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Microstructures and tensile properties of Ti2AlNb and Mo-modified Ti2AlNb alloys fabricated by hot isostatic pressing
Materials Science & Engineering A 776 (2020) 139043
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Microstructures and tensile properties of Ti2AlNb and Mo-modified Ti2AlNb alloys fabricated by hot isostatic pressing Yaran Zhang , Yongchang Liu *, Liming Yu , Hongyan Liang **, Yuan Huang , Zongqing Ma State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Ti2AlNb alloys Mo addition Mechanical properties Tensile strength Fracture
The Ti2AlNb and Mo-modified Ti2AlNb alloys were fabricated by hot isostatic pressing, and then underwent annealing treatment at 800 � C for 1, 2 and 3 h respectively. To evaluate the tensile properties of the two kinds of alloys, tensile tests were conducted at ambient temperature and 650 � C in this study. For the Ti2AlNb alloy, the room-temperature ultimate tensile strength (UTS) was 700–900 MPa, and the highest elongation to failure (εf) reached 6.2% after aging for 2 h at 800 � C. Since the Mo addition promoted the precipitation of the coarse B2þO colony structure near the grain boundary, the Mo-modified Ti2AlNb alloy exhibited decreased UTS and εf at room temperature. As for the high-temperature performance of both the Ti2AlNb and Mo-modified Ti2AlNb alloys, the UTS reached approximately 700 MPa and the εf exceeded 3.0%. Meanwhile, the two kinds of alloys presented the same tendency that both the high-temperature UTS and εf increased firstly and then decreased as the aging time increased. Among them the specimens after aging at 800 � C for 2 h for both alloys acquired the highest UTS and εf. Moreover, the Mo-modified alloy exhibited a better high-temperature ductility than that at room temperature, which is associated with the B2↔O phase transformation during the loading at 650 � C. The Mo addition effectively refined the lath of the precipitates in the grain and the grain size of the matrix, and it induced abundant dislocations. Thus, the Mo-modified Ti2AlNb alloy aged for 1 and 3 h exhibited superior hightemperature mechanical properties than those of the Ti2AlNb alloy.
1. Introduction The demand for lightweight high-temperature structural materials is urgently prompted by the manufacturing of advanced aeroengine [1–4]. Among these materials, the Ti2AlNb-based alloys were developed on the basis of intermetallic Ti3Al alloys by introducing a certain amount of Nb. The formation of an orthorhombic phase (O) results in excellent high-temperature strength and creep resistance, Hence, the Ti2AlNb alloys have received considerable attention as new potential structural materials in advanced gas turbine jet engine [5–9]. In addition to the O phase, the Ti2AlNb alloys may also contain body-centered cubic (B2) and Ti3Al (α2) phases, and particular microstructure is designed for specific performance by tuning the constitution of these three phases [10–13]. Although a lot of efforts have been made to improve the roomand high-temperature strength of the Ti2AlNb alloys in the past decades, the intrinsic brittleness has restricted their industrial applications [14–20]. Molybdenum has been used as a mature alloying element to modify
Ti2AlNb alloys, and the tensile properties and creep resistance of Momodified alloys have been mostly investigated on the basis of Ti–22Al–24Nb-0.5Mo [21–23]. It is found that Mo addition improves the creep resistance of Ti2AlNb alloys without sacrificing the ductility or yield strength [24], and Mo and Fe can improve the room-temperature ductility of Ti2AlNb alloys [25]. Lei et al. [26] investigated the origin of the enhanced mechanical properties in laser additive welded Ti2AlNb joints. A high Mo content induced equiaxed grains and high-angle grain boundaries, thus improving the tensile properties. Quast et al. compared the microstructure, tensile, and creep properties of Ti2AlNb alloys with different contents of Mo. The mechanical properties are associated with the volume fraction and contiguity of the α2 phase [27]. For the reason that Mo addition is a fascinating route for improving the ductility and enhancing the strength, it is essentially important to clarify the rela tionship between the mechanical performance and Mo microalloying. A cost-effective powder metallurgy technique, which is highly ad vantageous over casting and hot working considering the high utiliza tion of materials and free-designed shape of the workpieces, has been
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Liu), [email protected] (H. Liang). https://doi.org/10.1016/j.msea.2020.139043 Received 15 January 2020; Received in revised form 31 January 2020; Accepted 2 February 2020 Available online 4 February 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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microscopy (SEM, Hitachi S-4800 and JEOL 7800). The specimens were polished mechanically and then chemically etched using Kroll reagent. Detailed microstructures, such as dislocation and interfacial structures, were characterized by transmission electron microscopy (TEM, JEM2100f) operated at 200 kV. Room- and high-temperature (650 � C) ten sile tests were performed on the electronic universal testing machine (MTS C45.305) with a nominal strain rate of 10 4 s 1. The schematic graph of the tensile specimens is shown in Fig. 1. The overall length was 44 mm before testing, and the gauge length is 30 mm.
Fig. 1. Schematic diagram of the specimen for tensile test.
3. Results and discussion
used to fabricate Ti2AlNb alloys [28–32]. Wu et al. [28] investigated the influences of hot isostatic pressing (HIP) on the mechanical properties and microstructure of Mo-modified Ti2AlNb alloys. The HIP parameters were optimized to obtain homogenous microstructure and balanced mechanical properties. Recently, relative investigations about the microstructure evolution and mechanical properties were mostly con ducted on laser welded or linear friction welded Ti2AlNb alloys [26,29]. The room- and high-temperature tensile behaviors are mostly studied in Ti–22Al–24Nb-0.5Mo alloy fabricated by consumable vacuum arc re-melting [23]. However, the effects of heat treatment on the micro structure and mechanical properties of powder metallurgical Mo-modified Ti2AlNb alloys were barely mentioned. Hence, the effect of heat treatment on the room- and high-temperature mechanical proper ties of the powder metallurgical Mo-modified Ti2AlNb alloys need to be further explored. In the present study, the Ti2AlNb and Mo-modified Ti2AlNb alloys were fabricated via HIP sintering, the obtained samples were aged in the €tten O þ B2 O þ B2 phase region to obtain the typical Widmansta microstructure [17,33]. Room- and high-temperature tensile tests were performed to evaluate the mechanical properties of the sintered alloys, and the fracture morphologies were characterized. Moreover, the phase constituent, the lath dimension, and the dislocation distribution were investigated to elaborate the effect of Mo addition on the fracture behavior of the aged Ti2AlNb alloys.
Fig. 2 shows the recorded XRD patterns of the Ti2AlNb and Momodified Ti2AlNb alloys aged at 800 � C for 1, 2 and 3 h, respectively. The patterns indicated that the prepared alloys only contained B2 and O phases, and the O phase remained as the main phase with increase of the aging time. Mo addition hardly affected the phase constituent, but might have induced lattice distortion in the B2 phase. A similar phenomenon was observed in the spark plasma sintered Mo-modified Ti2AlNb alloy, and entangled dislocations were caused by such distortion [30]. In comparison with the (110) peak of B2 phase for the Ti2AlNb alloy (Fig. 2), the peak for the Mo-modified Ti2AlNb alloy shifted to a high-angle direction, whereas the peak positions of the O phase remained unchanged. This phenomenon can be attributed to the β-phase stabilizing element (Mo) which may replace Nb atoms in the B2 phase [34], thus decreasing the cell volume due to a smaller atomic radius of Mo than that of Nb. Subsequently, the interplanar spacing of the (110) face of the B2 phase was reduced. Moreover, the displacement of the B2 peak increased with the aging time increasing, which indicated that an
2. Experimental details Ti–22Al–25Nb and Ti–22Al-23.9Nb-1.1Mo (at.%) alloys were sin tered from Ti–22Al–25Nb pre-alloyed powder (~200 μm in particle size, single B2 phase) and Mo powder (~75 μm in diameter, 99.99% in pu rity) via HIP. The cylinder-shaped samples with the size of Φ30 � 80 mm were formed by sintering at 1150 � C for 3 h under 120 MPa. The applied heating and cooling rates were 6.5 � C⋅min 1 and 5.5 � C⋅min 1, respec tively. The obtained alloys were then aged at 800 � C for 1, 2 and 3 h respectively, followed by air cooling. X-ray diffraction (XRD) analysis was conducted with Cu Kα radiation to determine the phase constituent. Microstructural features were characterized via scanning electron
Fig. 3. Mass fraction of B2 and O phases in the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1, 2 and 3 h.
Fig. 2. Comparison of XRD patterns of the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for (a) 1 h, (b) 2 h, and (c) 3 h, respectively. 2
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Fig. 4. SEM images of the microstructures of the sintered samples: Ti2AlNb alloy aged at 800 � C for (a) 0 h, (b) 1 h, (c) 2 h and (d) 3 h; Mo-modified Ti2AlNb alloy aged at 800 � C for (e) 0 h, (f) 1 h, (g) 2 h and (h) 3 h.
increasing number of Mo atoms entered the B2 lattice during aging. In order to reveal the effect of Mo addition on the phase constituent during aging, the mass fraction (%) of the two phases was calculated and illustrated in Fig. 3 by using the Rietveld refinement method. For the Ti2AlNb alloy, the increasing mass fraction of O phase manifested that the phase transformation of B2→O occurred during aging. However, the Mo addition stabilized the B2 phase, which made the B2 phase hardly consumed as the aging proceeded with a stable mass fraction of approximately 36.0%. The initial microstructures of the sintered sam ples are shown in Fig. 4a and b. The average grain sizes of B2 phase for
the Ti2AlNb and the Mo-modified Ti2AlNb alloys are measured to be 466 � 45 μm and 148 � 15 μm, respectively. Fig. 4c–f presents the SEM images of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1, 2 and €tten structure was observed in the Ti2AlNb and Mo3 h. Full Widmansta modified Ti2AlNb alloys, whereas the homogeneity of the B2 grains is different. For the Ti2AlNb alloy, coarse and fine O þ B2 structures coexisted (in which the laths in light and dark colors represent the B2 phase and O phase, respectively). The average widths of the coarse O laths were 0.31 � 0.06, 0.42 � 0.08, and 0.42 � 0.07 μm for the Ti2AlNb alloy aged for 1, 2 and 3 h, respectively. By contrast, the Mo-modified 3
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segregation of Nb. On the basis of the energy dispersive spectrometer (EDS) measurement results, the Nb-rich region corresponds to the coarse structure, with a percentage of 26.8 � 1.7 at.%, and the fine laths contain 24.5 � 0.6 at.% Nb. Hence, Mo addition is beneficial for achieving the refinement and homogeneity of the microstructure within the grains, because there were barely coarse structure observed and €tten measured Nb content keeps around 24.5 at.% in the Widmansta structure of the Mo-modified Ti2AlNb alloy. It has been reported that the addition of 1 at.% Mo was capable of replacing 4.25 at.% Nb [36]. The substitution of Mo for Nb in the B2 lattice will generate some active Nb atoms with high diffusion rate, which promotes homogeneous distri bution of Nb in the Mo-modified Ti2AlNb alloy. Fig. 5 illustrates the measured room-temperature stress–strain curves of the Ti2AlNb and Mo-modified Ti2AlNb alloys. As seen from it, the UTS was larger than 700 MPa for all the Ti2AlNb alloy. The UTS (876 MPa) was acquired with an elongation to failure of 6.2% when the sintered Ti2AlNb alloy was aged for 2 h at 800 � C. By contrast, the Mo-modified Ti2AlNb alloy exhibited aggravating room-temperature UTS of 498, 481, and 380 MPa when aged for 1, 2, and 3 h, respectively, and their εf was uncompetitive with a value less than 2.5%. The Mo addition deterio rated the room-temperature mechanical performance of the Ti2AlNb alloys. Fig. 6 illustrates the room-temperature fracture morphologies of the Ti2AlNb and Mo-modified Ti2AlNb alloys. The fracture mode of the Ti2AlNb alloy was a cleavage type, wherein long river patterns caused by multiple slip systems extended along a typical cleavage fracture of the B2 phase [37]. The source of the “river” reflected that the crack initiated
Fig. 5. Measured room-temperature stress–strain curves for the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for 1, 2 and 3 h, respectively.
Ti2AlNb alloy displayed a more homogeneous O þ B2 structure, wherein the average widths of O laths were 0.25 � 0.05, 0.24 � 0.04, and 0.23 � 0.05 μm for those aged for 1, 2 and 3 h, respectively. It has been reported that Mo segregation resulted in curvy O þ B2 colonies [35], neverthe less, in this study no curvy O þ B2 colonies were observed, hence the Mo component was homogeneously distributed in the matrix. Elemental analysis indicated that the inhomogeneity originated from the
Fig. 6. SEM images of the fracture morphologies of the sintered samples after room-temperature tensile tests: Ti2AlNb aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3 h; Mo-modified Ti2AlNb aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
Fig. 7. SEM images indicating the differences along the grain boundary between the sintered (a) Ti2AlNb and (b) Mo-modified Ti2AlNb alloy after aging at 800 � C for 2 h. 4
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Fig. 8. TEM images of the microstructures of the sintered samples: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3 h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
from the grain boundary of the aged Ti2AlNb alloy. Short but crowded river patterns with dimples were observed in the Mo-modified Ti2AlNb alloy, showing a quasi-cleavage fracture. The increasing number of river patterns indicated that the Mo-modified alloy contained additional crack sources, which caused low strength and ductility. In the Mo-modified Ti2AlNb alloy, cracks initiated at the grain boundary as well. Therefore, the structure of the grain boundary of the Ti2AlNb and Mo-modified Ti2AlNb alloys were characterized. Fig. 7 shows the SEM images of the grain boundaries of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 2 h. It was found that a few O þ B2 colonies precipitated at the grain boundary of the Ti2AlNb alloy, whereas an increasing number of O þ B2 colonies precipitated at the grain boundary of the Mo-modified Ti2AlNb alloy. The phases precipitated at the grain boundary were detrimental to the mechanical properties of Ti2AlNb- based alloys [38], and the cracks preferred to initiate at the interface of B2 and O phase or along B2/O boundary [37,39]. This accounted for the low room-temperature strength and ductility of the Mo-modified Ti2AlNb alloys. Mao et al. [34] investigated the formation of colonies adjacent to grain boundary in the Mo- and Fe-added Ti2AlNb alloys, which was caused by the higher partition coefficient of Mo than that of Nb [40]. The strength is affected by the volume fractions of O phase, width of O lath, and substructure of dislocations [41]. The TEM images in Fig. 8 provide additional details about the microstructure of the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging. The aging process involved the movement and rearrangement of dislocations, and the density of dislo cations decreased as the aging time increased. The entanglement of
dislocations within the B2 phase in the Ti2AlNb alloy aged for 2 and 3 h, as indicated by arrows in the TEM images, significantly decreased (Fig. 8b and c) compared with that in the sample aged for 1 h (Fig. 8a). A similar phenomenon was observed in the Mo-modified Ti2AlNb alloy, as shown in Fig. 8d–f. For the Ti2AlNb alloy, the mass fraction increase of O phase was accompanied by the growth of O lath and the decrease of dislocation density. Precipitation strengthening then plays a key role in improving the tensile strength of the Ti2AlNb alloy aged for 2 h. For the Mo-modified Ti2AlNb alloy, although Mo addition induced further lat tice distortion with an increasing amount of dislocations based on the XRD patterns, the dislocation density decreased with the increasing aging time (Fig. 8d, e, and f). Therefore, the growth in the mass fraction of O phase and the refinement of O-lath accounted for the improved tensile strength for the alloy aged for 2 h. Fig. 9a and b show the TEM images of the Ti2AlNb alloy aged for 1 and 2 h. Stacking faults formed at the grain boundary, because the dislocations tended to be merged by the grain boundary. This condition indicated the movement of dislocations during aging. Cai et al. [42] reported that the high strength of alloy containing long O laths was associated with the B2/O interfaces, whereas the distortion at the interfaces might induce entangled dislo cations. The TEM images shown in Fig. 10 reveal the interface structure of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1 h. The atoms at the grain boundary of B2 and O phases were ordered in the Ti2AlNb alloy aged for 1 h (Fig. 10a). The introduced Mo atoms increased the density of dislocation and the grain boundary hindered the motion of additional dislocations, as a result dislocation pile-ups happened at the boundary. Disordered atoms were observed at the grain boundary
Fig. 9. TEM images of the sintered Ti2AlNb alloy after aging at 800 � C for (a) 1 h and (b) 2 h. 5
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Fig. 10. High-resolution TEM images of (a) Ti2AlNb and (b) Mo-modified Ti2AlNb alloy after aging at 800 � C for 1 h.
Fig. 11 shows the high-temperature stress–strain curves of the Ti2AlNb and Mo-modified Ti2AlNb alloys at 650 � C. The Ti2AlNb alloy aged for 3 h and all the Mo-modified Ti2AlNb alloy exhibited enhanced tensile strength and ductility at 650 � C compared with that at room temperature. Specifically, the high-temperature UTS was between 600 MPa and 700 MPa and the εf was between 3.0% and 4.0% for the Momodified Ti2AlNb alloy, whereas the high-temperature UTS ranged from 600 MPa to 900 MPa and the εf ranged from 0.5% to 5.0% for the Ti2AlNb alloy. The UTS and εf are comparable with the Ti–22Al–25Nb alloy fabricated via vacuum hot-pressing sintering [44] and the Ti–22Al–27Nb alloy prepared via laser beam welding [45]. Mo addition is beneficial to the high-temperature mechanical properties. Brittle fracture was confirmed in the Ti2AlNb and Mo-modified Ti2AlNb alloys, based on the fracture morphologies shown in Fig. 12. Grain size varia tion could also be observed from these SEM images. The grain of the Mo-modified Ti2AlNb alloy was refined compared with the Ti2AlNb alloy. Hence, grain refinement strengthening played a dominating role in improving the mechanical properties at high temperature for the Mo-modified Ti2AlNb alloy aged for 1 and 2 h. The highest UTS at 650 � C was recorded in the Ti2AlNb alloy aged for 2 h. Fig. 13 displays the microstructure of the specimens sectioned from the fractured high-temperature tensile specimens. The Ti2AlNb alloy €tten structure before the high temperature showed typical Widmansta tensile test (Fig. 4), and fine O þ B2 colonies with dendrite arms were observed after the tensile test, as indicated by circles in Fig. 13a–c. This indicated the reaction of B2→O þ B2 probably has occurred during the high-temperature tensile test. The Mo-modified Ti2AlNb alloy aged for
Fig. 11. Measured high-temperature stress–strain curves for the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for 1, 2 and 3 h, respectively.
between B2 and O phases in the Mo-modified Ti2AlNb alloy aged for 1 h, as indicated by arrows in Fig. 10b. Moreover, the initial orientation relationship of B2 and O phases in the Widmanst€ atten structure followed the crystallographic relationship of (110)B2//(002)O [43]. The angle between the B2 and O phases was 59� for the Mo-modified Ti2AlNb alloy aged for 1 h but only 35� for the Ti2AlNb alloy aged for 1 h, indicating that a large distortion was induced in the Mo-modified Ti2AlNb alloy.
Fig. 12. SEM images of the fracture morphologies (magnification: � 100) of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h. 6
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Fig. 13. SEM images of the cross-sectional microstructures of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
Fig. 14. SEM images of the fracture morphologies (magnification: � 1000) of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
1–3 h contained coarser B2 laths than the alloy before deformation. The average width were 0.27 � 0.07 μm, 0.26 � 0.06 μm, and 0.24 � 0.06 μm for the Mo-modified Ti2AlNb alloy aged for 1, 2 and 3 h, respectively. In addition, it was found that small equiaxed B2 phases precipitated in the Mo-modified Ti2AlNb alloy, as indicated by arrows in Fig. 13c–f. This implied the phase transformation of O→B2 might have occurred during the high-temperature tensile test. The low strength of Ti2AlNb alloy at high temperature was attributed to the reduced strength of grain boundary, and the reduction of B2 content caused the aggravation in high-temperature ductility. By contrast, the Mo-modified Ti2AlNb alloy exhibited better high-temperature ductility than room-temperature properties, due to the increasing amount of B2 phase, which provided additional slip systems. Gogia et al. [46] proved that the presence of β/B2 phase is essential in imparting ductility and toughness to this alloy system. Fig. 14 shows the fracture morphologies in the Ti2AlNb and Mo-modified Ti2AlNb alloys after high-temperature tensile test. Numerous tear ridges were observed in the Mo-modified Ti2AlNb alloy aged for 1 and 3 h (Fig. 14d and f), and these samples containing coarse laths exhibited higher tensile properties at 650 � C than the Ti2AlNb alloy aged for 1 and 3 h (Fig. 14a and c). For the Ti2AlNb and Mo-modified
Ti2AlNb alloy, the high-temperature UTS and elongation initially increased and then decreased with aging time increasing. Dimples appeared on the fracture surface of the Ti2AlNb alloy aged for 2 h (Fig. 14b), and this alloy exhibited the highest UTS of 917 MPa and εf of approximately 4.8%. When the Ti2AlNb alloy was aged for 3 h (Fig. 14c), the B2 laths were reduced and became globular. River pat terns emerged on the fracture surface, and this alloy showed decreased strength and elongation, in contrast with the one aged for 2 h. Coarse O þ B2 colonies formed in the Mo-modified Ti2AlNb alloy aged for 1, 2, and 3 h (Fig. 14d–f), and these samples had similar strength. The elon gation of the alloy aged for 2 h was higher due to the formation of B2 phase in a large area. 4. Conclusions The Ti2AlNb and Mo-modified Ti2AlNb alloys were prepared by hot isostatic pressing and annealed at 800 � C for 1, 2 and 3 h, respectively in this study. Tensile tests were conducted at ambient temperature and 650 � C. The main conclusions of this study are as follows:
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(1) The addition of Mo refined the lath of the O phase within the grain and the grain size of the B2 phase and induced abundant dislocations within the B2 phase. (2) The ultimate tensile strength and fracture elongation of the Ti2AlNb alloy were more favorable than those of the Mo-modified alloy at room temperature, whereas the high-temperature strength and elongation to failure became uncompetitive for the Ti2AlNb alloy compared with those for the Mo-modified Ti2AlNb alloy. (3) On the basis of the comparison of the O phase mass fraction, the width of O-phase laths, and the grain size of B2 phase, the roomtemperature tensile properties were determined through the precipitation strengthening of O phase, whereas the hightemperature tensile properties were dominated by the grain refinement strengthening of the B2 phase.
[16] B. Shao, S. Wan, D. Shan, B. Guo, Y. Zong, Hydrogen-induced improvement of the cylindrical drawing properties of a Ti-22Al-25Nb alloy, Adv. Eng. Mater. 19 (2017), 1600621. [17] M. Li, Q. Cai, Y. Liu, Z. Ma, Z. Wang, Microstructure and mechanical properties of Ti2AlNb-based Alloy synthesized by spark plasma sintering from pre-alloyed and ball-milled powder, Adv. Eng. Mater. 20 (2018), 1700659. [18] 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) 265–272. [19] J. Yang, Q. Cai, Z. Ma, Y. Huang, L. Yu, H. Li, Y. Liu, Effect of W addition on phase transformation and microstructure of powder metallurgic Ti-22Al-25Nb alloy during quenching and furnace cooling, Chin. J. Aeronaut. 32 (2019) 1343–1351. [20] W. Wang, W. Zeng, C. Xue, X. Liao, J. Zhang, Designed bimodal size lamellar O microstructures in Ti2AlNb based alloy: microstructural evolution, tensile and creep properties, Mater. Sci. Eng., A 618 (2014) 288–294. [21] L. Xu, R. Guo, J. Wu, Z. Lu, R. Yang, Progress in hot isostatic pressing technology of titanium alloy powder, Acta Metall. Sin. 54 (2018) 1537–1552. [22] X. Jiao, G. Liu, D. Wang, Y. Wu, Creep behavior and effects of heat treatment on creep resistance of Ti-22Al-24Nb-0.5 Mo alloy, Mater. Sci. Eng., A 680 (2017) 182–189. [23] H. Zhao, B. Lu, M. Tong, R. Yang, Tensile behavior of Ti-22Al-24Nb-0.5 Mo in the range 25-650� C, Mater. Sci. Eng., A 679 (2017) 455–464. [24] L. Germann, D. Banerjee, J. Gu� edou, J.-L. Strudel, Effect of composition on the mechanical properties of newly developed Ti2AlNb-based titanium aluminide, Intermetallics 13 (2005) 920–924. [25] S. Emura, K. Tsuzaki, K. Tsuchiya, Improvement of room temperature ductility for Mo and Fe modified Ti2AlNb alloy, Mater. Sci. Eng., A 528 (2010) 355–362. [26] Z. Lei, K. Zhang, H. Zhou, L. Ni, Y. Chen, A comparative study of microstructure and tensile properties of Ti2AlNb joints prepared by laser welding and laseradditive welding with the addition of filler powder, J. Mater. Process. Technol. 255 (2018) 477–487. [27] J. Quast, C. Boehlert, Comparison of the microstructure, tensile, and creep behavior for Ti-24Al-17Nb-0.66 Mo (atomic percent) and Ti-24Al-17Nb-2.3Mo (atomic percent) alloy, Metall. Mater. Trans. 38 (2007) 529–536. [28] J. Wu, R. Guo, L. Xu, Z. Lu, Y. Cui, R. Yang, Effect of hot isostatic pressing loading route on microstructure and mechanical properties of powder metallurgy Ti2AlNb alloy, J. Mater. Sci. Technol. 33 (2017) 172–178. [29] X. Chen, F. Xie, T. Ma, W. Li, X. Wu, Microstructure evolution and mechanical properties of linear friction welded Ti2AlNb alloy, J. Alloys Compd. 646 (2015) 490–496. [30] Y. Zhang, Q. Cai, Z. Ma, C. Li, L. Yu, Y. Liu, Solution treatment for enhanced hardness in Mo-modified Ti2AlNb-based alloy, J. Alloys Compd. 805 (2019) 1184–1190. [31] G. Wang, J. Yang, X. Jiao, Microstructure and mechanical properties of Ti-22Al25Nb alloy fabricated by elemental powder metallurgy, Mater. Sci. Eng., A 654 (2016) 69–76. [32] 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 Nb-Al system during mechanical alloying process, J. Alloys Compd. 779 (2019) 175–182. [33] M. Li, Q. Cai, Y. Liu, Z. Ma, Z. Wang, Y. Huang, H. Li, Formation of fine B2/βþO structure and enhancement of hardness in the aged Ti2AlNb-Based alloy prepared by spark plasma sintering, Metall. Mater. Trans. 48 (2017) 4365–4371. [34] Y. Mao, M. Hagiwara, S. Emura, Creep behavior and tensile properties of Mo and Fe added orthorhombic Ti-22Al-11Nb-2Mo-1Fe alloy, Scripta Mater. 57 (2007) 261–264. [35] Y. Zhang, Q. Cai, Y. Liu, Formation of diverse OþB2 structure and hardness of Momodified Ti-22Al-25Nb alloy upon cooling, Vacuum 165 (2019) 199–206. [36] F. Tang, S. Nakazawa, M. Hagiwara, The effect of quaternary additions on the microstructures and mechanical properties of orthorhombic Ti2AlNb-based alloy, Mater. Sci. Eng., A 329–331 (2002) 492–498. [37] K. Zhang, L. Ni, Z. Lei, Y. Chen, X. Hu, In situ investigation of the tensile deformation of laser welded Ti2AlNb joints, Mater. Char. 123 (2017) 51–57. [38] J. Kumpfert, Intermetallic alloy based on orthorhombic titanium aluminide, Adv. Eng. Mater. 3 (2001) 851–864. [39] W. Wang, W. Zeng, D. Li, B. Zhu, Y. Zheng, X. Liang, Microstructural evolution and tensile behavior of Ti2AlNb alloy based α2-phase decomposition, Mater. Sci. Eng., A 662 (2016) 120–128. [40] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria among α (hcp), β (bcc) and γ (L10) phases in Ti-Al base ternary alloy, Intermetallics 8 (2000) 855–867. [41] S. Emura, A. Araoka, M. Hagiwara, B2 grain size refinement and its effect on room temperature tensile properties of a Ti-22Al-27Nb orthorhombic intermetallic alloy, Scripta Mater. 48 (2003) 629–634. [42] Q. Cai, M. Li, Y. Zhang, Y. Liu, Z. Ma, C. Li, H. Li, Precipitation behavior of Widmanst€ atten O phase associated with interface in aged Ti2AlNb-based alloy, Mater. Char. 145 (2018) 413–422. [43] C. Boehlert, B. Majumdar, V. Seetharaman, D. Miracle, Part I. The microstructural evolution in Ti2AlNb OþBCC orthorhombic alloy, Metall. Mater. Trans. 30 (1999) 2305–2323.
CRediT authorship contribution statement Yaran Zhang: Conceptualization, Methodology, Investigation, Data curation, Software, Writing - original draft. Yongchang Liu: Concep tualization, Validation, Writing - review & editing, Resources. Liming Yu: Writing - review & editing. Hongyan Liang: Validation, Writing review & editing. Yuan Huang: Writing - review & editing. Zongqing Ma: Conceptualization, Validation, Writing - review & editing. Acknowledgements The authors are grateful to the National Natural Science Foundation of China (granted No. U1660201), the National High Technology Research and Development Program of China (Granted No. 2015AA042504) for financial support. References [1] K. Zhang, Z. Lei, Y. Chen, K. Yang, Y. Bao, Heat treatment of laser-additive welded Ti2AlNb joints: microstructure and tensile properties, Mater. Sci. Eng., A 744 (2019) 436–444. [2] X. Zhang, H. Li, M. Zhan, Z. Zheng, J. Gao, G. Shao, Electron force-induced dislocations annihilation and regeneration of a superalloy through electrical in-situ transmission electron microscopy observations, J. Mater. Sci. Technol. 36 (2020) 79–83. [3] J. Wu, Y.C. Liu, C. Li, Y.T. Wu, X.C. Xia, H.J. Li, Recent progress of microstructure evolution and performance of multiphase Ni3Al-based intermetallic alloy with high Fe and Cr contents, Acta Metall. Sin. 56 (2020) 21–35. [4] X. Li, G. Wang, Y. Gu, J. Yang, Electrically assisted diffusion bonding of Ti2AlNb alloy sheet using CP-Ti foil interlayer: microstructural characterization and mechanical tests, Mater. Sci. Eng., A 744 (2019) 733–745. [5] C. Koch, C. Liu, N. Stoloff, High-temperature ordered imermetallic alloy, MRS Syrup, Proc. 39 (1985) 351. [6] Z. Lu, J. Wu, L. Xu, X. Cui, R. Yang, Ring rolling forming and properties of Ti2AlNb special shaped ring prepared by powder metallurgy, Acta Metall. Sin. 55 (2019) 729–740. [7] J. Zhang, J. Liu, D. Xu, J. Wu, L. Xu, R. Yang, Characterization of the prior particle boundaries in a powder metallurgy Ti2AlNb alloy, J. Mater. Sci. Technol. 35 (2019) 2513–2525. [8] Y. Huang, Y. Liu, C. Li, Z. Ma, L. Yu, H. Li, Microstructure evolution and phase transformations in Ti-22Al-25Nb alloy tailored by super-transus solution treatment, Vacuum 161 (2019) 209–219. [9] H.P. Wang, P. Lü, X. Cai, B. Zhai, B. Wei, Rapid solidification kinetics and mechanical property characteristics of Ni–Zr eutectic alloys processed under electromagnetic levitation state, Mater. Sci. Eng., A 772 (2020), 138660. [10] D. Banerjee, A. Gogia, T. Nandi, V. Joshi, A new ordered orthorhombic phase in a Ti2AlNb alloy, Acta Metall. 36 (1988) 871–882. [11] D. Banerjee, The intermetallic Ti2AlNb, Prog. Mater. Sci. 42 (1997) 135–158. [12] B. Wu, M. Zinkevich, F. Aldinger, M. Chu, J. Shen, Prediction of the ordering behaviours of the orthorhombic phase based on Ti2AlNb alloy by combining thermodynamic model with ab initio calculation, Intermetallics 16 (2008) 42–51. [13] C. Xue, W. Zeng, W. Wang, X. Liang, J. Zhang, Quantitative analysis on microstructure evolution and tensile property for the isothermally forged Ti2AlNb based alloy during heat treatment, Mater. Sci. Eng., A 573 (2013) 183–189. [14] S. Wang, W. Xu, Y. Zong, X. Zhong, D. Shan, Effect of initial microstructures on hot deformation behavior and workability of Ti2AlNb-based alloy, Metals 8 (2018) 382. [15] Y. Wang, K. Zhang, B. Li, Microstructure and high temperature tensile properties of Ti-22Al-25Nb alloy prepared by reactive sintering with element powders, Mater. Sci. Eng., A 608 (2014) 229–233.
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[44] J. Jia, K. Zhang, S. Jiang, Microstructure and mechanical properties of Ti-22Al25Nb alloy fabricated by vacuum hot pressing sintering, Mater. Sci. Eng., A 616 (2014) 93–98. [45] Z. Lei, Z. Dong, Y. Chen, J. Zhang, R. Zhu, Microstructure and tensile properties of laser beam welded Ti-22Al-27Nb alloy, Mater. Des. 46 (2013) 151–156.
[46] A. Gogia, T. Nandy, D. Banerjee, T. Carisey, J. Strudel, J. Franchet, Microstructure and mechanical properties of orthorhombic alloy in the Ti2AlNb system, Intermetallics 6 (1998) 741–748.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructures and tensile properties of Ti2AlNb and Mo-modified Ti2AlNb alloys fabricated by hot isostatic pressing Yaran Zhang , Yongchang Liu *, Liming Yu , Hongyan Liang **, Yuan Huang , Zongqing Ma State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China
A R T I C L E I N F O
A B S T R A C T
Keywords: Ti2AlNb alloys Mo addition Mechanical properties Tensile strength Fracture
The Ti2AlNb and Mo-modified Ti2AlNb alloys were fabricated by hot isostatic pressing, and then underwent annealing treatment at 800 � C for 1, 2 and 3 h respectively. To evaluate the tensile properties of the two kinds of alloys, tensile tests were conducted at ambient temperature and 650 � C in this study. For the Ti2AlNb alloy, the room-temperature ultimate tensile strength (UTS) was 700–900 MPa, and the highest elongation to failure (εf) reached 6.2% after aging for 2 h at 800 � C. Since the Mo addition promoted the precipitation of the coarse B2þO colony structure near the grain boundary, the Mo-modified Ti2AlNb alloy exhibited decreased UTS and εf at room temperature. As for the high-temperature performance of both the Ti2AlNb and Mo-modified Ti2AlNb alloys, the UTS reached approximately 700 MPa and the εf exceeded 3.0%. Meanwhile, the two kinds of alloys presented the same tendency that both the high-temperature UTS and εf increased firstly and then decreased as the aging time increased. Among them the specimens after aging at 800 � C for 2 h for both alloys acquired the highest UTS and εf. Moreover, the Mo-modified alloy exhibited a better high-temperature ductility than that at room temperature, which is associated with the B2↔O phase transformation during the loading at 650 � C. The Mo addition effectively refined the lath of the precipitates in the grain and the grain size of the matrix, and it induced abundant dislocations. Thus, the Mo-modified Ti2AlNb alloy aged for 1 and 3 h exhibited superior hightemperature mechanical properties than those of the Ti2AlNb alloy.
1. Introduction The demand for lightweight high-temperature structural materials is urgently prompted by the manufacturing of advanced aeroengine [1–4]. Among these materials, the Ti2AlNb-based alloys were developed on the basis of intermetallic Ti3Al alloys by introducing a certain amount of Nb. The formation of an orthorhombic phase (O) results in excellent high-temperature strength and creep resistance, Hence, the Ti2AlNb alloys have received considerable attention as new potential structural materials in advanced gas turbine jet engine [5–9]. In addition to the O phase, the Ti2AlNb alloys may also contain body-centered cubic (B2) and Ti3Al (α2) phases, and particular microstructure is designed for specific performance by tuning the constitution of these three phases [10–13]. Although a lot of efforts have been made to improve the roomand high-temperature strength of the Ti2AlNb alloys in the past decades, the intrinsic brittleness has restricted their industrial applications [14–20]. Molybdenum has been used as a mature alloying element to modify
Ti2AlNb alloys, and the tensile properties and creep resistance of Momodified alloys have been mostly investigated on the basis of Ti–22Al–24Nb-0.5Mo [21–23]. It is found that Mo addition improves the creep resistance of Ti2AlNb alloys without sacrificing the ductility or yield strength [24], and Mo and Fe can improve the room-temperature ductility of Ti2AlNb alloys [25]. Lei et al. [26] investigated the origin of the enhanced mechanical properties in laser additive welded Ti2AlNb joints. A high Mo content induced equiaxed grains and high-angle grain boundaries, thus improving the tensile properties. Quast et al. compared the microstructure, tensile, and creep properties of Ti2AlNb alloys with different contents of Mo. The mechanical properties are associated with the volume fraction and contiguity of the α2 phase [27]. For the reason that Mo addition is a fascinating route for improving the ductility and enhancing the strength, it is essentially important to clarify the rela tionship between the mechanical performance and Mo microalloying. A cost-effective powder metallurgy technique, which is highly ad vantageous over casting and hot working considering the high utiliza tion of materials and free-designed shape of the workpieces, has been
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Y. Liu), [email protected] (H. Liang). https://doi.org/10.1016/j.msea.2020.139043 Received 15 January 2020; Received in revised form 31 January 2020; Accepted 2 February 2020 Available online 4 February 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.
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microscopy (SEM, Hitachi S-4800 and JEOL 7800). The specimens were polished mechanically and then chemically etched using Kroll reagent. Detailed microstructures, such as dislocation and interfacial structures, were characterized by transmission electron microscopy (TEM, JEM2100f) operated at 200 kV. Room- and high-temperature (650 � C) ten sile tests were performed on the electronic universal testing machine (MTS C45.305) with a nominal strain rate of 10 4 s 1. The schematic graph of the tensile specimens is shown in Fig. 1. The overall length was 44 mm before testing, and the gauge length is 30 mm.
Fig. 1. Schematic diagram of the specimen for tensile test.
3. Results and discussion
used to fabricate Ti2AlNb alloys [28–32]. Wu et al. [28] investigated the influences of hot isostatic pressing (HIP) on the mechanical properties and microstructure of Mo-modified Ti2AlNb alloys. The HIP parameters were optimized to obtain homogenous microstructure and balanced mechanical properties. Recently, relative investigations about the microstructure evolution and mechanical properties were mostly con ducted on laser welded or linear friction welded Ti2AlNb alloys [26,29]. The room- and high-temperature tensile behaviors are mostly studied in Ti–22Al–24Nb-0.5Mo alloy fabricated by consumable vacuum arc re-melting [23]. However, the effects of heat treatment on the micro structure and mechanical properties of powder metallurgical Mo-modified Ti2AlNb alloys were barely mentioned. Hence, the effect of heat treatment on the room- and high-temperature mechanical proper ties of the powder metallurgical Mo-modified Ti2AlNb alloys need to be further explored. In the present study, the Ti2AlNb and Mo-modified Ti2AlNb alloys were fabricated via HIP sintering, the obtained samples were aged in the €tten O þ B2 O þ B2 phase region to obtain the typical Widmansta microstructure [17,33]. Room- and high-temperature tensile tests were performed to evaluate the mechanical properties of the sintered alloys, and the fracture morphologies were characterized. Moreover, the phase constituent, the lath dimension, and the dislocation distribution were investigated to elaborate the effect of Mo addition on the fracture behavior of the aged Ti2AlNb alloys.
Fig. 2 shows the recorded XRD patterns of the Ti2AlNb and Momodified Ti2AlNb alloys aged at 800 � C for 1, 2 and 3 h, respectively. The patterns indicated that the prepared alloys only contained B2 and O phases, and the O phase remained as the main phase with increase of the aging time. Mo addition hardly affected the phase constituent, but might have induced lattice distortion in the B2 phase. A similar phenomenon was observed in the spark plasma sintered Mo-modified Ti2AlNb alloy, and entangled dislocations were caused by such distortion [30]. In comparison with the (110) peak of B2 phase for the Ti2AlNb alloy (Fig. 2), the peak for the Mo-modified Ti2AlNb alloy shifted to a high-angle direction, whereas the peak positions of the O phase remained unchanged. This phenomenon can be attributed to the β-phase stabilizing element (Mo) which may replace Nb atoms in the B2 phase [34], thus decreasing the cell volume due to a smaller atomic radius of Mo than that of Nb. Subsequently, the interplanar spacing of the (110) face of the B2 phase was reduced. Moreover, the displacement of the B2 peak increased with the aging time increasing, which indicated that an
2. Experimental details Ti–22Al–25Nb and Ti–22Al-23.9Nb-1.1Mo (at.%) alloys were sin tered from Ti–22Al–25Nb pre-alloyed powder (~200 μm in particle size, single B2 phase) and Mo powder (~75 μm in diameter, 99.99% in pu rity) via HIP. The cylinder-shaped samples with the size of Φ30 � 80 mm were formed by sintering at 1150 � C for 3 h under 120 MPa. The applied heating and cooling rates were 6.5 � C⋅min 1 and 5.5 � C⋅min 1, respec tively. The obtained alloys were then aged at 800 � C for 1, 2 and 3 h respectively, followed by air cooling. X-ray diffraction (XRD) analysis was conducted with Cu Kα radiation to determine the phase constituent. Microstructural features were characterized via scanning electron
Fig. 3. Mass fraction of B2 and O phases in the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1, 2 and 3 h.
Fig. 2. Comparison of XRD patterns of the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for (a) 1 h, (b) 2 h, and (c) 3 h, respectively. 2
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Fig. 4. SEM images of the microstructures of the sintered samples: Ti2AlNb alloy aged at 800 � C for (a) 0 h, (b) 1 h, (c) 2 h and (d) 3 h; Mo-modified Ti2AlNb alloy aged at 800 � C for (e) 0 h, (f) 1 h, (g) 2 h and (h) 3 h.
increasing number of Mo atoms entered the B2 lattice during aging. In order to reveal the effect of Mo addition on the phase constituent during aging, the mass fraction (%) of the two phases was calculated and illustrated in Fig. 3 by using the Rietveld refinement method. For the Ti2AlNb alloy, the increasing mass fraction of O phase manifested that the phase transformation of B2→O occurred during aging. However, the Mo addition stabilized the B2 phase, which made the B2 phase hardly consumed as the aging proceeded with a stable mass fraction of approximately 36.0%. The initial microstructures of the sintered sam ples are shown in Fig. 4a and b. The average grain sizes of B2 phase for
the Ti2AlNb and the Mo-modified Ti2AlNb alloys are measured to be 466 � 45 μm and 148 � 15 μm, respectively. Fig. 4c–f presents the SEM images of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1, 2 and €tten structure was observed in the Ti2AlNb and Mo3 h. Full Widmansta modified Ti2AlNb alloys, whereas the homogeneity of the B2 grains is different. For the Ti2AlNb alloy, coarse and fine O þ B2 structures coexisted (in which the laths in light and dark colors represent the B2 phase and O phase, respectively). The average widths of the coarse O laths were 0.31 � 0.06, 0.42 � 0.08, and 0.42 � 0.07 μm for the Ti2AlNb alloy aged for 1, 2 and 3 h, respectively. By contrast, the Mo-modified 3
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segregation of Nb. On the basis of the energy dispersive spectrometer (EDS) measurement results, the Nb-rich region corresponds to the coarse structure, with a percentage of 26.8 � 1.7 at.%, and the fine laths contain 24.5 � 0.6 at.% Nb. Hence, Mo addition is beneficial for achieving the refinement and homogeneity of the microstructure within the grains, because there were barely coarse structure observed and €tten measured Nb content keeps around 24.5 at.% in the Widmansta structure of the Mo-modified Ti2AlNb alloy. It has been reported that the addition of 1 at.% Mo was capable of replacing 4.25 at.% Nb [36]. The substitution of Mo for Nb in the B2 lattice will generate some active Nb atoms with high diffusion rate, which promotes homogeneous distri bution of Nb in the Mo-modified Ti2AlNb alloy. Fig. 5 illustrates the measured room-temperature stress–strain curves of the Ti2AlNb and Mo-modified Ti2AlNb alloys. As seen from it, the UTS was larger than 700 MPa for all the Ti2AlNb alloy. The UTS (876 MPa) was acquired with an elongation to failure of 6.2% when the sintered Ti2AlNb alloy was aged for 2 h at 800 � C. By contrast, the Mo-modified Ti2AlNb alloy exhibited aggravating room-temperature UTS of 498, 481, and 380 MPa when aged for 1, 2, and 3 h, respectively, and their εf was uncompetitive with a value less than 2.5%. The Mo addition deterio rated the room-temperature mechanical performance of the Ti2AlNb alloys. Fig. 6 illustrates the room-temperature fracture morphologies of the Ti2AlNb and Mo-modified Ti2AlNb alloys. The fracture mode of the Ti2AlNb alloy was a cleavage type, wherein long river patterns caused by multiple slip systems extended along a typical cleavage fracture of the B2 phase [37]. The source of the “river” reflected that the crack initiated
Fig. 5. Measured room-temperature stress–strain curves for the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for 1, 2 and 3 h, respectively.
Ti2AlNb alloy displayed a more homogeneous O þ B2 structure, wherein the average widths of O laths were 0.25 � 0.05, 0.24 � 0.04, and 0.23 � 0.05 μm for those aged for 1, 2 and 3 h, respectively. It has been reported that Mo segregation resulted in curvy O þ B2 colonies [35], neverthe less, in this study no curvy O þ B2 colonies were observed, hence the Mo component was homogeneously distributed in the matrix. Elemental analysis indicated that the inhomogeneity originated from the
Fig. 6. SEM images of the fracture morphologies of the sintered samples after room-temperature tensile tests: Ti2AlNb aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3 h; Mo-modified Ti2AlNb aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
Fig. 7. SEM images indicating the differences along the grain boundary between the sintered (a) Ti2AlNb and (b) Mo-modified Ti2AlNb alloy after aging at 800 � C for 2 h. 4
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Fig. 8. TEM images of the microstructures of the sintered samples: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3 h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
from the grain boundary of the aged Ti2AlNb alloy. Short but crowded river patterns with dimples were observed in the Mo-modified Ti2AlNb alloy, showing a quasi-cleavage fracture. The increasing number of river patterns indicated that the Mo-modified alloy contained additional crack sources, which caused low strength and ductility. In the Mo-modified Ti2AlNb alloy, cracks initiated at the grain boundary as well. Therefore, the structure of the grain boundary of the Ti2AlNb and Mo-modified Ti2AlNb alloys were characterized. Fig. 7 shows the SEM images of the grain boundaries of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 2 h. It was found that a few O þ B2 colonies precipitated at the grain boundary of the Ti2AlNb alloy, whereas an increasing number of O þ B2 colonies precipitated at the grain boundary of the Mo-modified Ti2AlNb alloy. The phases precipitated at the grain boundary were detrimental to the mechanical properties of Ti2AlNb- based alloys [38], and the cracks preferred to initiate at the interface of B2 and O phase or along B2/O boundary [37,39]. This accounted for the low room-temperature strength and ductility of the Mo-modified Ti2AlNb alloys. Mao et al. [34] investigated the formation of colonies adjacent to grain boundary in the Mo- and Fe-added Ti2AlNb alloys, which was caused by the higher partition coefficient of Mo than that of Nb [40]. The strength is affected by the volume fractions of O phase, width of O lath, and substructure of dislocations [41]. The TEM images in Fig. 8 provide additional details about the microstructure of the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging. The aging process involved the movement and rearrangement of dislocations, and the density of dislo cations decreased as the aging time increased. The entanglement of
dislocations within the B2 phase in the Ti2AlNb alloy aged for 2 and 3 h, as indicated by arrows in the TEM images, significantly decreased (Fig. 8b and c) compared with that in the sample aged for 1 h (Fig. 8a). A similar phenomenon was observed in the Mo-modified Ti2AlNb alloy, as shown in Fig. 8d–f. For the Ti2AlNb alloy, the mass fraction increase of O phase was accompanied by the growth of O lath and the decrease of dislocation density. Precipitation strengthening then plays a key role in improving the tensile strength of the Ti2AlNb alloy aged for 2 h. For the Mo-modified Ti2AlNb alloy, although Mo addition induced further lat tice distortion with an increasing amount of dislocations based on the XRD patterns, the dislocation density decreased with the increasing aging time (Fig. 8d, e, and f). Therefore, the growth in the mass fraction of O phase and the refinement of O-lath accounted for the improved tensile strength for the alloy aged for 2 h. Fig. 9a and b show the TEM images of the Ti2AlNb alloy aged for 1 and 2 h. Stacking faults formed at the grain boundary, because the dislocations tended to be merged by the grain boundary. This condition indicated the movement of dislocations during aging. Cai et al. [42] reported that the high strength of alloy containing long O laths was associated with the B2/O interfaces, whereas the distortion at the interfaces might induce entangled dislo cations. The TEM images shown in Fig. 10 reveal the interface structure of the Ti2AlNb and Mo-modified Ti2AlNb alloys aged for 1 h. The atoms at the grain boundary of B2 and O phases were ordered in the Ti2AlNb alloy aged for 1 h (Fig. 10a). The introduced Mo atoms increased the density of dislocation and the grain boundary hindered the motion of additional dislocations, as a result dislocation pile-ups happened at the boundary. Disordered atoms were observed at the grain boundary
Fig. 9. TEM images of the sintered Ti2AlNb alloy after aging at 800 � C for (a) 1 h and (b) 2 h. 5
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Fig. 10. High-resolution TEM images of (a) Ti2AlNb and (b) Mo-modified Ti2AlNb alloy after aging at 800 � C for 1 h.
Fig. 11 shows the high-temperature stress–strain curves of the Ti2AlNb and Mo-modified Ti2AlNb alloys at 650 � C. The Ti2AlNb alloy aged for 3 h and all the Mo-modified Ti2AlNb alloy exhibited enhanced tensile strength and ductility at 650 � C compared with that at room temperature. Specifically, the high-temperature UTS was between 600 MPa and 700 MPa and the εf was between 3.0% and 4.0% for the Momodified Ti2AlNb alloy, whereas the high-temperature UTS ranged from 600 MPa to 900 MPa and the εf ranged from 0.5% to 5.0% for the Ti2AlNb alloy. The UTS and εf are comparable with the Ti–22Al–25Nb alloy fabricated via vacuum hot-pressing sintering [44] and the Ti–22Al–27Nb alloy prepared via laser beam welding [45]. Mo addition is beneficial to the high-temperature mechanical properties. Brittle fracture was confirmed in the Ti2AlNb and Mo-modified Ti2AlNb alloys, based on the fracture morphologies shown in Fig. 12. Grain size varia tion could also be observed from these SEM images. The grain of the Mo-modified Ti2AlNb alloy was refined compared with the Ti2AlNb alloy. Hence, grain refinement strengthening played a dominating role in improving the mechanical properties at high temperature for the Mo-modified Ti2AlNb alloy aged for 1 and 2 h. The highest UTS at 650 � C was recorded in the Ti2AlNb alloy aged for 2 h. Fig. 13 displays the microstructure of the specimens sectioned from the fractured high-temperature tensile specimens. The Ti2AlNb alloy €tten structure before the high temperature showed typical Widmansta tensile test (Fig. 4), and fine O þ B2 colonies with dendrite arms were observed after the tensile test, as indicated by circles in Fig. 13a–c. This indicated the reaction of B2→O þ B2 probably has occurred during the high-temperature tensile test. The Mo-modified Ti2AlNb alloy aged for
Fig. 11. Measured high-temperature stress–strain curves for the Ti2AlNb and Mo-modified Ti2AlNb alloys after aging at 800 � C for 1, 2 and 3 h, respectively.
between B2 and O phases in the Mo-modified Ti2AlNb alloy aged for 1 h, as indicated by arrows in Fig. 10b. Moreover, the initial orientation relationship of B2 and O phases in the Widmanst€ atten structure followed the crystallographic relationship of (110)B2//(002)O [43]. The angle between the B2 and O phases was 59� for the Mo-modified Ti2AlNb alloy aged for 1 h but only 35� for the Ti2AlNb alloy aged for 1 h, indicating that a large distortion was induced in the Mo-modified Ti2AlNb alloy.
Fig. 12. SEM images of the fracture morphologies (magnification: � 100) of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h. 6
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Fig. 13. SEM images of the cross-sectional microstructures of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
Fig. 14. SEM images of the fracture morphologies (magnification: � 1000) of the sintered samples after high-temperature tensile tests: Ti2AlNb alloy aged at 800 � C for (a) 1 h, (b) 2 h and (c) 3h; Mo-modified Ti2AlNb alloy aged at 800 � C for (d) 1 h, (e) 2 h and (f) 3 h.
1–3 h contained coarser B2 laths than the alloy before deformation. The average width were 0.27 � 0.07 μm, 0.26 � 0.06 μm, and 0.24 � 0.06 μm for the Mo-modified Ti2AlNb alloy aged for 1, 2 and 3 h, respectively. In addition, it was found that small equiaxed B2 phases precipitated in the Mo-modified Ti2AlNb alloy, as indicated by arrows in Fig. 13c–f. This implied the phase transformation of O→B2 might have occurred during the high-temperature tensile test. The low strength of Ti2AlNb alloy at high temperature was attributed to the reduced strength of grain boundary, and the reduction of B2 content caused the aggravation in high-temperature ductility. By contrast, the Mo-modified Ti2AlNb alloy exhibited better high-temperature ductility than room-temperature properties, due to the increasing amount of B2 phase, which provided additional slip systems. Gogia et al. [46] proved that the presence of β/B2 phase is essential in imparting ductility and toughness to this alloy system. Fig. 14 shows the fracture morphologies in the Ti2AlNb and Mo-modified Ti2AlNb alloys after high-temperature tensile test. Numerous tear ridges were observed in the Mo-modified Ti2AlNb alloy aged for 1 and 3 h (Fig. 14d and f), and these samples containing coarse laths exhibited higher tensile properties at 650 � C than the Ti2AlNb alloy aged for 1 and 3 h (Fig. 14a and c). For the Ti2AlNb and Mo-modified
Ti2AlNb alloy, the high-temperature UTS and elongation initially increased and then decreased with aging time increasing. Dimples appeared on the fracture surface of the Ti2AlNb alloy aged for 2 h (Fig. 14b), and this alloy exhibited the highest UTS of 917 MPa and εf of approximately 4.8%. When the Ti2AlNb alloy was aged for 3 h (Fig. 14c), the B2 laths were reduced and became globular. River pat terns emerged on the fracture surface, and this alloy showed decreased strength and elongation, in contrast with the one aged for 2 h. Coarse O þ B2 colonies formed in the Mo-modified Ti2AlNb alloy aged for 1, 2, and 3 h (Fig. 14d–f), and these samples had similar strength. The elon gation of the alloy aged for 2 h was higher due to the formation of B2 phase in a large area. 4. Conclusions The Ti2AlNb and Mo-modified Ti2AlNb alloys were prepared by hot isostatic pressing and annealed at 800 � C for 1, 2 and 3 h, respectively in this study. Tensile tests were conducted at ambient temperature and 650 � C. The main conclusions of this study are as follows:
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(1) The addition of Mo refined the lath of the O phase within the grain and the grain size of the B2 phase and induced abundant dislocations within the B2 phase. (2) The ultimate tensile strength and fracture elongation of the Ti2AlNb alloy were more favorable than those of the Mo-modified alloy at room temperature, whereas the high-temperature strength and elongation to failure became uncompetitive for the Ti2AlNb alloy compared with those for the Mo-modified Ti2AlNb alloy. (3) On the basis of the comparison of the O phase mass fraction, the width of O-phase laths, and the grain size of B2 phase, the roomtemperature tensile properties were determined through the precipitation strengthening of O phase, whereas the hightemperature tensile properties were dominated by the grain refinement strengthening of the B2 phase.
[16] B. Shao, S. Wan, D. Shan, B. Guo, Y. Zong, Hydrogen-induced improvement of the cylindrical drawing properties of a Ti-22Al-25Nb alloy, Adv. Eng. Mater. 19 (2017), 1600621. [17] M. Li, Q. Cai, Y. Liu, Z. Ma, Z. Wang, Microstructure and mechanical properties of Ti2AlNb-based Alloy synthesized by spark plasma sintering from pre-alloyed and ball-milled powder, Adv. Eng. Mater. 20 (2018), 1700659. [18] 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) 265–272. [19] J. Yang, Q. Cai, Z. Ma, Y. Huang, L. Yu, H. Li, Y. Liu, Effect of W addition on phase transformation and microstructure of powder metallurgic Ti-22Al-25Nb alloy during quenching and furnace cooling, Chin. J. Aeronaut. 32 (2019) 1343–1351. [20] W. Wang, W. Zeng, C. Xue, X. Liao, J. Zhang, Designed bimodal size lamellar O microstructures in Ti2AlNb based alloy: microstructural evolution, tensile and creep properties, Mater. Sci. Eng., A 618 (2014) 288–294. [21] L. Xu, R. Guo, J. Wu, Z. Lu, R. Yang, Progress in hot isostatic pressing technology of titanium alloy powder, Acta Metall. Sin. 54 (2018) 1537–1552. [22] X. Jiao, G. Liu, D. Wang, Y. Wu, Creep behavior and effects of heat treatment on creep resistance of Ti-22Al-24Nb-0.5 Mo alloy, Mater. Sci. Eng., A 680 (2017) 182–189. [23] H. Zhao, B. Lu, M. Tong, R. Yang, Tensile behavior of Ti-22Al-24Nb-0.5 Mo in the range 25-650� C, Mater. Sci. Eng., A 679 (2017) 455–464. [24] L. Germann, D. Banerjee, J. Gu� edou, J.-L. Strudel, Effect of composition on the mechanical properties of newly developed Ti2AlNb-based titanium aluminide, Intermetallics 13 (2005) 920–924. [25] S. Emura, K. Tsuzaki, K. Tsuchiya, Improvement of room temperature ductility for Mo and Fe modified Ti2AlNb alloy, Mater. Sci. Eng., A 528 (2010) 355–362. [26] Z. Lei, K. Zhang, H. Zhou, L. Ni, Y. Chen, A comparative study of microstructure and tensile properties of Ti2AlNb joints prepared by laser welding and laseradditive welding with the addition of filler powder, J. Mater. Process. Technol. 255 (2018) 477–487. [27] J. Quast, C. Boehlert, Comparison of the microstructure, tensile, and creep behavior for Ti-24Al-17Nb-0.66 Mo (atomic percent) and Ti-24Al-17Nb-2.3Mo (atomic percent) alloy, Metall. Mater. Trans. 38 (2007) 529–536. [28] J. Wu, R. Guo, L. Xu, Z. Lu, Y. Cui, R. Yang, Effect of hot isostatic pressing loading route on microstructure and mechanical properties of powder metallurgy Ti2AlNb alloy, J. Mater. Sci. Technol. 33 (2017) 172–178. [29] X. Chen, F. Xie, T. Ma, W. Li, X. Wu, Microstructure evolution and mechanical properties of linear friction welded Ti2AlNb alloy, J. Alloys Compd. 646 (2015) 490–496. [30] Y. Zhang, Q. Cai, Z. Ma, C. Li, L. Yu, Y. Liu, Solution treatment for enhanced hardness in Mo-modified Ti2AlNb-based alloy, J. Alloys Compd. 805 (2019) 1184–1190. [31] G. Wang, J. Yang, X. Jiao, Microstructure and mechanical properties of Ti-22Al25Nb alloy fabricated by elemental powder metallurgy, Mater. Sci. Eng., A 654 (2016) 69–76. [32] 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 Nb-Al system during mechanical alloying process, J. Alloys Compd. 779 (2019) 175–182. [33] M. Li, Q. Cai, Y. Liu, Z. Ma, Z. Wang, Y. Huang, H. Li, Formation of fine B2/βþO structure and enhancement of hardness in the aged Ti2AlNb-Based alloy prepared by spark plasma sintering, Metall. Mater. Trans. 48 (2017) 4365–4371. [34] Y. Mao, M. Hagiwara, S. Emura, Creep behavior and tensile properties of Mo and Fe added orthorhombic Ti-22Al-11Nb-2Mo-1Fe alloy, Scripta Mater. 57 (2007) 261–264. [35] Y. Zhang, Q. Cai, Y. Liu, Formation of diverse OþB2 structure and hardness of Momodified Ti-22Al-25Nb alloy upon cooling, Vacuum 165 (2019) 199–206. [36] F. Tang, S. Nakazawa, M. Hagiwara, The effect of quaternary additions on the microstructures and mechanical properties of orthorhombic Ti2AlNb-based alloy, Mater. Sci. Eng., A 329–331 (2002) 492–498. [37] K. Zhang, L. Ni, Z. Lei, Y. Chen, X. Hu, In situ investigation of the tensile deformation of laser welded Ti2AlNb joints, Mater. Char. 123 (2017) 51–57. [38] J. Kumpfert, Intermetallic alloy based on orthorhombic titanium aluminide, Adv. Eng. Mater. 3 (2001) 851–864. [39] W. Wang, W. Zeng, D. Li, B. Zhu, Y. Zheng, X. Liang, Microstructural evolution and tensile behavior of Ti2AlNb alloy based α2-phase decomposition, Mater. Sci. Eng., A 662 (2016) 120–128. [40] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, Phase equilibria among α (hcp), β (bcc) and γ (L10) phases in Ti-Al base ternary alloy, Intermetallics 8 (2000) 855–867. [41] S. Emura, A. Araoka, M. Hagiwara, B2 grain size refinement and its effect on room temperature tensile properties of a Ti-22Al-27Nb orthorhombic intermetallic alloy, Scripta Mater. 48 (2003) 629–634. [42] Q. Cai, M. Li, Y. Zhang, Y. Liu, Z. Ma, C. Li, H. Li, Precipitation behavior of Widmanst€ atten O phase associated with interface in aged Ti2AlNb-based alloy, Mater. Char. 145 (2018) 413–422. [43] C. Boehlert, B. Majumdar, V. Seetharaman, D. Miracle, Part I. The microstructural evolution in Ti2AlNb OþBCC orthorhombic alloy, Metall. Mater. Trans. 30 (1999) 2305–2323.
CRediT authorship contribution statement Yaran Zhang: Conceptualization, Methodology, Investigation, Data curation, Software, Writing - original draft. Yongchang Liu: Concep tualization, Validation, Writing - review & editing, Resources. Liming Yu: Writing - review & editing. Hongyan Liang: Validation, Writing review & editing. Yuan Huang: Writing - review & editing. Zongqing Ma: Conceptualization, Validation, Writing - review & editing. Acknowledgements The authors are grateful to the National Natural Science Foundation of China (granted No. U1660201), the National High Technology Research and Development Program of China (Granted No. 2015AA042504) for financial support. References [1] K. Zhang, Z. Lei, Y. Chen, K. Yang, Y. Bao, Heat treatment of laser-additive welded Ti2AlNb joints: microstructure and tensile properties, Mater. Sci. Eng., A 744 (2019) 436–444. [2] X. Zhang, H. Li, M. Zhan, Z. Zheng, J. Gao, G. Shao, Electron force-induced dislocations annihilation and regeneration of a superalloy through electrical in-situ transmission electron microscopy observations, J. Mater. Sci. Technol. 36 (2020) 79–83. [3] J. Wu, Y.C. Liu, C. Li, Y.T. Wu, X.C. Xia, H.J. Li, Recent progress of microstructure evolution and performance of multiphase Ni3Al-based intermetallic alloy with high Fe and Cr contents, Acta Metall. Sin. 56 (2020) 21–35. [4] X. Li, G. Wang, Y. Gu, J. Yang, Electrically assisted diffusion bonding of Ti2AlNb alloy sheet using CP-Ti foil interlayer: microstructural characterization and mechanical tests, Mater. Sci. Eng., A 744 (2019) 733–745. [5] C. Koch, C. Liu, N. Stoloff, High-temperature ordered imermetallic alloy, MRS Syrup, Proc. 39 (1985) 351. [6] Z. Lu, J. Wu, L. Xu, X. Cui, R. Yang, Ring rolling forming and properties of Ti2AlNb special shaped ring prepared by powder metallurgy, Acta Metall. Sin. 55 (2019) 729–740. [7] J. Zhang, J. Liu, D. Xu, J. Wu, L. Xu, R. Yang, Characterization of the prior particle boundaries in a powder metallurgy Ti2AlNb alloy, J. Mater. Sci. Technol. 35 (2019) 2513–2525. [8] Y. Huang, Y. Liu, C. Li, Z. Ma, L. Yu, H. Li, Microstructure evolution and phase transformations in Ti-22Al-25Nb alloy tailored by super-transus solution treatment, Vacuum 161 (2019) 209–219. [9] H.P. Wang, P. Lü, X. Cai, B. Zhai, B. Wei, Rapid solidification kinetics and mechanical property characteristics of Ni–Zr eutectic alloys processed under electromagnetic levitation state, Mater. Sci. Eng., A 772 (2020), 138660. [10] D. Banerjee, A. Gogia, T. Nandi, V. Joshi, A new ordered orthorhombic phase in a Ti2AlNb alloy, Acta Metall. 36 (1988) 871–882. [11] D. Banerjee, The intermetallic Ti2AlNb, Prog. Mater. Sci. 42 (1997) 135–158. [12] B. Wu, M. Zinkevich, F. Aldinger, M. Chu, J. Shen, Prediction of the ordering behaviours of the orthorhombic phase based on Ti2AlNb alloy by combining thermodynamic model with ab initio calculation, Intermetallics 16 (2008) 42–51. [13] C. Xue, W. Zeng, W. Wang, X. Liang, J. Zhang, Quantitative analysis on microstructure evolution and tensile property for the isothermally forged Ti2AlNb based alloy during heat treatment, Mater. Sci. Eng., A 573 (2013) 183–189. [14] S. Wang, W. Xu, Y. Zong, X. Zhong, D. Shan, Effect of initial microstructures on hot deformation behavior and workability of Ti2AlNb-based alloy, Metals 8 (2018) 382. [15] Y. Wang, K. Zhang, B. Li, Microstructure and high temperature tensile properties of Ti-22Al-25Nb alloy prepared by reactive sintering with element powders, Mater. Sci. Eng., A 608 (2014) 229–233.
8
Y. Zhang et al.
Materials Science & Engineering A 776 (2020) 139043
[44] J. Jia, K. Zhang, S. Jiang, Microstructure and mechanical properties of Ti-22Al25Nb alloy fabricated by vacuum hot pressing sintering, Mater. Sci. Eng., A 616 (2014) 93–98. [45] Z. Lei, Z. Dong, Y. Chen, J. Zhang, R. Zhu, Microstructure and tensile properties of laser beam welded Ti-22Al-27Nb alloy, Mater. Des. 46 (2013) 151–156.
[46] A. Gogia, T. Nandy, D. Banerjee, T. Carisey, J. Strudel, J. Franchet, Microstructure and mechanical properties of orthorhombic alloy in the Ti2AlNb system, Intermetallics 6 (1998) 741–748.
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