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Precipitation behavior of the ωo phase in an annealed high Nb-TiAl alloy
Journal of Alloys and Compounds 701 (2017) 882e891
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Precipitation behavior of the uo phase in an annealed high Nb-TiAl alloy Teng Ye a, Lin Song b, **, Shubo Gao a, Yongfeng Liang a, Yanli Wang a, Junpin Lin a, * a b
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, Shaanxi, 710072, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 31 May 2016 Received in revised form 28 December 2016 Accepted 18 January 2017 Available online 20 January 2017
Ordered omega (uo) phases are considered to be equilibrium phases in high Nb-TiAl alloys at intermediate temperatures. However, most studies have focused on the evolution behaviors of these phases during annealing, while the precipitation behavior has not been systematically studied. In the present work, the precipitation behavior of the uo phase (B82 structure) in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y (at.%) alloy during annealing at 850 C was investigated using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After annealing at 1150 C for 2 h followed by water quenching, the uo phase in the as-cast alloy transformed to the bo phase. At the initial stage of annealing processing at 850 C, the uo phase re-precipitated at the g/bo boundaries and the size of the uo phase increased with the annealing time. Moreover, the g plates precipitated within the bo phase at this temperature became coarser and tended to globalize during annealing. In addition, the a2 and D88-u phases were also observed at the g/bo boundaries. Energy dispersive X-ray spectroscopy (EDS) and Scanning transmission electron microscopy (STEM) analysis showed that the D88-u phase is rich in W and Nb elements. The corresponding mechanisms of the transformation mentioned above were discussed. © 2017 Elsevier B.V. All rights reserved.
Keywords: Titanium aluminides Phase transformation TEM Ordered omega phases
1. Introduction High Nb-TiAl alloys have been considered to be potential materials for high-temperature applications due to their low density, high strength, good oxidation resistance and creep properties [1e3]. As reported in many studies, Nb is a b phase stabilizer that extends the b phase region to lower temperatures and higher Al contents in the phase diagram [4,5]. Thus, a certain amount of b phase can be retained in an as-cast ingot during the cooling process [5e7]. Although the retained b phase is expected to facilitate the hot workability of TiAl alloys [8e10], the ordered bo(B2 structure) phase transformed from the disordered bcc b(A2 structure) phase is considered to be detrimental to the ductility of TiAl alloys [5]. Recently, several studies have reported ordered u phase transitions in the bo phase regions in high Nb-TiAl alloys that involved chemical (replacive) and displacive order/disorder transitions [11,12]. Commonly, the ordered u phase in the B82 structure is
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (L. Song), [email protected] (J. Lin). http://dx.doi.org/10.1016/j.jallcom.2017.01.198 0925-8388/© 2017 Elsevier B.V. All rights reserved.
assigned as uo (space group: P63/mmc). Many reports have indicated that the uo phase in high Nb-TiAl alloys was an equilibrium phase at 700 C-900 C and that the observed orientation relationship (OR) between the uo and bo phases was [0001]uo//[111]bo; (1120)uo//(110)bo [13,14]. Huang et al. indicated that there is a three-phase region, namely, a2þgþbo(uo) (the ordered u formed within the bo phase) at 700 C in the Ti-44Al-8Nb-1B alloy, with further results showing that a2þg transformed into bo(uo) following thermal exposure at 700 C [15]. Meanwhile, g plates precipitated from the bo phase were also observed in Refs. [13,16e19]. Song et al. reported that the bo(uo)/g transformation can be accomplished by direct nucleation within the bo phase and growth of the pre-existing g phase in the Ti-45Al-8.5Nb(W, B, Y) alloy after annealing at 900 C [16]. Similar results were reported by Tang et al. in an alloy with the same composition after annealing at 700 C and 800 C [17]. It was indicated that due to the enhanced diffusion process, the g plates became coarser and globalized when the annealing temperature increased from 700 C to 800 C. Meanwhile, the D88-u phase (space group: P63/mcm) was also observed in a TiAl alloy with Nb and Zr additions formed through the occupation of the 2b positions by vacancies in the uo phase [11,20e22]. Tewari et al. proposed that Nb had a significant
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effect on the stabilization of vacancies during the formation of the D88 structure in the (Zr3Al)-Nb alloy [21]. Huang reported that the orientation relationship between the D88-u, uo and bo phases in the Ti-44Al-4Nb-4Zr-0.2Si-1B alloy was 〈0001〉D88-u//〈0001〉uo// 〈111〉bo and {1010}D88-u//{1120}uo//{110}bo and the structure transformation path in the retained b after annealing at 700 for 1000 h was b/bo/uo/D88-u [20]. Similar results were also reported by Song et al. in Ref. [16] where the vacancies produced by bo(uo)/g phase transformation acted as the main cause for D88-u formation and that D88-u could be dissolved into the g phase at an area far away from the g/bo(uo) boundary; it was suggested that the formation of D88-u was significantly affected by the vacancy concentration. These results suggest that the phase transformation that occurs in high Nb-TiAl alloys is strongly dependent on the local composition. However, investigations of the phase transformation mechanism related to composition are rare. Therefore, the formation mechanism of the D88-u phase and decomposition of the bo(uo) phase during annealing should be investigated more carefully, and a complete understanding of the phase transformation in high Nb-TiAl alloys at service temperatures is meaningful for the application of these alloys. In the present work, the phase transformation processes in the bo phase were investigated during annealing at 850 C and the corresponding transformation mechanisms were discussed in terms of composition analysis. 2. Experiments An ingot of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy used in this study was prepared using a plasma cold hearth melting furnace. Specimens with sizes of 15 10 10 mm were cut from the center of the ingot by electric-discharge machining. To reveal the precipitation behavior of the alloy at intermediate temperature, the samples were first annealed at 1150 C for 2 h, followed by water quenching (designed specimen A) to remove the ordered omega phases in the as-cast alloy [23], and then, some specimens were annealed at 850 C for 30min, 1 h, 3 h and 10 h, followed by water quenching (designed as specimens B, C, D, E, respectively) as listed in Table 1. The microstructures after the heat treatments were examined using a Zeiss Supra 55 scanning electron microscope (SEM) in backscatter mode operating at 15 kV. Discs with diameters of 3 mm and thicknesses of 0.3 mm were cut from the specimens by spark discharge cutting and were then mechanically ground to a thickness of 50e60 mm. Thin foils used for TEM observation were prepared by twin-jet electro-polishing in a solution of 65 vol% methanol, 30 vol% butanol, and 5 vol% perchloric acid at 30 V and 30 C. TEM analysis was conducted using a Tecnai G2 F30 field emission transmission electron microscope operating at 300 kV. The compositions and the element distributions of the phases were obtained by an energy dispersive X-ray spectroscopy (EDS) on TEM and Scanning transmission electron microscopy (STEM), respectively. The composition measurements were processed in the adjacent region to ensure the uniform thickness of the specimen. Each parameter was obtained as an average of the values measured
Table 1 Heat treatments of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy. Specimen
Step 1
A B C D E
1150 1150 1150 1150 1150
C/2 C/2 C/2 C/2 C/2
Step 2 h/WQ h/WQ h/WQ h/WQ h/WQ
e 850 850 850 850
C/30mins/WQ C/1 h/WQ C/3 h/WQ C/10 h/WQ
883
at more than five locations. 3. Results and discussion 3.1. Dissolution and re-precipitation behavior of uo phase As described in Ref. [23], the microstructure of the as-cast alloy was nearly fully lamellar and the volume fraction of the bo phase was measured to be approximately 9.5% by SEM imaging, with most of the phase distributed on the boundaries or at the triple-junctions of the lamellar colonies. A large number of uo phase regions shown by gray contrast were found within the bo phase. Fig. 1 shows the back-scattered electron (BSE) microstructure of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy after annealing. The microstructure is composed of (a2/g) colonies and bo phases. Similar to the as-cast alloy, the bo phases are mostly situated between (a2/g) lamellar colonies and the boundaries of the g phases. However, the uo particles found within the bo phase in the as-cast alloy disappeared after annealing at 1150 C for 2 h (as shown in Fig. 1a). This result was consistent with the observation that the uo phase solvus is below 950 C, as reported by Chladil et al. [24]. Fig. 1bed shows the microstructures of the alloy after annealing at 1150 C for 2 h; followed by transfer to 850 C for 1 h, 3 h and 10 h; and water quenching (the microstructure of the specimen annealed at 850 C for 30 min is similar to that found by SEM investigations for the specimen annealed for 1 h, which is not shown here). It was found that after annealing at 850 C for 1 h and 3 h, no uo phase could be observed by SEM, except for some g plates (shown by dark contrast, as indicated by arrows in Fig. 1bec) found within the bo region at the micrometer level. By contrast, the uo particles were observed readily to have a micro-size of approximately 1 mm in the specimen annealed at 850 C for 10 h (as shown by gray contrast and indicated by arrows in Fig. 1d). According to these results, it can be deduced that the uo phase precipitates from the bo phase at this temperature and that the g plates become coarser and tend to globalize with increased annealing time. Another interesting phenomenon is that serrated boundaries of bo and g grains were observed (the magnified images of the circled areas in Fig. 1b and c are shown in the insert images), which will be studied in detail by TEM. The phase constitutions in all of the specimens were examined by X-ray diffraction (not shown here) (all the specimens were crushed into approximate 200 mesh fine powder). However, it was difficult to clearly distinguish the relative variations among the diffraction intensities of these specimens. Probably this was because these specimens were held at 850 C only for a short time so that the volume fraction of the phases hardly changed. As shown in the magnified insert image in Fig. 1c, the equiaxed g grains at the bo/g boundary coarsened into the bo phase, similar to the results previously reported in Refs. [16,17]. In addition, the precipitation of nano-sized uo particles in the bo phase is difficult to distinguish by SEM. However, when these specimens were examined by TEM, the morphologies of the bo areas appeared to be different and were dependent on the annealing time. Since some studies reported that the Nb is an uo phase stabilized element and W is a bo phase stabilized element, the precipitation behavior of the uo from the bo phase is significantly related to the redistribution of these elements. This section will be interpreted by means of analyzing the composition via EDS. Fig. 2 shows TEM images of the specimen after annealing at 1150 C for 2 h followed by water quenching. The morphology of the bo phase corresponds to that observed by SEM, i.e., no uo phase was found in this region. According to the selected area diffraction (SAD) analysis at the center of the bo phase, weak streaks are present in the SAD pattern (shown in Fig. 2b), indicating the existence of an u-related structure. The diffuse intensities indicated the early
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Fig. 1. Backscatter SEM images of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy: (a) specimen A; (b) specimen C and the insert image in (b) is the magnified image; (c) specimen D and the inserted image in (c) is the magnified image; (d) specimen E.
Fig. 2. TEM images of the specimen A: (a) bright field image; (b) corresponding SAD pattern of the center of the bo phase.
stage of the formation of the uo phase [12]. This suggested that diffuse intensities appeared during the cooling process due to the dissolution of the uo phase into the bo phase when the temperature increased above the uo phase solvus at approximate 950 C [24e26], which is below the annealing temperature of 1150 C used in the present study. Fig. 3 shows the uo phases of approximately ten to hundreds of nanometers in size precipitated at the bo phase boundary with an increasing annealing time at 850 C for 30 min (Specimen B), 1 h (Specimen C) and 10 h (Specimen E). The corresponding SAD patterns and the DF images taken from the uo phase diffraction spots are shown in Fig. 3dei, respectively. It is important to note that the intensities of one variant of the uo-spots and superposition spots of the uo and bo phases are significantly increased, while the other variants of uo phase are weakened in Specimen B (as shown in
Fig. 3d). The corresponding dark field (DF) image (Fig. 3g) taken from the intensified uo-spot circled in Fig. 3d indicates that the intensified-variant uo-particles may have a larger size and greater volume fraction. Comparison of Figs. 2 and 3 shows that the growth of the uo phase is observed, indicating that the size of the uo phase increased so that it almost replaced the bo phase with increasing annealing time. The compositions of the uo and bo phases were measured by EDS, with the results shown in Table 2. According to the EDS results, it is found that the Nb content increases while the Al concentration decreases in uo phase with increasing annealing time at 850 C. However, there is a slight difference in the W concentration in the uo and bo phases in the EDS results, and experimental errors may affect the accuracy of the obtained composition. Song et al. reported that the uo phase is obviously rich in Nb, but depleted in W,
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Fig. 3. TEM images of the specimens after annealing at 850 C for (a) 30 min; (b) 1 h, the particle arrowed in (b) is D88-u and the corresponding SAD pattern is shown in the insert image in (b); (c) 10 h, the arrows indicate the residual uo particles in the transformed g, as discussed in Section 3.2; (d)e(f) the SAD patterns taken from the circled areas in (a)e(c); (g)e(i) DF images obtained for the circled region in (d)e(f).
uo phase is poor in W is consistent with reports that W is a bo phase stabilizer but not an uo phase stabilizer element in the TiAl based alloy with Nb and W additions [16].
Table 2 EDS results for different phases. Specimen
Phase
Composition (at.%) Ti
A B C
bo g bo uo bo uo
D88-u
D
bo uo
E
bo uo
D88-u
D88-u
gp a2
49.3 46.7 49.8 49.2 49.4 49.9 48.4 49.8 49.3 48.5 50.4 50.1 48.5 45.6 57.2
Al ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.7 0.7 0.5 0.4 0.6 0.6 0.8 0.6 0.7 0.6 0.5 0.7 0.5 0.4 0.3
34.3 44.6 33.0 31.3 33.3 31.7 22.3 33.4 31.9 22.7 32.3 30.6 22.5 45.1 33.4
± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.4 0.6 0.3 0.3 0.4 0.3 0.5 0.4 0.4 0.5 0.3 0.3 0.3 0.3 0.3
Nb
W
15.4 ± 0.6 8.1 ± 0.3 16.2 ± 0.6 18.8 ± 0.6 16.2 ± 0.7 17.8 ± 0.6 27.4 ± 1.0 15.6 ± 0.8 17.8 ± 0.6 27.0 ± 0.7 16.3 ± 0.6 18.4 ± 0.8 27.1 ± 0.8 8.6 ± 0.5 8.5 ± 0.4
0.9 0.5 0.9 0.6 1.1 0.6 1.8 1.1 0.9 1.8 1.0 0.8 1.9 0.6 0.9
± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.2 0.2 0.2 0.2 0.2 0.2 0.3 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2
compared to the bo phase in the as-cast alloy after annealing at 900 C for 100 h because of the uo phase formed from the solidification process may undergo a composition change and approach equilibrium considering the large size of the ingot and the relatively low cooling rate. The result obtained in the present study that the
3.2. Precipitation of the g phase Fig. 4aeb shows TEM images of the g plates and the coarser g phase observed in specimens B and E, respectively. The corresponding SAD patterns of the g phases are shown in the insert images in Fig. 4a and b. According to the SAD pattern in the insert image of Fig. 4a, the OR between the g plates and bo phase was {111} g//{110}bo and 〈011〉g//〈111〉bo, indicating that the g plates nucleated in the bo phase. Meanwhile, the weak diffraction spots of uo indicated that a small amount of uo phase formed in the bo phase at the same time. However, no OR was observed between the coarsened g and bo phases. The EDS results (Table 2) show that in contrast to the bo and uo phases, in g phase, the Ti and Nb contents decrease, while the Al content increases. The formation of the g phase with different compositions between the bo and uo phases could be controlled by diffusion at high temperature. Fig. 5aeb shows that the g phase grew into the bo region in specimen E, as was confirmed by using the diffraction spot of g to obtain the DF image (Fig. 5b), indicating that the g phases embedded in the bo(uo) region have the same orientation
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Fig. 4. TEM images of the g plates and the coarser g phases: (a) specimen B; (b) specimen E. Corresponding SAD patterns of the circled area are shown in the insert images in (a) and (b).
Fig. 5. TEM images of specimen E: (a) g phase near the bo(uo) phase transformed into the bo phase and the SAD pattern of the circled area are shown in the insert image; (b) is the corresponding DF image obtained by taking the circled spot in the SAD pattern in (a); (c) g laths within a2/g lamellar colonies grew into the bo region, the circled area shows the remaining uo particles in the transformed g phase, and the SAD patterns of the uo arrowed in (c) are shown in the insert image; (d) is the corresponding DF images by taking the uo phase diffraction spot as indicated by circle in the insert image in (c).
relationship as those of the grains next to the bo region. It is interesting to note that small particles existed in the g phases, as shown by the circles in Fig. 5c (similar results are shown in Fig. 3c).
By taking the diffraction spot of uo, indicated by the arrow in Fig. 5c for DF image, Fig. 5d shows that a small amount of uo particles (indicated by arrows in Fig. 5d) existed in the g phase near the g/bo
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boundary. Considering the crystal structure and compositional differences of the g and uo phases (Table 2), the uo to g transformation is difficult because the uo phase is a hexagonal structure with a low diffusion rate and the Al and Nb contents of these two phases are very different. Thus, the residual uo phases acted as pinning particles during boundary migration and dissolved into the g phase at some distance from the boundary as a result of diffusion. Consequently, the aligned residual uo phases, shown by circles in Fig. 5c, may indicate the trace of the initial boundary of the bo and g phases. According to the phase diagram published by Chen et al. [27], the equilibrium phase at 850 C in this alloy may consist of a2 and g phases. The calculated phase diagram has the similar results, except for the uo phase existing as an equilibrium phase below 800 C, as shown in Fig. 6, by using the Pandat software with the database from Ref. [28]. It is suggested that the retained bo phase from the solidification process may decompose into more stable phases, such as g and uo phases [13,23,26]. However, there are no reports indicated that bo/a2 transformation occurred during the decomposition process of the bo phase at 700e900 C in high Nb-TiAl alloys. This case was observed by TEM and is presented in Fig. 10 in the present study. Cheng et al. reported that Tb, the lowest temperature at which the stable b phase existed, was 1290 C in Ti44Al-8Nb, and it has been shown that the retained bo phase partially can transform to the g phase by a process analogous to the discontinuous coarsening of g lamellae and the decomposition of
887
the bo phase directly to the faceted g phase because of the direct nucleation of the g phase after the furnace is cooled from 1350 C [18]. As described above, a similar phenomenon was observed in this alloy after annealing at 850 C. It is suggested that the g phase nucleated directly in the bo phase while the g lamellae consume the bo phase through the migration of the bo/g boundaries during the annealing process. Both mechanisms were accomplished by diffusion. In fact, the coarsening process of the g phase was more commonly observed in the bo phase during annealing at 850 C, and no orientation relationship between bo and the coarsened g was observed in this study. It is reasonable to infer that the high-angle bo/g boundaries act as rapid diffusion channels, possibly resulting in the coarsening process of the g phase. 3.3. Precipitation behavior of D88-u phase Observation of the D88-u phase was interesting. Fig. 7 shows the crystal structure of the uo and D88-u phases. The unit cell of D88-u (Ti6Al6Nb4) is composed of three uo unit cells, with a vacancy occupying the 2b position; meanwhile, the Wyckoff site 4d was occupied by Nb in contrast to the uo (Ti4Al3Nb) phase [11,12]. According to the stoichiometric composition of uo and D88-u, D88-u is rich in Nb and thus the excess Nb may result in the formation of D88-u. Fig. 8 shows that the D88-u phase is present at the g/bo boundary in Specimen B (similar results are observed in Fig. 3b)
Fig. 6. Calculated phase diagram of the Ti-45Al-8.5Nb alloy.
Fig. 7. Schematic diagram of the uo and D88-u phases: (a) uo; (b) D88-u.
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Fig. 8. TEM images of the D88-u phase in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy: (a) specimen B (the arrows show the D88-u phase); (b)e(d) are corresponding SAD patterns in different zone axes.
with the corresponding SAD patterns in different zone axes, as shown in Fig. 8bed. Fig. 9 shows the STEM images of Specimen D; the presence of particles located at the bo boundary (denoted by arrows in Fig. 9a) that have a brighter contrast than that of the bo region suggests that the particles are rich in heavy elements. As shown in the image presented in the insert of Fig. 9a, the bright contrast particle is D88u. Fig. 9c shows the distribution of the elements in the square area, in good agreement with the EDS measurement; it is evident that D88-u is more rich in W and Nb than the bo and uo phases. Comparing the investigations of Ti-45Al-10Nb and Ti-44Al-8Nb1B [29], Huang proposed that the combination of Zr and Nb could play an important role in vacancy stabilization in the retained bo phase, promoting the formation of the D88-u structure [20]. Song et al. reported the precipitation of D88-u particles in the as-cast Ti45Al-8.5Nb-0.2W-0.2B-0.02Y alloy after annealing at 900 C and concluded that all small D88-u particles precipitated near the g/bo boundary and that the interface between the D88-u and g phases produced by bo(uo)/g phase transformation acted as the vacancy source for D88-u formation [16]. In this study, the pre-existing uo phases dissolve into the bo phase, and this condition is conducive to studying the formation mechanism of D88-u. As shown in Figs. 3b and 8a, a D88-u phase approximately several hundred nanometers in size was retained at the g/bo boundary, even if the reprecipitated uo phases are several times smaller than the D88-u
phase after annealing for a short time. It can be suggested that the formation of the D88-u phase is controlled by the interface migration at the g/bo boundary. Due to its large size, it is more conveniently to perform composition analysis of the D88-u phase by using EDS (show in Table 2). The EDS and STEM results indicated that the D88-u phase has more Nb and W than the uo phase, indicating that Nb and W are both stabilizer elements in the D88-u phase. Shang et al. calculated the vacancy formation energies and vacancy activation energies of Al, Nb, W in a first-principles study of the pure elements [30]. For the vacancy formation and vacancy activation energies, it was found that Al has the lowest values, in contrast to W, which showed the largest values. Thus, the vacancy concentration and self-diffusion coefficients are smaller for W and Nb than for Al at the same temperature. Although the vacancy formation and vacancy activation energies of Al, Nb and W were calculated for the pure element, the experimental investigations of the diffusion process in TiAl alloys found that Nb and W show lower diffusion rates by the vacancy diffusion mechanism, as reported in Refs. [31e34]. Considering these results, the formation process of the D88-u phase can be suggested as follows. During annealing, the bo/uo phase transformation was accomplished by enrichment of Nb, accompanied by depletion of Al and W in the uo phase. Thus, the un-transformed bo regions may have a high concentration of W and Al. Meanwhile, the bo/g phase transformation resulted in a diffusion flow of Al from the bo phase to the g phase and the
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Fig. 9. (a) STEM images of the specimen D; (b) magnified image of (a); (c) distribution of the elements in the square area.
opposite diffusion flow of Nb and W. Due to the different diffusion rates, Al could diffuse quickly from bo to g, but Nb and W concentrated in the bo phase at the g/bo boundary. Vacancies formed due to mass transport, and the g/bo boundary migrated from g to bo. Thus, the D88-u phase was formed with the re-ordering process of the vacancies in the uo phase and then dissolved into the g phase with the diffusion of Nb and W. Another phenomenon related to the composition change is the a2 phase. This phase was found in the bo phase and at the g/bo boundary in Specimens D and E, respectively. Fig. 10a and b shows a2 phase precipitation in the bo phase, close to the uo precipitates. It is suggested that the a2 phase may preferentially nucleate at the bo/ uo boundary within the bo region. According to the SAD pattern shown in the circled area in Fig. 10a, the ORs between these phases can be described as: {1120}a2//{0002}uo//{111}bo and 〈0001〉a2// 〈1120〉uo//〈110〉bo. Fig. 10c shows the a2 phase precipitated in the bo at the g/bo boundary in Specimen E, and the D88-u phase circled by dashed line in Fig. 10c was embedded in the transformed g phase during the bo/g transformation in Specimen E (the corresponding SAD patterns are shown in the inserted image). Fig. 10d shows the corresponding SAD pattern taken at the area circled by the solid line in Fig. 10c. This pattern indicated that there is a Blackburn orientation relationship between the a2 and the g phases near the
bo region, described as {0001}a2//{111}g and 〈1120〉a2//〈110〉g [35].
This can be interpreted as meaning that the un-transformed region in the bo phase was poor in W, Nb and Al after bo/uo, bo(uo)/g and bo(uo)/D88-u transformations and that the local composition is close to the a2 phase, favoring the formation of the a2 phase. Thus, the a2 phase precipitated at the bo/uo and bo(uo)/g boundaries after a considerable amount of the uo phase precipitated and was distributed within the bo phase accompanied by local composition changes. 4. Conclusions In this work, the precipitation behaviors of the uo, g, a2 and D88u phases were investigated in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y
alloy during annealing at 850 C for different times. The main results are summarized as follows: 1. The uo phases precipitate at the g/bo boundaries during annealing at 850 C, and the sizes of uo phase sizes increased with increasing annealing time. 2. The g plates precipitated within the bo phase and then became coarser and tended to globalize with the increasing annealing time due to the diffusion-controlled growth processing. 3. D88-u is observed at the g/bo boundaries in the present study. According to the analysis by TEM with EDS and STEM, the
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Fig. 10. TEM images of the a2 phase in specimens D and E: (a) a2 phase precipitated in the bo in specimen D, and the square area is magnified in the insert image; (b) the SAD pattern of the circled area in the insert image in (a); (c) a2 phase precipitated in the bo phase at the g/bo boundary in specimen E, and the SAD pattern of the D88-u phase circled by the dashed line is show in the insert image in (c); (d) corresponding SAD pattern of the a2 and g circled in (c) by the solid line.
formation process of D88-u is attributed to net vacancies that flow at the g/bo boundaries, while Nb and W are the stabilizer elements of the D88-u phase. 4. The a2 phase is found in the bo region and at the g/bo boundary; this is interpreted to mean that the un-transformed region in the bo phase is poor in W, Nb and Al and that the local composition is close to the a2 phase after the of the bo/uo, bo(uo)/g and bo(uo)/D88-u transformations. Acknowledgments This research was supported by the National Natural Science Foundation of China under Contract Nos. 51671016, 51601146, 51271016 and 51371144. References [1] Y. Kim, Gamma titanium aluminides, JOM 47 (1995) 38. [2] J.D.H. Paul, F. Appel, R. Wagner, The compression behaviour of niobium alloyed g-titanium alumindies, Acta Mater. 46 (1998) 1075. [3] Y.H. Wang, J.P. Lin, Y.H. He, C.K. Zu, G.L. Chen, Pore structures and thermal insulating properties of high Nb containing TiAl porous alloys, J. Alloy Compd. 492 (2010) 213. [4] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria among a (hcp), b (bcc) and g (L10) phases in TieAl base ternary alloys, Intermetallics 8 (2000) 855.
[5] 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) 625. [6] F.S. Sun, C.X. Cao, M.G. Yan, S.E. Kim, Y. Tailee, Alloying mechanism of beta stabilizers in a TiAl alloy, Metall. Mater. Trans. A 32 (2001) 1573. [7] X.J. Xu, J.P. Lin, Z.K. Teng, Y.L. Wang, G.L. Chen, On the microsegregation of Tie45Ale(8e9)Nbe(W, B, Y) alloy, Mater. Lett. 61 (2007) 369. [8] H. Clemens, W. Wallgram, S. Kremmer, V. Güther, A. Otto, A. Bartels, Design of novel b-solidifying TiAl alloys with adjustable b/B2-phase fraction and excellent hot-workability, Adv. Eng. Mater. 10 (2008) 707. [9] T. Voisin, J.P. Monchoux, M. Hantcherli, S. Mayer, H. Clemens, A. Couret, Microstructures and mechanical properties of a multi-phase b-solidifying TiAl alloy densified by spark plasma sintering, Acta Mater. 73 (2014) 107. [10] B.H. Li, Y.Y. Chen, Z.Q. Hou, F.T. Kong, Microstructure and mechanical properties of as-cast Tie43Ale9Ve0.3Y alloy, J. Alloys Compd. 473 (2009) 123. [11] L.A. Bendersky, W.J. Boettinger, B.P. Burton, F.S. Biancaniello, Ordered u-derivatives in a Ti-37.5Al-20Nb at% alloy, Scripta Metall. Mater. 24 (1990) 1541. [12] 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) 931. € 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) 241. [14] L. Song, X.J. Xu, L. You, Y.F. Liang, Y.L. Wang, J.P. Lin, Ordered a2 to uo phase transformations in high Nb-containing TiAl alloys, Acta Mater. 91 (2015) 330. [15] Z.W. Huang, T. Cong, Microstructural instability and embrittlement behaviour of an Al-lean, high-Nb g-TiAl-based alloy subjected to a long-term thermal exposure in air, Intermetallics 18 (2010) 161. [16] 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) 483. [17] B. Tang, W. Ou, H. Kou, J. Li, Experimental study of the dissolution and
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reprecipitation behaviors of uo phase in high Nb containing TiAl alloy, Mater. Charact. 109 (2015) 122. T.T. Cheng, M.H. Loretto, The decomposition of the beta phase in Tie44Ale8Nb and Tie44Ale4Nbe4Zre0.2Si alloys, Acta Mater. 46 (1998) 4801. L. Song, X.J. Xu, L. You, Y.F. Liang, J.P. Lin, Ordered u phase transformations in Ti-45Al-8.5Nb-0.2B alloy, Intermetallics 65 (2015) 22. Z.W. Huang, Ordered u phases in a 4Zre4Nb-containing TiAl-based alloy, Acta Mater. 56 (2008) 1689. R. Tewari, P. Mukhopadhyay, S. Banerjee, L.A. Bendersky, Evolution of ordered u phases in (Zr3Al)eNb alloys, Acta Mater. 47 (1999) 1307. S. Banerjee, R. Tewari, P. Mukhopadhyay, Coupling of displacive and replacive ordering, Prog. Mater. Sci. 42 (1997) 109. L. Song, L.Q. Zhang, X.J. Xu, J. Sun, J.P. Lin, Omega phase in as-cast high-Nbcontaining TiAl alloy, Scr. Mater. 68 (2013) 929. H.F. Chladil, H. Clemens, G.A. Zickler, M. Takeyama, E. Kozeschnik, A. Bartels, T. Buslaps, R. Gerling, S. Kremmer, L. Yeoh, K.D. Liss, Experimental studies and thermodynamic simulation of phase transformations in high Nb containing gTiAl based alloys, Int. J. Mater. Res. 98 (2007) 1131. A. Stark, A. Bartels, H. Clemens, F.P. Schimansky, On the formation of ordered u-phase in high Nb containing g-TiAl based alloys, Adv. Eng. Mater. 10 (2008) 929e934. H. Jiang, D. Hu, X. Wu, Thermal stability of the omega phase in Zr-containing TiAl alloys, J. Alloy Compd. 475 (2009) 134.
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[27] G.L. Chen, W.J. Zhang, Z.C. Liu, S.J. Li, Y.W. Kim, Gamma Titanium Aluminides, TMS, Warrendale, PA, 1999, p. 371. [28] V.T. Witusiewicz, A.A. Bondar, U. Hecht, T.Y. Velikanova, The AleBeNbeTi system: IV. Experimental study and thermodynamic re-evaluation of the binary AleNb and ternary AleNbeTi systems, J. Alloy Compd. 472 (2009) 133. [29] Z.W. Huang, D.G. Zhu, Thermal stability of Tie44Ale8Nbe1B alloy, Intermetallics 16 (2008) 156. [30] S. Shang, B. Zhou, W.Y. Wang, A.J. Ross, X.L. Liu, Y. Hu, H. Fang, Y. Wang, Z. Liu, A comprehensive first-principles study of pure elements: vacancy formation and migration energies and self-diffusion coefficients, Acta Mater. 109 (2016) 128. [31] C. Herzig, T. Przeorski, M. Friesel, F. Hisker, S. Divinski, Tracer solute diffusion of Nb, Zr, Cr, Fe, and Ni in g-TiAl: effect of preferential site occupation, Intermetallics 9 (2001) 461. [32] Y. Mishin, C. Herzig, Diffusion in the TieAl system, Acta Mater. 48 (2000) 589. [33] Y.H. Wang, J.P. Lin, Y.H. He, Y.L. Wang, G.L. Chen, Diffusion behavior of Nb element in high Nb containing TiAl Alloys by reactive hot pressing, Rare Met. 25 (2006) 349. [34] F. Appel, Diffusion assisted dislocation climb in intermetallic gamma TiAl, Mater. Sci. Eng. A 317 (2001) 115. [35] M.J. Blackburn, Some aspects of phase transformations in titanium alloys, in: R.I. Jafee, N.E. Promisel (Eds.), The Science, Technology and Application of Titanium, Pergamon Press, Oxford, 1970, p. 633.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Precipitation behavior of the uo phase in an annealed high Nb-TiAl alloy Teng Ye a, Lin Song b, **, Shubo Gao a, Yongfeng Liang a, Yanli Wang a, Junpin Lin a, * a b
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, 100083, China State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an, Shaanxi, 710072, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 31 May 2016 Received in revised form 28 December 2016 Accepted 18 January 2017 Available online 20 January 2017
Ordered omega (uo) phases are considered to be equilibrium phases in high Nb-TiAl alloys at intermediate temperatures. However, most studies have focused on the evolution behaviors of these phases during annealing, while the precipitation behavior has not been systematically studied. In the present work, the precipitation behavior of the uo phase (B82 structure) in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y (at.%) alloy during annealing at 850 C was investigated using scanning electron microscopy (SEM) and transmission electron microscopy (TEM). After annealing at 1150 C for 2 h followed by water quenching, the uo phase in the as-cast alloy transformed to the bo phase. At the initial stage of annealing processing at 850 C, the uo phase re-precipitated at the g/bo boundaries and the size of the uo phase increased with the annealing time. Moreover, the g plates precipitated within the bo phase at this temperature became coarser and tended to globalize during annealing. In addition, the a2 and D88-u phases were also observed at the g/bo boundaries. Energy dispersive X-ray spectroscopy (EDS) and Scanning transmission electron microscopy (STEM) analysis showed that the D88-u phase is rich in W and Nb elements. The corresponding mechanisms of the transformation mentioned above were discussed. © 2017 Elsevier B.V. All rights reserved.
Keywords: Titanium aluminides Phase transformation TEM Ordered omega phases
1. Introduction High Nb-TiAl alloys have been considered to be potential materials for high-temperature applications due to their low density, high strength, good oxidation resistance and creep properties [1e3]. As reported in many studies, Nb is a b phase stabilizer that extends the b phase region to lower temperatures and higher Al contents in the phase diagram [4,5]. Thus, a certain amount of b phase can be retained in an as-cast ingot during the cooling process [5e7]. Although the retained b phase is expected to facilitate the hot workability of TiAl alloys [8e10], the ordered bo(B2 structure) phase transformed from the disordered bcc b(A2 structure) phase is considered to be detrimental to the ductility of TiAl alloys [5]. Recently, several studies have reported ordered u phase transitions in the bo phase regions in high Nb-TiAl alloys that involved chemical (replacive) and displacive order/disorder transitions [11,12]. Commonly, the ordered u phase in the B82 structure is
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (L. Song), [email protected] (J. Lin). http://dx.doi.org/10.1016/j.jallcom.2017.01.198 0925-8388/© 2017 Elsevier B.V. All rights reserved.
assigned as uo (space group: P63/mmc). Many reports have indicated that the uo phase in high Nb-TiAl alloys was an equilibrium phase at 700 C-900 C and that the observed orientation relationship (OR) between the uo and bo phases was [0001]uo//[111]bo; (1120)uo//(110)bo [13,14]. Huang et al. indicated that there is a three-phase region, namely, a2þgþbo(uo) (the ordered u formed within the bo phase) at 700 C in the Ti-44Al-8Nb-1B alloy, with further results showing that a2þg transformed into bo(uo) following thermal exposure at 700 C [15]. Meanwhile, g plates precipitated from the bo phase were also observed in Refs. [13,16e19]. Song et al. reported that the bo(uo)/g transformation can be accomplished by direct nucleation within the bo phase and growth of the pre-existing g phase in the Ti-45Al-8.5Nb(W, B, Y) alloy after annealing at 900 C [16]. Similar results were reported by Tang et al. in an alloy with the same composition after annealing at 700 C and 800 C [17]. It was indicated that due to the enhanced diffusion process, the g plates became coarser and globalized when the annealing temperature increased from 700 C to 800 C. Meanwhile, the D88-u phase (space group: P63/mcm) was also observed in a TiAl alloy with Nb and Zr additions formed through the occupation of the 2b positions by vacancies in the uo phase [11,20e22]. Tewari et al. proposed that Nb had a significant
T. Ye et al. / Journal of Alloys and Compounds 701 (2017) 882e891
effect on the stabilization of vacancies during the formation of the D88 structure in the (Zr3Al)-Nb alloy [21]. Huang reported that the orientation relationship between the D88-u, uo and bo phases in the Ti-44Al-4Nb-4Zr-0.2Si-1B alloy was 〈0001〉D88-u//〈0001〉uo// 〈111〉bo and {1010}D88-u//{1120}uo//{110}bo and the structure transformation path in the retained b after annealing at 700 for 1000 h was b/bo/uo/D88-u [20]. Similar results were also reported by Song et al. in Ref. [16] where the vacancies produced by bo(uo)/g phase transformation acted as the main cause for D88-u formation and that D88-u could be dissolved into the g phase at an area far away from the g/bo(uo) boundary; it was suggested that the formation of D88-u was significantly affected by the vacancy concentration. These results suggest that the phase transformation that occurs in high Nb-TiAl alloys is strongly dependent on the local composition. However, investigations of the phase transformation mechanism related to composition are rare. Therefore, the formation mechanism of the D88-u phase and decomposition of the bo(uo) phase during annealing should be investigated more carefully, and a complete understanding of the phase transformation in high Nb-TiAl alloys at service temperatures is meaningful for the application of these alloys. In the present work, the phase transformation processes in the bo phase were investigated during annealing at 850 C and the corresponding transformation mechanisms were discussed in terms of composition analysis. 2. Experiments An ingot of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy used in this study was prepared using a plasma cold hearth melting furnace. Specimens with sizes of 15 10 10 mm were cut from the center of the ingot by electric-discharge machining. To reveal the precipitation behavior of the alloy at intermediate temperature, the samples were first annealed at 1150 C for 2 h, followed by water quenching (designed specimen A) to remove the ordered omega phases in the as-cast alloy [23], and then, some specimens were annealed at 850 C for 30min, 1 h, 3 h and 10 h, followed by water quenching (designed as specimens B, C, D, E, respectively) as listed in Table 1. The microstructures after the heat treatments were examined using a Zeiss Supra 55 scanning electron microscope (SEM) in backscatter mode operating at 15 kV. Discs with diameters of 3 mm and thicknesses of 0.3 mm were cut from the specimens by spark discharge cutting and were then mechanically ground to a thickness of 50e60 mm. Thin foils used for TEM observation were prepared by twin-jet electro-polishing in a solution of 65 vol% methanol, 30 vol% butanol, and 5 vol% perchloric acid at 30 V and 30 C. TEM analysis was conducted using a Tecnai G2 F30 field emission transmission electron microscope operating at 300 kV. The compositions and the element distributions of the phases were obtained by an energy dispersive X-ray spectroscopy (EDS) on TEM and Scanning transmission electron microscopy (STEM), respectively. The composition measurements were processed in the adjacent region to ensure the uniform thickness of the specimen. Each parameter was obtained as an average of the values measured
Table 1 Heat treatments of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy. Specimen
Step 1
A B C D E
1150 1150 1150 1150 1150
C/2 C/2 C/2 C/2 C/2
Step 2 h/WQ h/WQ h/WQ h/WQ h/WQ
e 850 850 850 850
C/30mins/WQ C/1 h/WQ C/3 h/WQ C/10 h/WQ
883
at more than five locations. 3. Results and discussion 3.1. Dissolution and re-precipitation behavior of uo phase As described in Ref. [23], the microstructure of the as-cast alloy was nearly fully lamellar and the volume fraction of the bo phase was measured to be approximately 9.5% by SEM imaging, with most of the phase distributed on the boundaries or at the triple-junctions of the lamellar colonies. A large number of uo phase regions shown by gray contrast were found within the bo phase. Fig. 1 shows the back-scattered electron (BSE) microstructure of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy after annealing. The microstructure is composed of (a2/g) colonies and bo phases. Similar to the as-cast alloy, the bo phases are mostly situated between (a2/g) lamellar colonies and the boundaries of the g phases. However, the uo particles found within the bo phase in the as-cast alloy disappeared after annealing at 1150 C for 2 h (as shown in Fig. 1a). This result was consistent with the observation that the uo phase solvus is below 950 C, as reported by Chladil et al. [24]. Fig. 1bed shows the microstructures of the alloy after annealing at 1150 C for 2 h; followed by transfer to 850 C for 1 h, 3 h and 10 h; and water quenching (the microstructure of the specimen annealed at 850 C for 30 min is similar to that found by SEM investigations for the specimen annealed for 1 h, which is not shown here). It was found that after annealing at 850 C for 1 h and 3 h, no uo phase could be observed by SEM, except for some g plates (shown by dark contrast, as indicated by arrows in Fig. 1bec) found within the bo region at the micrometer level. By contrast, the uo particles were observed readily to have a micro-size of approximately 1 mm in the specimen annealed at 850 C for 10 h (as shown by gray contrast and indicated by arrows in Fig. 1d). According to these results, it can be deduced that the uo phase precipitates from the bo phase at this temperature and that the g plates become coarser and tend to globalize with increased annealing time. Another interesting phenomenon is that serrated boundaries of bo and g grains were observed (the magnified images of the circled areas in Fig. 1b and c are shown in the insert images), which will be studied in detail by TEM. The phase constitutions in all of the specimens were examined by X-ray diffraction (not shown here) (all the specimens were crushed into approximate 200 mesh fine powder). However, it was difficult to clearly distinguish the relative variations among the diffraction intensities of these specimens. Probably this was because these specimens were held at 850 C only for a short time so that the volume fraction of the phases hardly changed. As shown in the magnified insert image in Fig. 1c, the equiaxed g grains at the bo/g boundary coarsened into the bo phase, similar to the results previously reported in Refs. [16,17]. In addition, the precipitation of nano-sized uo particles in the bo phase is difficult to distinguish by SEM. However, when these specimens were examined by TEM, the morphologies of the bo areas appeared to be different and were dependent on the annealing time. Since some studies reported that the Nb is an uo phase stabilized element and W is a bo phase stabilized element, the precipitation behavior of the uo from the bo phase is significantly related to the redistribution of these elements. This section will be interpreted by means of analyzing the composition via EDS. Fig. 2 shows TEM images of the specimen after annealing at 1150 C for 2 h followed by water quenching. The morphology of the bo phase corresponds to that observed by SEM, i.e., no uo phase was found in this region. According to the selected area diffraction (SAD) analysis at the center of the bo phase, weak streaks are present in the SAD pattern (shown in Fig. 2b), indicating the existence of an u-related structure. The diffuse intensities indicated the early
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Fig. 1. Backscatter SEM images of the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy: (a) specimen A; (b) specimen C and the insert image in (b) is the magnified image; (c) specimen D and the inserted image in (c) is the magnified image; (d) specimen E.
Fig. 2. TEM images of the specimen A: (a) bright field image; (b) corresponding SAD pattern of the center of the bo phase.
stage of the formation of the uo phase [12]. This suggested that diffuse intensities appeared during the cooling process due to the dissolution of the uo phase into the bo phase when the temperature increased above the uo phase solvus at approximate 950 C [24e26], which is below the annealing temperature of 1150 C used in the present study. Fig. 3 shows the uo phases of approximately ten to hundreds of nanometers in size precipitated at the bo phase boundary with an increasing annealing time at 850 C for 30 min (Specimen B), 1 h (Specimen C) and 10 h (Specimen E). The corresponding SAD patterns and the DF images taken from the uo phase diffraction spots are shown in Fig. 3dei, respectively. It is important to note that the intensities of one variant of the uo-spots and superposition spots of the uo and bo phases are significantly increased, while the other variants of uo phase are weakened in Specimen B (as shown in
Fig. 3d). The corresponding dark field (DF) image (Fig. 3g) taken from the intensified uo-spot circled in Fig. 3d indicates that the intensified-variant uo-particles may have a larger size and greater volume fraction. Comparison of Figs. 2 and 3 shows that the growth of the uo phase is observed, indicating that the size of the uo phase increased so that it almost replaced the bo phase with increasing annealing time. The compositions of the uo and bo phases were measured by EDS, with the results shown in Table 2. According to the EDS results, it is found that the Nb content increases while the Al concentration decreases in uo phase with increasing annealing time at 850 C. However, there is a slight difference in the W concentration in the uo and bo phases in the EDS results, and experimental errors may affect the accuracy of the obtained composition. Song et al. reported that the uo phase is obviously rich in Nb, but depleted in W,
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Fig. 3. TEM images of the specimens after annealing at 850 C for (a) 30 min; (b) 1 h, the particle arrowed in (b) is D88-u and the corresponding SAD pattern is shown in the insert image in (b); (c) 10 h, the arrows indicate the residual uo particles in the transformed g, as discussed in Section 3.2; (d)e(f) the SAD patterns taken from the circled areas in (a)e(c); (g)e(i) DF images obtained for the circled region in (d)e(f).
uo phase is poor in W is consistent with reports that W is a bo phase stabilizer but not an uo phase stabilizer element in the TiAl based alloy with Nb and W additions [16].
Table 2 EDS results for different phases. Specimen
Phase
Composition (at.%) Ti
A B C
bo g bo uo bo uo
D88-u
D
bo uo
E
bo uo
D88-u
D88-u
gp a2
49.3 46.7 49.8 49.2 49.4 49.9 48.4 49.8 49.3 48.5 50.4 50.1 48.5 45.6 57.2
Al ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.7 0.7 0.5 0.4 0.6 0.6 0.8 0.6 0.7 0.6 0.5 0.7 0.5 0.4 0.3
34.3 44.6 33.0 31.3 33.3 31.7 22.3 33.4 31.9 22.7 32.3 30.6 22.5 45.1 33.4
± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.4 0.6 0.3 0.3 0.4 0.3 0.5 0.4 0.4 0.5 0.3 0.3 0.3 0.3 0.3
Nb
W
15.4 ± 0.6 8.1 ± 0.3 16.2 ± 0.6 18.8 ± 0.6 16.2 ± 0.7 17.8 ± 0.6 27.4 ± 1.0 15.6 ± 0.8 17.8 ± 0.6 27.0 ± 0.7 16.3 ± 0.6 18.4 ± 0.8 27.1 ± 0.8 8.6 ± 0.5 8.5 ± 0.4
0.9 0.5 0.9 0.6 1.1 0.6 1.8 1.1 0.9 1.8 1.0 0.8 1.9 0.6 0.9
± ± ± ± ± ± ± ± ± ± ± ± ± ± ±
0.2 0.2 0.2 0.2 0.2 0.2 0.3 0.2 0.2 0.2 0.2 0.2 0.2 0.2 0.2
compared to the bo phase in the as-cast alloy after annealing at 900 C for 100 h because of the uo phase formed from the solidification process may undergo a composition change and approach equilibrium considering the large size of the ingot and the relatively low cooling rate. The result obtained in the present study that the
3.2. Precipitation of the g phase Fig. 4aeb shows TEM images of the g plates and the coarser g phase observed in specimens B and E, respectively. The corresponding SAD patterns of the g phases are shown in the insert images in Fig. 4a and b. According to the SAD pattern in the insert image of Fig. 4a, the OR between the g plates and bo phase was {111} g//{110}bo and 〈011〉g//〈111〉bo, indicating that the g plates nucleated in the bo phase. Meanwhile, the weak diffraction spots of uo indicated that a small amount of uo phase formed in the bo phase at the same time. However, no OR was observed between the coarsened g and bo phases. The EDS results (Table 2) show that in contrast to the bo and uo phases, in g phase, the Ti and Nb contents decrease, while the Al content increases. The formation of the g phase with different compositions between the bo and uo phases could be controlled by diffusion at high temperature. Fig. 5aeb shows that the g phase grew into the bo region in specimen E, as was confirmed by using the diffraction spot of g to obtain the DF image (Fig. 5b), indicating that the g phases embedded in the bo(uo) region have the same orientation
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Fig. 4. TEM images of the g plates and the coarser g phases: (a) specimen B; (b) specimen E. Corresponding SAD patterns of the circled area are shown in the insert images in (a) and (b).
Fig. 5. TEM images of specimen E: (a) g phase near the bo(uo) phase transformed into the bo phase and the SAD pattern of the circled area are shown in the insert image; (b) is the corresponding DF image obtained by taking the circled spot in the SAD pattern in (a); (c) g laths within a2/g lamellar colonies grew into the bo region, the circled area shows the remaining uo particles in the transformed g phase, and the SAD patterns of the uo arrowed in (c) are shown in the insert image; (d) is the corresponding DF images by taking the uo phase diffraction spot as indicated by circle in the insert image in (c).
relationship as those of the grains next to the bo region. It is interesting to note that small particles existed in the g phases, as shown by the circles in Fig. 5c (similar results are shown in Fig. 3c).
By taking the diffraction spot of uo, indicated by the arrow in Fig. 5c for DF image, Fig. 5d shows that a small amount of uo particles (indicated by arrows in Fig. 5d) existed in the g phase near the g/bo
T. Ye et al. / Journal of Alloys and Compounds 701 (2017) 882e891
boundary. Considering the crystal structure and compositional differences of the g and uo phases (Table 2), the uo to g transformation is difficult because the uo phase is a hexagonal structure with a low diffusion rate and the Al and Nb contents of these two phases are very different. Thus, the residual uo phases acted as pinning particles during boundary migration and dissolved into the g phase at some distance from the boundary as a result of diffusion. Consequently, the aligned residual uo phases, shown by circles in Fig. 5c, may indicate the trace of the initial boundary of the bo and g phases. According to the phase diagram published by Chen et al. [27], the equilibrium phase at 850 C in this alloy may consist of a2 and g phases. The calculated phase diagram has the similar results, except for the uo phase existing as an equilibrium phase below 800 C, as shown in Fig. 6, by using the Pandat software with the database from Ref. [28]. It is suggested that the retained bo phase from the solidification process may decompose into more stable phases, such as g and uo phases [13,23,26]. However, there are no reports indicated that bo/a2 transformation occurred during the decomposition process of the bo phase at 700e900 C in high Nb-TiAl alloys. This case was observed by TEM and is presented in Fig. 10 in the present study. Cheng et al. reported that Tb, the lowest temperature at which the stable b phase existed, was 1290 C in Ti44Al-8Nb, and it has been shown that the retained bo phase partially can transform to the g phase by a process analogous to the discontinuous coarsening of g lamellae and the decomposition of
887
the bo phase directly to the faceted g phase because of the direct nucleation of the g phase after the furnace is cooled from 1350 C [18]. As described above, a similar phenomenon was observed in this alloy after annealing at 850 C. It is suggested that the g phase nucleated directly in the bo phase while the g lamellae consume the bo phase through the migration of the bo/g boundaries during the annealing process. Both mechanisms were accomplished by diffusion. In fact, the coarsening process of the g phase was more commonly observed in the bo phase during annealing at 850 C, and no orientation relationship between bo and the coarsened g was observed in this study. It is reasonable to infer that the high-angle bo/g boundaries act as rapid diffusion channels, possibly resulting in the coarsening process of the g phase. 3.3. Precipitation behavior of D88-u phase Observation of the D88-u phase was interesting. Fig. 7 shows the crystal structure of the uo and D88-u phases. The unit cell of D88-u (Ti6Al6Nb4) is composed of three uo unit cells, with a vacancy occupying the 2b position; meanwhile, the Wyckoff site 4d was occupied by Nb in contrast to the uo (Ti4Al3Nb) phase [11,12]. According to the stoichiometric composition of uo and D88-u, D88-u is rich in Nb and thus the excess Nb may result in the formation of D88-u. Fig. 8 shows that the D88-u phase is present at the g/bo boundary in Specimen B (similar results are observed in Fig. 3b)
Fig. 6. Calculated phase diagram of the Ti-45Al-8.5Nb alloy.
Fig. 7. Schematic diagram of the uo and D88-u phases: (a) uo; (b) D88-u.
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Fig. 8. TEM images of the D88-u phase in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y alloy: (a) specimen B (the arrows show the D88-u phase); (b)e(d) are corresponding SAD patterns in different zone axes.
with the corresponding SAD patterns in different zone axes, as shown in Fig. 8bed. Fig. 9 shows the STEM images of Specimen D; the presence of particles located at the bo boundary (denoted by arrows in Fig. 9a) that have a brighter contrast than that of the bo region suggests that the particles are rich in heavy elements. As shown in the image presented in the insert of Fig. 9a, the bright contrast particle is D88u. Fig. 9c shows the distribution of the elements in the square area, in good agreement with the EDS measurement; it is evident that D88-u is more rich in W and Nb than the bo and uo phases. Comparing the investigations of Ti-45Al-10Nb and Ti-44Al-8Nb1B [29], Huang proposed that the combination of Zr and Nb could play an important role in vacancy stabilization in the retained bo phase, promoting the formation of the D88-u structure [20]. Song et al. reported the precipitation of D88-u particles in the as-cast Ti45Al-8.5Nb-0.2W-0.2B-0.02Y alloy after annealing at 900 C and concluded that all small D88-u particles precipitated near the g/bo boundary and that the interface between the D88-u and g phases produced by bo(uo)/g phase transformation acted as the vacancy source for D88-u formation [16]. In this study, the pre-existing uo phases dissolve into the bo phase, and this condition is conducive to studying the formation mechanism of D88-u. As shown in Figs. 3b and 8a, a D88-u phase approximately several hundred nanometers in size was retained at the g/bo boundary, even if the reprecipitated uo phases are several times smaller than the D88-u
phase after annealing for a short time. It can be suggested that the formation of the D88-u phase is controlled by the interface migration at the g/bo boundary. Due to its large size, it is more conveniently to perform composition analysis of the D88-u phase by using EDS (show in Table 2). The EDS and STEM results indicated that the D88-u phase has more Nb and W than the uo phase, indicating that Nb and W are both stabilizer elements in the D88-u phase. Shang et al. calculated the vacancy formation energies and vacancy activation energies of Al, Nb, W in a first-principles study of the pure elements [30]. For the vacancy formation and vacancy activation energies, it was found that Al has the lowest values, in contrast to W, which showed the largest values. Thus, the vacancy concentration and self-diffusion coefficients are smaller for W and Nb than for Al at the same temperature. Although the vacancy formation and vacancy activation energies of Al, Nb and W were calculated for the pure element, the experimental investigations of the diffusion process in TiAl alloys found that Nb and W show lower diffusion rates by the vacancy diffusion mechanism, as reported in Refs. [31e34]. Considering these results, the formation process of the D88-u phase can be suggested as follows. During annealing, the bo/uo phase transformation was accomplished by enrichment of Nb, accompanied by depletion of Al and W in the uo phase. Thus, the un-transformed bo regions may have a high concentration of W and Al. Meanwhile, the bo/g phase transformation resulted in a diffusion flow of Al from the bo phase to the g phase and the
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Fig. 9. (a) STEM images of the specimen D; (b) magnified image of (a); (c) distribution of the elements in the square area.
opposite diffusion flow of Nb and W. Due to the different diffusion rates, Al could diffuse quickly from bo to g, but Nb and W concentrated in the bo phase at the g/bo boundary. Vacancies formed due to mass transport, and the g/bo boundary migrated from g to bo. Thus, the D88-u phase was formed with the re-ordering process of the vacancies in the uo phase and then dissolved into the g phase with the diffusion of Nb and W. Another phenomenon related to the composition change is the a2 phase. This phase was found in the bo phase and at the g/bo boundary in Specimens D and E, respectively. Fig. 10a and b shows a2 phase precipitation in the bo phase, close to the uo precipitates. It is suggested that the a2 phase may preferentially nucleate at the bo/ uo boundary within the bo region. According to the SAD pattern shown in the circled area in Fig. 10a, the ORs between these phases can be described as: {1120}a2//{0002}uo//{111}bo and 〈0001〉a2// 〈1120〉uo//〈110〉bo. Fig. 10c shows the a2 phase precipitated in the bo at the g/bo boundary in Specimen E, and the D88-u phase circled by dashed line in Fig. 10c was embedded in the transformed g phase during the bo/g transformation in Specimen E (the corresponding SAD patterns are shown in the inserted image). Fig. 10d shows the corresponding SAD pattern taken at the area circled by the solid line in Fig. 10c. This pattern indicated that there is a Blackburn orientation relationship between the a2 and the g phases near the
bo region, described as {0001}a2//{111}g and 〈1120〉a2//〈110〉g [35].
This can be interpreted as meaning that the un-transformed region in the bo phase was poor in W, Nb and Al after bo/uo, bo(uo)/g and bo(uo)/D88-u transformations and that the local composition is close to the a2 phase, favoring the formation of the a2 phase. Thus, the a2 phase precipitated at the bo/uo and bo(uo)/g boundaries after a considerable amount of the uo phase precipitated and was distributed within the bo phase accompanied by local composition changes. 4. Conclusions In this work, the precipitation behaviors of the uo, g, a2 and D88u phases were investigated in the Ti-45Al-8.5Nb-0.2W-0.2B-0.02Y
alloy during annealing at 850 C for different times. The main results are summarized as follows: 1. The uo phases precipitate at the g/bo boundaries during annealing at 850 C, and the sizes of uo phase sizes increased with increasing annealing time. 2. The g plates precipitated within the bo phase and then became coarser and tended to globalize with the increasing annealing time due to the diffusion-controlled growth processing. 3. D88-u is observed at the g/bo boundaries in the present study. According to the analysis by TEM with EDS and STEM, the
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Fig. 10. TEM images of the a2 phase in specimens D and E: (a) a2 phase precipitated in the bo in specimen D, and the square area is magnified in the insert image; (b) the SAD pattern of the circled area in the insert image in (a); (c) a2 phase precipitated in the bo phase at the g/bo boundary in specimen E, and the SAD pattern of the D88-u phase circled by the dashed line is show in the insert image in (c); (d) corresponding SAD pattern of the a2 and g circled in (c) by the solid line.
formation process of D88-u is attributed to net vacancies that flow at the g/bo boundaries, while Nb and W are the stabilizer elements of the D88-u phase. 4. The a2 phase is found in the bo region and at the g/bo boundary; this is interpreted to mean that the un-transformed region in the bo phase is poor in W, Nb and Al and that the local composition is close to the a2 phase after the of the bo/uo, bo(uo)/g and bo(uo)/D88-u transformations. Acknowledgments This research was supported by the National Natural Science Foundation of China under Contract Nos. 51671016, 51601146, 51271016 and 51371144. References [1] Y. Kim, Gamma titanium aluminides, JOM 47 (1995) 38. [2] J.D.H. Paul, F. Appel, R. Wagner, The compression behaviour of niobium alloyed g-titanium alumindies, Acta Mater. 46 (1998) 1075. [3] Y.H. Wang, J.P. Lin, Y.H. He, C.K. Zu, G.L. Chen, Pore structures and thermal insulating properties of high Nb containing TiAl porous alloys, J. Alloy Compd. 492 (2010) 213. [4] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria among a (hcp), b (bcc) and g (L10) phases in TieAl base ternary alloys, Intermetallics 8 (2000) 855.
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