates the phase transition of B2 → ω″ has not been confirmed. In the present study, the B2 → ωo transition in Ti-45Al-8.5Nb-(W, B, Y) (at.%) was investigated. The processes of atom redistribution and lattice collapse in B2 → ωo transition were characterized using transmission electron microscopy (TEM) combined with high resolution electron microscopy (HREM). The transition mechanism was discussed on the basis of experimental evidence. 2. Materials and Experimental Procedures The nominal composition of Ti-45Al-8.5Nb-0.2 W-0.2B-0.02Y (at. %) ingot was produced by a plasma cold hearth melting furnace. Samples in size of 8 mm × 8 mm × 8 mm were wire-cut from the ingot. A pre-heat treatment, annealing at 1100 °C for 1 h followed by water quenching, in order to eliminate the ωo phase in the B2 matrix. Then these samples were annealed at 850 °C for 20 min in a preheated resistance furnace and then were quenched in water. To eliminate the influence of the high-temperature oxidation on the transition of B2 → ω, in this study, the TME specimens were cut from the center of these
Materials Characterization 130 (2017) 135–138
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Fig. 1. Schematic illustrations of crystal structures and Wyckoff positions of (a) B2, (b) ω′, (c) ω″ and (d) ωo.
heat-treated samples, further grinded to 50 μm in thickness and finally thinned using twin-jet electro-polisher in an electrolyte of 5 vol% perchloric acid, 30 vol% butan-1-ol and 65 vol% methanol with operating at 30 V and − 30 °C. TEM observations were performed on Tecnai G2 F30 operating at 300 kV.
the [110]B2 zone axis, indicating that four variants exist in B2 phase region including two < 1120>ωo and two < 1102>ωo zone axes, as also revealed in Ref. [16]. After quenching from 1100 °C, no ωo phase particles can be found in the B2 area (Fig. 2b). However, from the selected area electron diffraction pattern along the [110]B2 zone axis, in addition to the {110} B2 , {001} B2 and {111}B2, the diffuse maxima related to ω phase (marked by blue circles) are visible, which indicates that some ω-related structures exist in the B2 matrix [10,18]. These ω structures with the habit plane {0001}ω//{111}B2 in a range of several nanometers distribute within the B2 matrix uniformly [5]. After annealed at 850 °C for 20 min, granular ωo particles precipitate along the B2 grain boundaries, as presented in Fig. 2c. Fig. 2d indicates that the average diameter of the precipitated ω particles is
3. Results and Discussion 3.1. The Reprecipitation of ω Phase Fig.2a shows the microstructure of the as-cast Ti-45Al-8.5Nb-(W, B, Y) alloy. A large number of ωo particles with the size of 0.3–1.0 μm distribute within B2 matrix. The inset in the corner of Fig. 2a shows the selected area diffraction pattern (SADP) of this area that viewed along
Fig. 2. TEM observation of B2(ω) region: (a) as-cast Ti-45Al-8.5Nb-(W, B, Y) alloy; (b) water-quenched sample; (c) annealing at 850 °C for 20 min; (d) magnified image of the blue dotted box in Fig. 2c. (For interpretation of the references to colour in this figure, the reader is referred to the web version of this article.)
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about 80 nm. The growth rate of ω particles is high, which is because the composition of Nb-segregation areas in this alloy is very close to that of ωo phase [16] and the free energy of ωo (Gωo) phase is much lower than that of B2 phase (GB2) at 850 °C [17]. It is worthy of noting that no ωo reflections arising in the SADP inserted in the top-right corner of Fig. 2c, which suggests that no ω-related phases form in the center region of B2 phase. It can be speculated that the difference between Gωo and GB2 is smaller than the energy barrier of B2 → ωo transition at 850 °C. The higher defect density can provide additional nucleation energy for the ωo phase [19], so the ωo particles precipitate first at the grain boundaries at 850 °C. Results in [20] revealed that the point defects (interstitial atoms) could induce the nucleation of the ω phase from β phase. In this study, interstitial atoms are likely to concentrate at grain boundaries, acting as the point defects. These point defects would induce local distortions in B2 matrix that containing the ωo structure fluctuations domains. On the other hand, the stress arising from quenching easily concentrates at the grain boundaries, which could promote the B2 → ωo transition when annealing at 850 °C [21]. Therefore, the B2 grain boundaries are the preferential area for the nucleation of ωo phase.
disordered mixture of A1 and Nb. Moreover, the contrast in (0001)ω layers (correspond to layers 1 and 7 in Fig. 1) gradually becomes brighter from left to right in Fig. 3c, compared with that Nb content in these layers gradually increases during the transition (Fig. 1c), indicating that the Nb atoms exhibits as a high brightness in this IFFT graph. It is generally believed that two main steps (the lattice collapse and the atoms diffusion) are involved in the B2 → ωo transition. However, it is not clear which step occurs first, that is, whether the diffusion induces the collapse or the opposite path is still in debate. From Fig. 3c, on the one hand, the Ti atoms (where the blue arrows point) are on the yellow dotted line and no displacement/collapse take places, so there is no partial collapse of the lattice during this phase transition stage; On the other hand, it can be distinctly observed that the brightness of the 1b sites (occupied by Al and Nb atoms) gradually becomes darker from the left side to the right side (white arrows point), which indicates that the Nb atoms gradually transfer out of 1b sites. A reasonable speculation can be made: Nb atoms diffuse from the 1b sites in the second and sixth layers to the sites in the first, fourth and seventh layers, and Ti atoms are diffusing from the first and seventh layers to the fourth layer. Similarly, the calculation in Ref. [17] indicated that during the B2 → ω″ transition, the atoms exchange first and then the lattice collapses to an intermediate structure, unlike Ti-V-Cu and Ti-Mo alloys in which there is no such short distance atoms reconstruction during β → ω [23,24]. This transition occurs without an energy barrier except the activation energy of the atoms rearrangement. When the B2 phase transfers to the ω″ type structure, the layers 2 and 3 as well as 5 and 6 of the unit cell begin to move toward each other, and reach an intermediate stable structure with the minimum free energy. Therefore, it can be confirmed that the mechanism of the B2 → ω″ transition is the “diffusion-collapse” controlled pathway. Along the A-B line marked in Fig. 3a, the intensity of an array of atoms was plotted according to the contrast and shown in Fig. 3d. Notably, the seventh intensity peak appears earlier about 0.156 nm. In order to explain this phenomenon, to get started, it can be assumed that this is because of the mismatch of crystal lattice in the phase transition. If there is no crystal lattice mismatch, the interplanar spacing of (1010)ω planes is equal to three times that of (112)B2 . In this research, the3d (112)B2 = 0.393 nm and the interplanar spacing of (1010)ω is 0.395 nm based on the HRTEM image. Therefore, the size of precipitated ω particle can be roughly estimated as:
3.2. Process of B2 → ω″ Transition To understand the process of B2 → ωo phase transition, the boundary between B2 and ω phase was characterized. Fig. 3a shows the inverse fast Fourier transform (IFFT) graph of the B2 → ω″ area along the [110]B2 zone axis. The left and the right side represent the B2 and ω phase respectively. Importantly, the middle zone reveals the transition of B2 → ω. Fig. 3b shows the magnified view of the B2 phase. Lattice fringes with the spacing of 0.322 and 0.228 nm corresponding to the (001)B2 and the (110)B2 plane of the matrix B2 phase structure, respectively. The lattice parameter of B2 in the left area of Fig. 3a is a = 0.322 nm and that of ω-related phase in the right area of Fig. 3a are a = b = 0.456 nm, c = 0.565 nm. It can be seen that there are two different levels of brightness along the (110)B2 planes. In the present work, Fig. 3a was obtained at a perfect Scherzer defocus state, the atoms attributed to a same element are represented by a specific contrast [22]. As presented in Fig. 1, the structure of layers 1 and 7 in ω phase completely inherit that of (111)B2 layers which are occupied by Ti atoms. Considering that the (0001)h layers in ω phase are connected with the (111)B2 layers that occupied by bright atoms, it can be deduced from Fig. 3a that the bright sites are Ti and the dark sites are by a
Fig. 3. (a) IFFT graph of the HRTEM graph in B2 → ω transition area; (b) and (c) are the magnified image of area I and II, respectively; (d) the line profile taken from the line AB marked in Fig. 3a. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)
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Fig. 4. (a) IFFT graph of the HRTEM graph in ω″ → ωo transition area, (b) three parallel intensity profiles along the [0001]ω direction from the rows marked as A, B and C in (a). The three intensity profiles start from the same (0001)ω plane.
dω = 2⋅L⋅(d (1010)ω − 3d (112)B2)−1⋅d (1010)ω
of ωo particles is very fast. (2) The first stage of B2 → ωo transition is B2 → ω″, in which the atoms redistribute first and then the lattice partially collapses. It is a “diffusion-collapse” controlled pathway. (3) The second stage of the phase transition is ω″ → ωo. During this stage, the collapse process among different layers is out of sync. The collapse of (0001)ω along the [0001]ω is sequential.
(1)
Where, dω, L, d (1010)ω and d (112)B2 are the size of ω particle, the total atomic displacement (as marked in Fig.3d), the interplanar spacings of the (1010)ω and the (112)B2 , respectively. Based on the Formula (1), the dω in this work was calculated and the value is 87.9 nm which agrees well with the result in Fig. 2d. So this phenomenon can be responsible for the mismatch of crystal lattic. 3.3. The Early Stage of ω″ → ωo Transition
Acknowledgements
For the ω″ → ωo transition, the atomic structure transition was also identified by HRTEM. Further, the IFFT graph is obtained and presented in Fig. 4a. The layers 2 and 3 (5 and 6) collapse toward each other and form a double-atom layer, which is similar to the typical ω″ structure (displayed in Fig. 1c). To evaluate the collapse along [0001]ω, three representative intensity profiles along [0001]ω (marked as A, B and C in Fig. 4a) beginning at a same (0001)ω plane were plotted. The collapsed displacements (l) of layers 2–3 (l23) and 5–6 (l56) are measured based on Fig. 4a. It should be noted that the l23 is evidently different with the l56. l56 is a constant that is 0.07 nm, which indicates that layers 5–6 do not collapse in the early stage of ω″ → ωo transition. In contrast, l23 gradually decreases from 0.05 nm in the left to 0.02 nm in the right, as revealed in Fig. 4b. So it is seems that the collapses of the layers 2–3 and the layers 5–6 are two unsynchronized processes. Therefore, the collapse of (0001)ω planes along [0001]ω zone axis is a sequential (similar to the near-β titanium alloys [12]) and uneven process. This may be attributed to the high energy barrier during ω″ → ωo transition [17], this uneven structure transition could decrease the energy barrier. It is worth to note that the stage of ω″ → ωo may be a mixed-mode transition, which is similar with the study in [25]. For further study, some thermodynamic calculations will be performed to understand the structure transition in a quantitative way. The change of the lattice parameters of the ω phase during this stage is difficult to be detected owing to the resolution limitation. Here, only the average lattice parameters in this area are given: a = b = 0.458 nm, c = 0.570 nm.
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4. Conclusions In this work, the process of B2 → ωo phase transformation in Ti45Al-8.5Nb-0.2 W-0.2B-0.02Y alloy was investigated by TEM. Following are the main results: (1) At 850 °C, the B2 phase is thermodynamic unstable, granular ωo particles precipitate along the B2 grain boundaries and the growth
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