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Solution treatment for enhanced hardness in Mo-modified Ti2AlNb-based alloys
Journal of Alloys and Compounds 805 (2019) 1184e1190
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Solution treatment for enhanced hardness in Mo-modified Ti2AlNb-based alloys Yaran Zhang a, Qi Cai b, Zongqing Ma a, *, Chong Li a, Liming Yu a, Yongchang Liu a, ** a
State Key Laboratory of Hydraulic Engineering Simulation and Safety, School of Materials Science and Engineering, Tianjin University, Tianjin, 300354, China b Beijing Inst Technol, Sch Mat Sci & Engn, Beijing, 100081, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 March 2019 Received in revised form 11 July 2019 Accepted 15 July 2019 Available online 16 July 2019
The present work investigates the change in the microstructure and the improvement in the micro€tten hardness in the Mo-modified Ti2AlNb alloys induced by solution treatment. Typical Widmansta O þ B2 structure was observed in the alloys directly aged in the O þ B2 phase region, with a lower hardness value compared to that of the alloy without Mo addition. Well-arranged O þ B2 colonies were observed in the specimen solution treated in single B2 phase region, and it exhibited a favorable hardness performance of up to 620 HV. Besides the precipitation hardening of the O phase, the solution of Mo in B2 phase should be responsible for the enhancement of hardness. Further microstructural observation proved that the substitution of Mo for Nb in the lattice of B2 phase induced the lattice distortion mainly the (200) and (211) planes, and the induced entanglement of dislocations also contributed to the enhancement of hardness. © 2019 Elsevier B.V. All rights reserved.
Keywords: Ti2AlNb alloys Mo addition Solution treatment Microstructure Hardness
1. Introduction Intermetallic Ti2AlNb alloys, based on the orthorhombic (O) phase, have exhibited great potential for application in the components of engine, due to the excellent high-temperature properties. Typical Ti2AlNb-based alloys include Tie25Ale17Nb, Tie22Ale25Nb, and Tie23Ale27Nb, and these alloys containing high amounts of O phase have shown favorable mechanical properties, e.g. higher strength, and excellent creep resistance, compared with the conventional Ti3Al alloys [1e3]. Apart from the O phase, the Ti2AlNb-based alloys may also contain B2 phase (bodycentered cubic, bcc) and a2 phase (DO19 structure, Ti3Al) [4e6], and the precipitation strengthening of the O/a2 phase is currently the main concern for enhancement of the mechanical properties. For instance, fine acicular O/a2 phase was confirmed to be favorable to the tensile strength, whereas coarse O/a2 phase might improve the ductility of the Ti2AlNb-based alloys [7]. To further explore the possible application of Ti2AlNb-based alloys in aerospace, systematic investigation has been performed on the precipitation behavior of O/a2 phase, including the volume fraction, shape, and dimension,
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Z. Ma), [email protected] (Y. Liu). https://doi.org/10.1016/j.jallcom.2019.07.149 0925-8388/© 2019 Elsevier B.V. All rights reserved.
which are strongly related to the processing technique and heat treatment [8e11]. Adding b/a alloying elements is a convenient route to regulate the microstructure of the Ti2AlNb alloys. Mao et al. [12] have reported the substitution of Ta for Nb improved the strength and ductility of the Ti2AlNb alloy, and the yield strength of the deformed Tie22Ale20Nbe7Ta alloy reached 1200 MPa. Hagiwara et al. [13] produced W-modified Tie22Ale27Nb alloys with excellent high-temperature strength and creep properties. Alloying procedure is conveniently accomplished by directly adding the element powder into pre-alloyed powder, i.e. a powder metallurgic route, which makes it widely employed to adjust the composition of various alloys [14e17], Moreover, the outstanding advantage of powder metallurgy also lies in the direct fabrication of Ti2AlNb alloys with desirable shape, microstructure and properties. Yang et al. [18] have performed the W alloying by powder metal€tten structure was significantly refined, lurgy, and the Widmansta which was responsible to the enhancement of the hardness for the W-modified Ti2AlNb-based alloys. In addition to the refinement of €tten structure, diverse O þ B2 structures have been the Widmansta observed in the Ti2AlNb-based alloys. Emura et al. [19] have observed ‘Van Gogh's Sky’ structure with high elongation-to-failure in the annealed Mo- and Fe-modified Tie25Ale14Nb alloys. Yang et al. [20] have obtained herringbone O þ B2 structure by ageing the W-modified Ti2AlNb alloys in the O þ B2 phase region, and the
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alloys exhibited high hardness of 750 HV. The orientation relationships between the three phases are normally (0001)a2// (011)B2 and (110)B2//(001)O [4,21], but additional orientation may occur to form special structure. Studies also revealed that Mo addition could reduce the contiguity of a2 phase, which lowered the creep rate of the Ti2AlNbbased alloy [22]. The alloy with lath structure showed higher creep resistance than the three-phase microstructure with equiaxed phase [23]. As far as the mechanical properties of Mo-modified Ti2AlNb-based alloys, Zhao et al. [24] have investigated the tensile behavior of Tie22Ale24Nb-0.5Mo in the range of 25e650 C. The alloy with a2þB2 possessed the largest elongation, and the one with O/a2þB2 structure has the optimal balance between strength and plasticity. Jiao et al. [25] studied the microstructure and creep behavior of Tie22Ale24Nb-0.5Mo alloy. The creep resistance was significantly improved by eliminating coarse O þ B2 laths. However, the microhardness of the alloys, associated with the microstructure, has not been systematically investigated by solution treatment and ageing in different phase region; and moreover, the improvement or depression of hardness in Mo-added Ti2AlNb alloys was to date barely investigated from the viewpoint of lattice distortion. In this study, the sintered Mo-modified Ti2AlNb-based alloys are heat treated following different routes, either directly aged in O þ B2 and Oþa2þB2 phase regions, or solution treated in single B2 phase region and subsequent aged in O þ B2 and Oþa2þB2 phase regions. Different O þ B2 microstructure was obtained via these heat treatment process, and the resultant hardness performance was evaluated. Based on measuring the dimension and volume faction of O/a2, and investigating the lattice distortion by Mo substitution for Nb, the strengthening mechanism of the O/a2 phase was analyzed. 2. Experimental details Molybdenum-modified Ti2AlNb alloys, with the nominal composition of Tie22Al-23.9Nb-1.1Mo (at.%), were spark plasma sintered from spherical Tie22Ale25Nb pre-alloyed powder (~200 mm in size, single B2 phase) and Mo powder (~75 mm in diameter, 99.99% in purity) at 1100 C for 10 min and under 40 MPa. Two types of heat treatment were carried out on the as-SPSed alloys. The obtained alloys were either directly aged in different phase regions, or experienced solution treatment and subsequent ageing in different regions. The direct ageing was at 750 C, 800 C, 850 C, 900 C, 950 C and 1000 C for 2 h, respectively. The solution treatment was carried out at 1300 C for 6 h, followed by water quenching, and the subsequent ageing was at 750 C, 800 C, 850 C, 900 C, 950 C and 1000 C for 2 h, respectively, followed by air cooling. Fig. 1 shows the vertical section of the AleNbeTi phase diagram at 22 at.% Al [26], in which the phase regions for ageing are indicated. A pipe furnace (SK-G06123K) was used for the heat treatment after the samples were sealed in quartz tubes to avoid oxidation. The Mo-free Tie22Ale25Nb alloy were also aged at 850 C for 2 h without solution treatment as a control sample. The phase composition of the Ti2AlNb-based alloys were characterized by X-ray diffraction (XRD, Bruker D8 Advanced) using Cu Ka radiation, and the microstructure were observed by scanning electron microscopy (SEM, Hitachi S-4800), respectively. The alloys were chemically etched using Kroll's reagent. The orientation relation of the B2 and O/a2 was confirmed by transmission electron microscopy operated at 200 kV (TEM, JEOL-2100f). The roomtemperature Vickers microhardness was obtained under 0.2 kg load. The room temperature tensile tests were conducted on the electronic universal testing machine (Instron 5848) with the strain rate was 0.04 min1。
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Fig. 1. Vertical section of the AleNbeTi phase diagram at 22 at.% Al. Data are taken from Ref. [26].
3. Results Fig. 2 shows the XRD patterns of the as-SPSed Tie22Al-23.9Nb1.1Mo alloy and those after direct ageing. The as-SPSed alloy was in single B2 phase, and ageing induced the precipitation of O/a2 phase. The SEM images of the aged alloys is shown in Fig. 3. Only O phase generated in the alloys aged at 750e850 C, i.e. in O þ B2 phase region, and the O phase was in the lath shape, forming Wid€tten O þ B2 structure (Fig. 3aec). The rest alloys aged at mansta 900e1000 C, i.e. in Oþa2þB2 phase region, were composed of B2, a2, and O phases, and the O/a2 phase was rod-like, randomly distributed on the B2 matrix (Fig. 3def). The O phase in the alloys directly aged in the B2þO phase region is fine and dense, while that in the alloys aged in three-phase region is coarse and sparse.
Fig. 2. XRD patterns of the as-SPSed Tie22Al-23.9Nb-1.1Mo alloy and the alloys directly aged at 750 Ce1000 C.
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Fig. 3. SEM images of the Tie22Al-23.9Nb-1.1Mo alloys directly aged at (a) 750 C, (b) 800 C, (c) 850 C, (d) 900 C, (e) 950 C, and (f) 1000 C.
Therefore, phase transformation of B2/O þ B2 occurred when the Mo-modified Ti2AlNb alloys were aged in the O þ B2 phase region, which agrees with the observation in the Mo-free Ti2AlNb alloys aged in the same phase regions [27]. Moreover, the absence of the characteristic peaks of O phase at 2q ¼ 40.2 , 42.0 , and 54.5 indicated that ageing at 1000 C triggered the phase transformation of O/a2þB2 [28], and the intensity of peaks for a2 and B2 increased. The hardness of the aged alloys is shown in Fig. 4. The Mo-modified Ti2AlNb alloys aged in the O þ B2 phase region with €tten O laths showed the hardness of about 420 HV, Widmansta higher than those aged in the Oþa2þB2 phase region with coarse rod-like O. However, the hardness was lower than that of the Mofree ones (450 HV) aged in the same region [27], which is indicated in Fig. 4. With aim of regulating the microstructure for high hardness performance, the solution treatment was carried out before ageing for the SPSed Tie22Al-23.9Nb-1.1Mo alloys. Fig. 5 shows the XRD patterns of the Tie22Al-23.9Nb-1.1Mo
Fig. 5. XRD patterns of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at 750 Ce1000 C.
Fig. 4. Vickers hardness of the Tie22Al-23.9Nb-1.1Mo alloys directly aged at 750 Ce1000 C.
alloys by solution treatment and ageing. The alloy contained B2 as the main phase after solution treatment, and O phase became dominated after ageing in the O þ B2 phase region. Similar to the directly aged alloys, only the O phase precipitated when the ageing temperature was below 850 C, while the O/a2 phase precipitated in the alloys aged above 850 C. The a2 phase began to precipitate when the alloy was aged at 850 C, lower than the onset precipitation temperature of a2 (900 C) for the directly aged alloys. The SEM images of the Tie22Al-23.9Nb-1.1Mo alloys after solution treatment and ageing are shown in Fig. 6. Complete O þ B2 colonies, also reported as parallel O lamellae [9], were observed in the alloys aged at 750e850 C, indicating the phase transformation of B2/O þ B2 occurred during ageing. Specially, the colonies were well arranged in the same direction, differed from the structure of the directly aged alloy. When the ageing temperature was higher than 850 C, rod-like O phase precipitated from the B2 matrix, which is analogous to the directly aged alloys. The dispersion of the
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Fig. 6. SEM images of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at (a) 750 C, (b) 800 C, (c) 850 C, (d) 900 C, (e) 950 C, and (f) 1000 C.
O phase also shows similar tendency to that in the directly aged alloys. For a certain ageing temperature, the O-phase in the alloy experienced solution treatment is coarser and sparser than that in the directly aged one. Fig. 7 shows the hardness of the Tie22Al23.9Nb-1.1Mo alloys after solution treatment and ageing process. The hardness was significantly enhanced in contrast with that of the directly aged alloys. The alloys aged in the O þ B2 phase region exhibited the hardness of over 600 HV, and those aged in the Oþa2þB2 phase region of about 420 HV. 4. Discussion The hardness is associated with the content and the dimension of the O/a2 phase in the Ti2AlNb alloys. The volume fraction of the B2 and a2/O phases was calculated from the peak intensity of the XRD patterns in Figs. 2 and 5, and the corresponding values are collected in Fig. 8. The alloys aged in the O þ B2 phase region were
Fig. 8. vol fraction of the Tie22Al-23.9Nb-1.1Mo alloys experienced direct ageing and those experienced solution treatment and ageing.
Fig. 7. Vickers hardness of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at 750 Ce1000 C.
dominated by O phase, while those aged in the Oþa2þB2 phase region contained more B2 phase, which accorded with the previous results for the Ti2AlNb-based alloy containing Mo [29]. It is suggested the phase transformation of O/B2 or O/a2 occurred when the alloys were aged in the three-phase region [25]. Overall, those O-phase-dominated alloys exhibited higher microhardness than the B2-phase-dominated ones (Fig. 4). For the Tie22Al-23.9Nb1.1Mo alloys, the content of the O phase was about 80 vol% when the alloys were directly aged at 750e850 C. As the ageing temperature increased, the phase transformation of O/B2 or O/a2 proceeded by consuming the O phase that brought precipitation hardening effect. The alloys aged at 900e1000 C contained about 20 vol% O/a2 phase, fewer than the alloys aged at 750e850 C, which should be the main reason for the depression of the
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hardness. Similar tendency is observed in the illustration of hardness for the alloys experienced solution treatment and ageing (Fig. 7). Those aged in the B2þO phase region containing more O phase exhibited higher hardness than those aged in the Oþa2þB2 phase region. It is then compared the content of O/a2 phase in the directly aged alloys and those experienced solution treatment and ageing. For those aged at 750 C, the one experienced solution treatment contained 22% higher content of O phase than the directly aged one. The corresponding enhancement of hardness is about 47%. For those solution treated and aged at 800 C and 850 C, the hardness is enhanced by 51% and 47%; however, the content of O phase is decreased by 1% and 21%, respectively, in contrast with the directly aged ones. In this regard, the shape and width of the O phase might provide positive effects on the microhardness. The width of the lath and rod-like O is illustrated in Fig. 9. The average width of the O laths was about 50 nm for the Mo-modified alloys aged in the O þ B2 phase region, and it increased with the elevated ageing temperature to a limited degree. Compared with the values shown in Fig. 9 for the Mo-free alloys [27] experienced €tten O laths was refined the same ageing treatment, the Widmansta by Mo addition, which is due to the stabilization of B2 by Mo element that inhibited the B2/O þ B2. The refinement of O phase could be another reason for the enhancement of hardness for the Tie22Al-23.9Nb-1.1Mo alloys [11,18]. When the alloys were directly aged in the three-phase region, the width of the O laths became larger as the ageing temperature increased, and the value was 0.35 mm for the alloy aged at 1000 C, which should be responsible for the poor hardness performance (smaller than 380 HV). For the Tie22Al-23.9Nb-1.1Mo alloys experienced solution treatment and aged at 750e850 C, the O laths were further refined in contrast with those of the directly aged alloys, and the hardness of these alloys was significantly improved, up to 620 HV. For the alloys aged at 750 C and 800 C, the ones experienced solution treatment contained the O-phase laths with 39% and 58% smaller width than the directly aged ones, respectively. The corresponding enhancement of hardness is about 47% and 51%. Since the O-phase content is decreased for the alloy solution treated and aged at 800 C (66.5 vol%), in contrast with the directly aged one (67.9 vol%), the increase of the hardness is dominated by the reduction of the lath width. For the alloys aged at 850 C, the content of O phase is decreased by 21%, and the lath width is increased by 2% for the one experienced solution treatment and ageing. These should result in
Fig. 10. XRD patterns of the SPSed Tie22Ale25Nb alloy, the SPSed Tie22Al-23.9Nb1.1Mo alloy, and the solution treated Tie22Al-23.9Nb-1.1Mo alloy.
the decrease of the hardness. On the contrary, the hardness is enhanced by 47%. Therefore, the lattice distortion and dislocations are characterized for these Mo-modified alloys. We have attributed the refinement of O phase to the solution of Mo in the B2 phase, which inhibited the phase transformation of B2/O þ B2. In return, the solution of Mo lowered the formation temperature of O/B2, since the initial ageing temperature for the formation of a2 phase was at 900 C for the directly aged alloy, while the temperature was 850 C for the solution treated and aged alloy (Fig. 8). The evidence for the solution of Mo could be traced from the XRD patterns of the SPSed Tie22Ale25Nb alloy, the SPSed Tie22Al-23.9Nb-1.1Mo alloy, and the solution treated Tie22Al23.9Nb-1.1Mo alloy (Fig. 10). Since the atomic radius of Mo is smaller than that of Nb, the Mo substitution for Nb will reduce the volume of the lattice [30]. For the SPSed Tie22Al-23.9Nb-1.1Mo alloy, the (200) peak of B2 phase shifted to the high-angle direction, while the (110) and (211) peaks remained at the same position. Moreover, it was suggested spark plasma sintering could only achieve a limited amount of Mo substitution for Nb, and the initial stage of this substitution would induce the lattice distortion along the (200) lattice planes. When the SPSed Tie22Al-23.9Nb-1.1Mo alloy experienced solution treatment, an increasing amount of Mo substituted the Nb sites in the B2 lattice, as the (110), (200), and (211) peaks of the B2 phase shifted to the high-angle direction, and the deviation is 0.11, 0.25 , and 0.17, respectively. The interplanar spacing and the lattice parameters of the B2 phase is calculated by Rietveld refinement, as shown in Table 1. The Mo addition hardly decreased the lattice parameter and interplanar spacing during SPS, indicating limited Mo atoms have entered the lattice of B2 phase. After solution treatment, the lattice parameter a and the interplanar spacing of (110), (200), and (211) faces were obviously decreased, suggesting the Mo substitution that reduced the volume
Table 1 Lattice parameter of B2 phase and interplanar spacing of (110), (200), and (211) faces for Ti2AlNb, Mo-modified Ti2AlNb, and Mo-modified Ti2AlNb solution treated at 1300 C.
Fig. 9. Width of the laths for the aged Tie22Al-23.9Nb-1.1Mo alloys without solution treatment and those experienced solution treatment and ageing.
Sample
a/Å
d(110)/Å
d(200)/Å
d(211)/Å
Ti2AlNb Mo-modified Ti2AlNb Mo-modified Ti2AlNb, 1300 C
3.2482 3.2481 3.2454
2.2968 2.2964 2.2948
1.6241 1.6238 1.6227
1.3260 1.3258 1.3249
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Table 2 Lattice parameter of B2 phase and interplanar spacing of (110), (200), and (211) faces for MoeTi2AlNb alloys experienced different heat treatment. Mo-modified Ti2AlNb 750 C 800 C 850 C
Ageing Solution treatment þ ageing Ageing Solution treatment þ ageing Ageing Solution treatment þ ageing
a/Å
d(110)/Å
d(200)/Å
d(211)/Å
3.2778 3.2630 3.2805 3.2798 3.2689 3.2676
2.3178 2.3073 2.3193 2.3192 2.3115 2.3105
1.6389 1.6315 1.6400 1.6399 1.6345 1.6338
1.3382 1.3321 1.3391 1.3390 1.3346 1.3340
of the unit cell. Due to the anisotropy of the B2 lattice, the (200) and (211) planes suffered larger distortion than the (110) plane. Hence, it is evidenced the solution of Mo in the B2 lattice, which is likely to further affect the structure of the remained B2 phase and the precipitation of O phase. Table 2 shows the lattice parameter of B2 phase for MoeTi2AlNb alloys experienced different ageing treatment, as well as the interplanar spacing of (110), (200), and (211) faces. Overall, the values of lattice parameter a for the alloys experienced solution treatment and ageing is smaller than those for the alloys directly aged at the same temperature, which is caused by the remained Mo atoms in the lattice. The decrease of the interplanar spacing also provided evidence for the incorporation of Mo. However, it only decreased by 104 Å for the alloys aged at 800 C. This suggested that the lattice distortion may bring few effects on the enhancement of hardness for those aged at 800 C. As far as the alloys aged at 750 C and 850 C, Fig. 11 shows the TEM images of the directly aged Mo-modified Ti2AlNb alloys and those experienced solution treatment and ageing. Few dislocations are observed in the directly aged samples, whereas a considerable number of dislocations entangled in the alloys experienced solution treatment and ageing. Hence, the Mo addition induced lattice distortion in the B2 phase, and an increasing number of dislocations formed and entangled within the B2 laths should be responsible to the enhancement of the hardness as well. Considering the reduced O-phase content and the increased width of O laths, the entangled dislocations would be the main reason for the enhancement of hardness in the Mo-modified Ti2AlNb alloys solution treated and aged at 850 C.
Fig. 11. TEM images of the Mo-modified Ti2AlNb alloys directly aged at (a) 750 C and (b) 850 C, and the alloys solution treated at 1300 C followed by ageing at (c) 750 C and (d) 850 C.
5. Conclusions In summary, the effects of Mo addition on the microstructure and hardness were investigated in the spark plasma sintered Tie22Al-23.9Nb-1.1Mo alloys. The Mo-modified Ti2AlNb alloy solution treated and aged in the O þ B2 phase region exhibited favorable hardness of about 620 HV. The enhancement of hardness resulted from the increased content of O phase, the reduced width of O laths, and the entangled dislocations, which are induced by the Mo substitution. The Mo substitution for Nb induced severe lattice distortion along the (200) and (211) crystal planes of the B2 phase, but less distortion along the (110) planes. This could be the reason for the entangled dislocations. Specifically, the enhancement of hardness is dominated by the reduced width of O laths at 800 C, and it is dominated by the entangled dislocations at 850 C. Acknowledgements The authors are grateful to the National Natural Science Foundation of China (granted No. 51804195 and 51822404), the National High Technology Research and Development Program of China (Granted No. 2015AA042504), and the China Postdoctoral Science Foundation Grant (No. 2019T120329) for financial support. References [1] D. Banerjee, T.K. Nandi, V.A. Joshi, A new ordered orthorhombic phase in a Ti3Al-Nb alloy, Acta Metall. 36 (1988) 871e882. [2] D. Banerjee, The intermetallic Ti2AlNb, Prog. Mater. Sci. 42 (1997) 135e158. [3] J. Kumpfert, Intermetallic alloys based on orthorhombic titanium aluminide, Adv. Eng. Mater. 3 (2001) 851e864. [4] C.J. Boehlert, B.S. Majumdar, V. Seetharaman, D.B. Miracle, Part I. The microstructural evolution in Ti-Al-NbO plus Bcc orthorhombic alloys, Metall. Mater. Trans. A 30 (1999) 2305e2323. [5] K. Muraleedharan, A.K. Gogia, T.K. Nandy, D. Banerjee, S. Lele, Transformations in a Ti-24AI-15Nb alloy: Part I. Phase equilibria and microstructure, Metall. Trans. A 23 (1992) 401e415. [6] N.V. Kazantseva, S.V. Lepikhin, Study of the Ti-Al-Nb phase diagram, Phys. Met. Metallogr. 102 (2006) 169e180. [7] H.Z. Niu, Y.F. Chen, D.L. Zhang, Y.S. Zhang, J.W. Lu, W. Zhang, P.X. Zhang, Fabrication of a powder metallurgy Ti2AlNb-based alloy by spark plasma sintering and associated microstructure optimization, Mater. Des. 89 (2016) 823e829. [8] J. Wu, R.P. Guo, L. Xu, Z.G. Lu, Y.Y. Cui, R. Yang, Effect of hot isostatic pressing loading route on microstructure and mechanical properties of powder metallurgy Ti2AlNb alloys, J. Mater. Sci. Technol. 33 (2017) 172e178. [9] Y.P. Zheng, W.D. Zeng, D. Li, J.W. Xu, X. Ma, X.B. Liang, J.W. Zhang, Orthorhombic precipitate variant selection in a Ti2AlNb based alloy, Mater. Des. 158 (2018) 46e61. [10] J.L. Yang, G.F. Wang, W.C. Zhang, W.Z. Chen, X.Y. Jiao, K.F. Zhang, Microstructure evolution and mechanical properties of P/M Ti-22Al-25Nb alloy during hot extrusion, Mater. Sci. Eng., A 699 (2017) 210e216. [11] K.H. Sim, G.F. Wang, R.C. Son, S.L. Choe, Influence of mechanical alloying on the microstructure and mechanical properties of powder metallurgy Ti2AlNbbased alloy, Powder Technol. 317 (2017) 133e141. [12] Y. Mao, S.Q. Li, J.W. Zhang, J.H. Peng, D.X. Zou, Z.Y. Zhong, Microstructure and tensile properties of orthorhombic Ti-Al-Nb-Ta alloys, Intermetallics 8 (2000) 659e662. [13] M. Hagiwara, S. Emura, A. Araoka, B.O. Kong, F. Tang, Enhanced mechanical properties of orthorhombic Ti2AlNb-based intermetallic alloy, Met. Mater. Int. 9 (2003) 265e272. [14] X. Li, X. Wen, H. Zhao, Z. Ma, L. Yu, C. Li, C. Liu, Q. Guo, Y. Liu, The formation and evolution mechanism of amorphous layer surrounding Nb nano-grains in
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(atomic percent) alloys, Metall. Mater. Trans. A 38a (2007) 529e536. [23] J.W. Zhang, C.S. Lee, D.X. Zou, S.Q. Li, J.K.L. Lai, Microstructure and creep behavior of an orthorhombic Ti-25Al-17Nb-1Mo alloy, Metall. Mater. Trans. A 29 (1998) 559e564. [24] H.Z. Zhao, B. Lu, M. Tong, R. Yang, Tensile behavior of Ti-22Al-24Nb-0.5Mo in the range 25-650 oC, Mater. Sci. Eng., A 679 (2017) 455e464. [25] X.Y. Jiao, G. Liu, D.J. Wang, Y. Wu, Creep behavior and effects of heat treatment on creep resistance of Ti-22Al-24Nb-0.5Mo alloy, Mater. Sci. Eng., A 680 (2017) 182e189. [26] V. Raghavan, Al-Nb-Ti (Aluminum-Niobium-Titanium), J. Phase Equilibria Diffusion 26 (2005) 360e368. [27] M.C. Li, Q. Cai, Y.C. Liu, Z.Q. Ma, Z.M. Wang, Y. Huang, H.J. Li, Formation of fine B2/beta plus O structure and enhancement of hardness in the aged Ti2AlNbbased alloys prepared by spark plasma sintering, Metall. Mater. Trans. A 48a (2017) 4365e4371. [28] B. Shao, Y.Y. Zong, D.S. Wen, Y.T. Tian, D.B. Shan, Investigation of the phase transformations in Ti-22Al-25Nb alloy, Mater. Char. 114 (2016) 75e78. [29] A.A. Popov, A.G. Illarionov, S.V. Grib, S.L. Demakov, M.S. Karabanalov, O.A. Elkina, Phase and structural transformations in the alloy on the basis of the orthorhombic titanium aluminide, Phys. Met. Metallogr. 106 (2008) 399e410. [30] Y. Mao, M. Hagiwara, S. Emura, Creep behavior and tensile properties of Moand Fe-added orthorhombic Tie22Ale11Nbe2Moe1Fe alloy, Scr. Mater. 57 (2007) 261e264.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Solution treatment for enhanced hardness in Mo-modified Ti2AlNb-based alloys Yaran Zhang a, Qi Cai b, Zongqing Ma a, *, Chong Li a, Liming Yu a, Yongchang Liu a, ** a
State Key Laboratory of Hydraulic Engineering Simulation and Safety, School of Materials Science and Engineering, Tianjin University, Tianjin, 300354, China b Beijing Inst Technol, Sch Mat Sci & Engn, Beijing, 100081, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 21 March 2019 Received in revised form 11 July 2019 Accepted 15 July 2019 Available online 16 July 2019
The present work investigates the change in the microstructure and the improvement in the micro€tten hardness in the Mo-modified Ti2AlNb alloys induced by solution treatment. Typical Widmansta O þ B2 structure was observed in the alloys directly aged in the O þ B2 phase region, with a lower hardness value compared to that of the alloy without Mo addition. Well-arranged O þ B2 colonies were observed in the specimen solution treated in single B2 phase region, and it exhibited a favorable hardness performance of up to 620 HV. Besides the precipitation hardening of the O phase, the solution of Mo in B2 phase should be responsible for the enhancement of hardness. Further microstructural observation proved that the substitution of Mo for Nb in the lattice of B2 phase induced the lattice distortion mainly the (200) and (211) planes, and the induced entanglement of dislocations also contributed to the enhancement of hardness. © 2019 Elsevier B.V. All rights reserved.
Keywords: Ti2AlNb alloys Mo addition Solution treatment Microstructure Hardness
1. Introduction Intermetallic Ti2AlNb alloys, based on the orthorhombic (O) phase, have exhibited great potential for application in the components of engine, due to the excellent high-temperature properties. Typical Ti2AlNb-based alloys include Tie25Ale17Nb, Tie22Ale25Nb, and Tie23Ale27Nb, and these alloys containing high amounts of O phase have shown favorable mechanical properties, e.g. higher strength, and excellent creep resistance, compared with the conventional Ti3Al alloys [1e3]. Apart from the O phase, the Ti2AlNb-based alloys may also contain B2 phase (bodycentered cubic, bcc) and a2 phase (DO19 structure, Ti3Al) [4e6], and the precipitation strengthening of the O/a2 phase is currently the main concern for enhancement of the mechanical properties. For instance, fine acicular O/a2 phase was confirmed to be favorable to the tensile strength, whereas coarse O/a2 phase might improve the ductility of the Ti2AlNb-based alloys [7]. To further explore the possible application of Ti2AlNb-based alloys in aerospace, systematic investigation has been performed on the precipitation behavior of O/a2 phase, including the volume fraction, shape, and dimension,
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Z. Ma), [email protected] (Y. Liu). https://doi.org/10.1016/j.jallcom.2019.07.149 0925-8388/© 2019 Elsevier B.V. All rights reserved.
which are strongly related to the processing technique and heat treatment [8e11]. Adding b/a alloying elements is a convenient route to regulate the microstructure of the Ti2AlNb alloys. Mao et al. [12] have reported the substitution of Ta for Nb improved the strength and ductility of the Ti2AlNb alloy, and the yield strength of the deformed Tie22Ale20Nbe7Ta alloy reached 1200 MPa. Hagiwara et al. [13] produced W-modified Tie22Ale27Nb alloys with excellent high-temperature strength and creep properties. Alloying procedure is conveniently accomplished by directly adding the element powder into pre-alloyed powder, i.e. a powder metallurgic route, which makes it widely employed to adjust the composition of various alloys [14e17], Moreover, the outstanding advantage of powder metallurgy also lies in the direct fabrication of Ti2AlNb alloys with desirable shape, microstructure and properties. Yang et al. [18] have performed the W alloying by powder metal€tten structure was significantly refined, lurgy, and the Widmansta which was responsible to the enhancement of the hardness for the W-modified Ti2AlNb-based alloys. In addition to the refinement of €tten structure, diverse O þ B2 structures have been the Widmansta observed in the Ti2AlNb-based alloys. Emura et al. [19] have observed ‘Van Gogh's Sky’ structure with high elongation-to-failure in the annealed Mo- and Fe-modified Tie25Ale14Nb alloys. Yang et al. [20] have obtained herringbone O þ B2 structure by ageing the W-modified Ti2AlNb alloys in the O þ B2 phase region, and the
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alloys exhibited high hardness of 750 HV. The orientation relationships between the three phases are normally (0001)a2// (011)B2 and (110)B2//(001)O [4,21], but additional orientation may occur to form special structure. Studies also revealed that Mo addition could reduce the contiguity of a2 phase, which lowered the creep rate of the Ti2AlNbbased alloy [22]. The alloy with lath structure showed higher creep resistance than the three-phase microstructure with equiaxed phase [23]. As far as the mechanical properties of Mo-modified Ti2AlNb-based alloys, Zhao et al. [24] have investigated the tensile behavior of Tie22Ale24Nb-0.5Mo in the range of 25e650 C. The alloy with a2þB2 possessed the largest elongation, and the one with O/a2þB2 structure has the optimal balance between strength and plasticity. Jiao et al. [25] studied the microstructure and creep behavior of Tie22Ale24Nb-0.5Mo alloy. The creep resistance was significantly improved by eliminating coarse O þ B2 laths. However, the microhardness of the alloys, associated with the microstructure, has not been systematically investigated by solution treatment and ageing in different phase region; and moreover, the improvement or depression of hardness in Mo-added Ti2AlNb alloys was to date barely investigated from the viewpoint of lattice distortion. In this study, the sintered Mo-modified Ti2AlNb-based alloys are heat treated following different routes, either directly aged in O þ B2 and Oþa2þB2 phase regions, or solution treated in single B2 phase region and subsequent aged in O þ B2 and Oþa2þB2 phase regions. Different O þ B2 microstructure was obtained via these heat treatment process, and the resultant hardness performance was evaluated. Based on measuring the dimension and volume faction of O/a2, and investigating the lattice distortion by Mo substitution for Nb, the strengthening mechanism of the O/a2 phase was analyzed. 2. Experimental details Molybdenum-modified Ti2AlNb alloys, with the nominal composition of Tie22Al-23.9Nb-1.1Mo (at.%), were spark plasma sintered from spherical Tie22Ale25Nb pre-alloyed powder (~200 mm in size, single B2 phase) and Mo powder (~75 mm in diameter, 99.99% in purity) at 1100 C for 10 min and under 40 MPa. Two types of heat treatment were carried out on the as-SPSed alloys. The obtained alloys were either directly aged in different phase regions, or experienced solution treatment and subsequent ageing in different regions. The direct ageing was at 750 C, 800 C, 850 C, 900 C, 950 C and 1000 C for 2 h, respectively. The solution treatment was carried out at 1300 C for 6 h, followed by water quenching, and the subsequent ageing was at 750 C, 800 C, 850 C, 900 C, 950 C and 1000 C for 2 h, respectively, followed by air cooling. Fig. 1 shows the vertical section of the AleNbeTi phase diagram at 22 at.% Al [26], in which the phase regions for ageing are indicated. A pipe furnace (SK-G06123K) was used for the heat treatment after the samples were sealed in quartz tubes to avoid oxidation. The Mo-free Tie22Ale25Nb alloy were also aged at 850 C for 2 h without solution treatment as a control sample. The phase composition of the Ti2AlNb-based alloys were characterized by X-ray diffraction (XRD, Bruker D8 Advanced) using Cu Ka radiation, and the microstructure were observed by scanning electron microscopy (SEM, Hitachi S-4800), respectively. The alloys were chemically etched using Kroll's reagent. The orientation relation of the B2 and O/a2 was confirmed by transmission electron microscopy operated at 200 kV (TEM, JEOL-2100f). The roomtemperature Vickers microhardness was obtained under 0.2 kg load. The room temperature tensile tests were conducted on the electronic universal testing machine (Instron 5848) with the strain rate was 0.04 min1。
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Fig. 1. Vertical section of the AleNbeTi phase diagram at 22 at.% Al. Data are taken from Ref. [26].
3. Results Fig. 2 shows the XRD patterns of the as-SPSed Tie22Al-23.9Nb1.1Mo alloy and those after direct ageing. The as-SPSed alloy was in single B2 phase, and ageing induced the precipitation of O/a2 phase. The SEM images of the aged alloys is shown in Fig. 3. Only O phase generated in the alloys aged at 750e850 C, i.e. in O þ B2 phase region, and the O phase was in the lath shape, forming Wid€tten O þ B2 structure (Fig. 3aec). The rest alloys aged at mansta 900e1000 C, i.e. in Oþa2þB2 phase region, were composed of B2, a2, and O phases, and the O/a2 phase was rod-like, randomly distributed on the B2 matrix (Fig. 3def). The O phase in the alloys directly aged in the B2þO phase region is fine and dense, while that in the alloys aged in three-phase region is coarse and sparse.
Fig. 2. XRD patterns of the as-SPSed Tie22Al-23.9Nb-1.1Mo alloy and the alloys directly aged at 750 Ce1000 C.
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Fig. 3. SEM images of the Tie22Al-23.9Nb-1.1Mo alloys directly aged at (a) 750 C, (b) 800 C, (c) 850 C, (d) 900 C, (e) 950 C, and (f) 1000 C.
Therefore, phase transformation of B2/O þ B2 occurred when the Mo-modified Ti2AlNb alloys were aged in the O þ B2 phase region, which agrees with the observation in the Mo-free Ti2AlNb alloys aged in the same phase regions [27]. Moreover, the absence of the characteristic peaks of O phase at 2q ¼ 40.2 , 42.0 , and 54.5 indicated that ageing at 1000 C triggered the phase transformation of O/a2þB2 [28], and the intensity of peaks for a2 and B2 increased. The hardness of the aged alloys is shown in Fig. 4. The Mo-modified Ti2AlNb alloys aged in the O þ B2 phase region with €tten O laths showed the hardness of about 420 HV, Widmansta higher than those aged in the Oþa2þB2 phase region with coarse rod-like O. However, the hardness was lower than that of the Mofree ones (450 HV) aged in the same region [27], which is indicated in Fig. 4. With aim of regulating the microstructure for high hardness performance, the solution treatment was carried out before ageing for the SPSed Tie22Al-23.9Nb-1.1Mo alloys. Fig. 5 shows the XRD patterns of the Tie22Al-23.9Nb-1.1Mo
Fig. 5. XRD patterns of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at 750 Ce1000 C.
Fig. 4. Vickers hardness of the Tie22Al-23.9Nb-1.1Mo alloys directly aged at 750 Ce1000 C.
alloys by solution treatment and ageing. The alloy contained B2 as the main phase after solution treatment, and O phase became dominated after ageing in the O þ B2 phase region. Similar to the directly aged alloys, only the O phase precipitated when the ageing temperature was below 850 C, while the O/a2 phase precipitated in the alloys aged above 850 C. The a2 phase began to precipitate when the alloy was aged at 850 C, lower than the onset precipitation temperature of a2 (900 C) for the directly aged alloys. The SEM images of the Tie22Al-23.9Nb-1.1Mo alloys after solution treatment and ageing are shown in Fig. 6. Complete O þ B2 colonies, also reported as parallel O lamellae [9], were observed in the alloys aged at 750e850 C, indicating the phase transformation of B2/O þ B2 occurred during ageing. Specially, the colonies were well arranged in the same direction, differed from the structure of the directly aged alloy. When the ageing temperature was higher than 850 C, rod-like O phase precipitated from the B2 matrix, which is analogous to the directly aged alloys. The dispersion of the
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Fig. 6. SEM images of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at (a) 750 C, (b) 800 C, (c) 850 C, (d) 900 C, (e) 950 C, and (f) 1000 C.
O phase also shows similar tendency to that in the directly aged alloys. For a certain ageing temperature, the O-phase in the alloy experienced solution treatment is coarser and sparser than that in the directly aged one. Fig. 7 shows the hardness of the Tie22Al23.9Nb-1.1Mo alloys after solution treatment and ageing process. The hardness was significantly enhanced in contrast with that of the directly aged alloys. The alloys aged in the O þ B2 phase region exhibited the hardness of over 600 HV, and those aged in the Oþa2þB2 phase region of about 420 HV. 4. Discussion The hardness is associated with the content and the dimension of the O/a2 phase in the Ti2AlNb alloys. The volume fraction of the B2 and a2/O phases was calculated from the peak intensity of the XRD patterns in Figs. 2 and 5, and the corresponding values are collected in Fig. 8. The alloys aged in the O þ B2 phase region were
Fig. 8. vol fraction of the Tie22Al-23.9Nb-1.1Mo alloys experienced direct ageing and those experienced solution treatment and ageing.
Fig. 7. Vickers hardness of the Tie22Al-23.9Nb-1.1Mo alloys solution treated at 1300 C and aged at 750 Ce1000 C.
dominated by O phase, while those aged in the Oþa2þB2 phase region contained more B2 phase, which accorded with the previous results for the Ti2AlNb-based alloy containing Mo [29]. It is suggested the phase transformation of O/B2 or O/a2 occurred when the alloys were aged in the three-phase region [25]. Overall, those O-phase-dominated alloys exhibited higher microhardness than the B2-phase-dominated ones (Fig. 4). For the Tie22Al-23.9Nb1.1Mo alloys, the content of the O phase was about 80 vol% when the alloys were directly aged at 750e850 C. As the ageing temperature increased, the phase transformation of O/B2 or O/a2 proceeded by consuming the O phase that brought precipitation hardening effect. The alloys aged at 900e1000 C contained about 20 vol% O/a2 phase, fewer than the alloys aged at 750e850 C, which should be the main reason for the depression of the
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hardness. Similar tendency is observed in the illustration of hardness for the alloys experienced solution treatment and ageing (Fig. 7). Those aged in the B2þO phase region containing more O phase exhibited higher hardness than those aged in the Oþa2þB2 phase region. It is then compared the content of O/a2 phase in the directly aged alloys and those experienced solution treatment and ageing. For those aged at 750 C, the one experienced solution treatment contained 22% higher content of O phase than the directly aged one. The corresponding enhancement of hardness is about 47%. For those solution treated and aged at 800 C and 850 C, the hardness is enhanced by 51% and 47%; however, the content of O phase is decreased by 1% and 21%, respectively, in contrast with the directly aged ones. In this regard, the shape and width of the O phase might provide positive effects on the microhardness. The width of the lath and rod-like O is illustrated in Fig. 9. The average width of the O laths was about 50 nm for the Mo-modified alloys aged in the O þ B2 phase region, and it increased with the elevated ageing temperature to a limited degree. Compared with the values shown in Fig. 9 for the Mo-free alloys [27] experienced €tten O laths was refined the same ageing treatment, the Widmansta by Mo addition, which is due to the stabilization of B2 by Mo element that inhibited the B2/O þ B2. The refinement of O phase could be another reason for the enhancement of hardness for the Tie22Al-23.9Nb-1.1Mo alloys [11,18]. When the alloys were directly aged in the three-phase region, the width of the O laths became larger as the ageing temperature increased, and the value was 0.35 mm for the alloy aged at 1000 C, which should be responsible for the poor hardness performance (smaller than 380 HV). For the Tie22Al-23.9Nb-1.1Mo alloys experienced solution treatment and aged at 750e850 C, the O laths were further refined in contrast with those of the directly aged alloys, and the hardness of these alloys was significantly improved, up to 620 HV. For the alloys aged at 750 C and 800 C, the ones experienced solution treatment contained the O-phase laths with 39% and 58% smaller width than the directly aged ones, respectively. The corresponding enhancement of hardness is about 47% and 51%. Since the O-phase content is decreased for the alloy solution treated and aged at 800 C (66.5 vol%), in contrast with the directly aged one (67.9 vol%), the increase of the hardness is dominated by the reduction of the lath width. For the alloys aged at 850 C, the content of O phase is decreased by 21%, and the lath width is increased by 2% for the one experienced solution treatment and ageing. These should result in
Fig. 10. XRD patterns of the SPSed Tie22Ale25Nb alloy, the SPSed Tie22Al-23.9Nb1.1Mo alloy, and the solution treated Tie22Al-23.9Nb-1.1Mo alloy.
the decrease of the hardness. On the contrary, the hardness is enhanced by 47%. Therefore, the lattice distortion and dislocations are characterized for these Mo-modified alloys. We have attributed the refinement of O phase to the solution of Mo in the B2 phase, which inhibited the phase transformation of B2/O þ B2. In return, the solution of Mo lowered the formation temperature of O/B2, since the initial ageing temperature for the formation of a2 phase was at 900 C for the directly aged alloy, while the temperature was 850 C for the solution treated and aged alloy (Fig. 8). The evidence for the solution of Mo could be traced from the XRD patterns of the SPSed Tie22Ale25Nb alloy, the SPSed Tie22Al-23.9Nb-1.1Mo alloy, and the solution treated Tie22Al23.9Nb-1.1Mo alloy (Fig. 10). Since the atomic radius of Mo is smaller than that of Nb, the Mo substitution for Nb will reduce the volume of the lattice [30]. For the SPSed Tie22Al-23.9Nb-1.1Mo alloy, the (200) peak of B2 phase shifted to the high-angle direction, while the (110) and (211) peaks remained at the same position. Moreover, it was suggested spark plasma sintering could only achieve a limited amount of Mo substitution for Nb, and the initial stage of this substitution would induce the lattice distortion along the (200) lattice planes. When the SPSed Tie22Al-23.9Nb-1.1Mo alloy experienced solution treatment, an increasing amount of Mo substituted the Nb sites in the B2 lattice, as the (110), (200), and (211) peaks of the B2 phase shifted to the high-angle direction, and the deviation is 0.11, 0.25 , and 0.17, respectively. The interplanar spacing and the lattice parameters of the B2 phase is calculated by Rietveld refinement, as shown in Table 1. The Mo addition hardly decreased the lattice parameter and interplanar spacing during SPS, indicating limited Mo atoms have entered the lattice of B2 phase. After solution treatment, the lattice parameter a and the interplanar spacing of (110), (200), and (211) faces were obviously decreased, suggesting the Mo substitution that reduced the volume
Table 1 Lattice parameter of B2 phase and interplanar spacing of (110), (200), and (211) faces for Ti2AlNb, Mo-modified Ti2AlNb, and Mo-modified Ti2AlNb solution treated at 1300 C.
Fig. 9. Width of the laths for the aged Tie22Al-23.9Nb-1.1Mo alloys without solution treatment and those experienced solution treatment and ageing.
Sample
a/Å
d(110)/Å
d(200)/Å
d(211)/Å
Ti2AlNb Mo-modified Ti2AlNb Mo-modified Ti2AlNb, 1300 C
3.2482 3.2481 3.2454
2.2968 2.2964 2.2948
1.6241 1.6238 1.6227
1.3260 1.3258 1.3249
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Table 2 Lattice parameter of B2 phase and interplanar spacing of (110), (200), and (211) faces for MoeTi2AlNb alloys experienced different heat treatment. Mo-modified Ti2AlNb 750 C 800 C 850 C
Ageing Solution treatment þ ageing Ageing Solution treatment þ ageing Ageing Solution treatment þ ageing
a/Å
d(110)/Å
d(200)/Å
d(211)/Å
3.2778 3.2630 3.2805 3.2798 3.2689 3.2676
2.3178 2.3073 2.3193 2.3192 2.3115 2.3105
1.6389 1.6315 1.6400 1.6399 1.6345 1.6338
1.3382 1.3321 1.3391 1.3390 1.3346 1.3340
of the unit cell. Due to the anisotropy of the B2 lattice, the (200) and (211) planes suffered larger distortion than the (110) plane. Hence, it is evidenced the solution of Mo in the B2 lattice, which is likely to further affect the structure of the remained B2 phase and the precipitation of O phase. Table 2 shows the lattice parameter of B2 phase for MoeTi2AlNb alloys experienced different ageing treatment, as well as the interplanar spacing of (110), (200), and (211) faces. Overall, the values of lattice parameter a for the alloys experienced solution treatment and ageing is smaller than those for the alloys directly aged at the same temperature, which is caused by the remained Mo atoms in the lattice. The decrease of the interplanar spacing also provided evidence for the incorporation of Mo. However, it only decreased by 104 Å for the alloys aged at 800 C. This suggested that the lattice distortion may bring few effects on the enhancement of hardness for those aged at 800 C. As far as the alloys aged at 750 C and 850 C, Fig. 11 shows the TEM images of the directly aged Mo-modified Ti2AlNb alloys and those experienced solution treatment and ageing. Few dislocations are observed in the directly aged samples, whereas a considerable number of dislocations entangled in the alloys experienced solution treatment and ageing. Hence, the Mo addition induced lattice distortion in the B2 phase, and an increasing number of dislocations formed and entangled within the B2 laths should be responsible to the enhancement of the hardness as well. Considering the reduced O-phase content and the increased width of O laths, the entangled dislocations would be the main reason for the enhancement of hardness in the Mo-modified Ti2AlNb alloys solution treated and aged at 850 C.
Fig. 11. TEM images of the Mo-modified Ti2AlNb alloys directly aged at (a) 750 C and (b) 850 C, and the alloys solution treated at 1300 C followed by ageing at (c) 750 C and (d) 850 C.
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