α-Ti composite: DRX and globularization behavior

Journal of Alloys and Compounds 827 (2020) 154170

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Microstructure evolution during forging deformation of (TiBþTiCþY2O3)/a-Ti composite: DRX and globularization behavior Jianhui Yang a, Shulong Xiao a, *, Yuyong Chen a, b, **, Lijuan Xu a, Xiaopeng Wang a, Jing Tian a, Dongdong Zhang a, Zhuangzhuang Zheng a a b

School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China Panzhihua University, Panzhihua, 617000, China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 25 November 2019 Received in revised form 2 February 2020 Accepted 3 February 2020 Available online 8 February 2020

In this work, the b and (aþb) forging deformation were conducted on (TiB þ TiC þ Y2O3)/a-Ti composite to investigate the effect of reinforcement on dynamic recrystallization and globularization behavior. During b deformation, the reinforcements not only accelerated discontinuous dynamic recrystallization by particle stimulated nucleation mechanism, but also promoted the heterogeneous nucleation of a phase thus conducted the {0001}a pole figure of composite more widely distributed. However, during (aþb) deformation, the microstructure was characteristic of kinked a colonies morphology. Moreover, compared to matrix, the (TiB þ TiC þ Y2O3)/a-Ti composite depicted a more homogeneous microstructure and larger crystallography rotation both along or transverse to the a lamellar. The distribution of Schmid factor value indicated that the deformation behavior of matrix showed varies with the a colony while it was more homogeneous in composite. Reinforced with (TiB þ TiC þ Y2O3) reinforcements, on the one hand, the random SF value of composite conducted the incompatibility slip between adjacent a lamellar thus at last leaded the formation of HABs in inter-lamellar; on the other hand, the reinforcement itself could obstruct dislocation that accelerated the dynamic globularization behavior. Therefore, the (TiB þ TiC þ Y2O3)/a-Ti composite had a great advantage in optimizing microstructure especially globularization that made the composite of great application prospect. © 2020 Elsevier B.V. All rights reserved.

Keywords: Titanium matrix composites (TiBþTiCþY2O3) reinforcements Microstructure evolution Deformation behavior Dynamic recrystallization Globularization

1. Introduction Titanium alloys have an extensive applications own to its high strength-to-weight ratio, good resistance to many corrosive environments and high service temperature [1e3]. With the rapid development of aerospace technology, the traditional hightemperature titanium alloys, such as IMI834, Ti1100 and Ti60, can’t meet the demands of supersonic aircraft and structural parts. Recently, the titanium matrix composites (TMCs) reinforced by ceramic particles or whiskers provide a good opportunity to solve this unsolvable problem [4e6]. Especially the discontinuously reinforced titanium matrix composites (DRTMCs), which are fabricated by in situ technique, have an incomparable advantage

* Corresponding author. ** Corresponding author. School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China. E-mail addresses: [email protected] (S. Xiao), [email protected] (Y. Chen). https://doi.org/10.1016/j.jallcom.2020.154170 0925-8388/© 2020 Elsevier B.V. All rights reserved.

because of their low cost, easy-to-manufacture, good interfacial integrity and isotropic properties [7,8]. Deformation is an efficient and effective way to achieve good mechanical properties [9]. However, the control of microstructure and final mechanical properties are significantly sensitive to deformation process parameters [10]. Ma [11e13] obtained Widmanstatten microstructure and duplex microstructure (DP microstructure) through various thermal mechanical processes (TMPs) in TiC/Ti-1100 composite. Zhang [14] found that the as-rolled 2.5 vol% (TiBw þ TiCp)/Ti composites achieved highest strength and ductility under 85% reduction at RT and 600  C. Roy [15] investigated the influence of temperature and strain rate on deformation response and microstructural evolution during hot compression of a Tie6Ale4Ve0.1B alloy by isothermal hot compression. Besides, Gaisin [16] and Imayev [17] also devoted to study of microstructure and mechanical properties of discontinuously reinforced composite materials based on TiB/near a-Ti and (TiB þ TiC)/near a-Ti composites subjected to multiple hot forging. Consensus had been reached that the basketweave microstructure which was associated

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with final forging in b phase field possessed excellent strength and creep properties [18]. While the DP microstructure which was associated with final forging in (aþb) phase field was the desired microstructure for aero-engines own its good balance of fatigue and creep properties [10]. And globularization of a lamellar is the most significant microstructure evolution phenomenon as well as the goal of microstructure tailoring in the subtransus processing [19]. Generally speaking, the globularization mechanism included: boundary splitting and microstructure coarsening (termination migration and Ostwald ripening) [20,21]. The former was controlled by substructure while the later was effected by temperature and annealing time [19,21]. As we all know, lots of literatures were concerning on microstructure control and mechanical properties. But limited information was available so far on deformation behavior and deformation mechanism for TMCs. As Roy [22] pointed out that a trace boron could influent the b-b grain boundary and then had a direct effect on subsequent b/as phase transformation in as-cast Tie6Ale4Ve0.1B. Therefore, there were considerable different in deformation behavior between matrix and composite. In the present study, the electron back-scattered diffraction (EBSD) analysis was conducted to explore the deformation behavior and mechanism in (TiB þ TiC þ Y2O3)/near a-Ti composite. For comparison, the matrix was also introduced at same condition. 2. Materials and experimental procedure In the present study, a near a titanium alloy with a normal composition Tie6Ale4Sne7Zr-0.8Moe1Nbe1W-0.25Si (in wt.%) was referred as matrix. The composite was differed by 1.5 vol% TiB, 1.5 vol% TiC and 0.4 vol% Y2O3 which was referred as (TiB þ TiC þ Y2O3)/a-Ti composite. The ingot casting was obtained by vacuum arc melting furnace. As depicted in Fig. 1, the matrix was characteristic of Widmanstatten microstructure with a lamellar width of 1.5 mm (Fig. 1 a). The grain boundary a (aGB) with a width of 4 mm was embellished at prior b grain boundary. On the contrary, the microstructure of (TiB þ TiC þ Y2O3)/a-Ti composite was complex which was between Widmanstatten and basketweave microstructure. What’s more, the a lamellar was more equiaxedlike with a width of 4 mm compared to matrix. As depicted in Fig. 1 b, TiB particle was rod-like (or whisker like); TiC and Y2O3 particles were near equiaxed like (identified with arrows). Two cylinders with F55  80 mm were cut from central of each ingot for subsequent forging. The b transformation temperature of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite were about 990  C and 1010  C, respectively. The sample was hold for 2 h at 1100  C before forging due to eliminate casting stress and homogenize microstructure. Then the sample was keep the deformation temperature (1050  C & 950  C) 30min to homogenize temperature,

after that the cylinder was broken down at 1050  C (or 950  C) from the height of 80 mme12 mm with a strain rate of 0.5 s1. After deformation, the sample was annealing at 650  C for 2 h and subsequent cooling with the furnace. The specific thermos-mechanical process was illustrated in Fig. 2. For microstructural observation, the deformed specimens were sectioned parallel to their compression axes. Samples for scanning electron microscope (SEM) observation were cut from the forging cake then prepared with conventional metallographic technique using grounding and electropolished (600 ml methanolþ300 ml butoxyethanolþ100 ml perchloric acid), then they were etched with Kroll’s reagent (10 ml HF, 30 ml HNO3 and 500 ml H2O). The microstructures observation was conducted on JXA-8230. The sample prepared for electron backscatter diffraction (EBSD) analysis was same with microstructural observation without etching. The EBSD data acquisition and analysis were conducted on ZEISS SUPRA55 and HKL channel 5 software, respectively. Thin discs for TEM were prepared by electrical-discharge machining followed by mechanical grinding to a thickness of less than 30 mm prior to ion reduction. Transmission electron microscopy (TEM) micrographs observation were conducted on Talos F200x. 3. Results and discussion 3.1. General observations of microstructure evolution The typical microstructure of the deformed titanium alloy was depicted in Fig. 3. As one can see, the microstructure of b forging and (aþb) forging was definitely distinct. The b grain was elongated along the metal flow direction (i.e. perpendicular to compression

Fig. 2. Thermomechanical process flow chart of the matrix and (TiB þ TiC þ Y2O3)/a-Ti composite.

Fig. 1. Typical optical metallography of the as-cast matrix (a) and (TiB þ TiC þ Y2O3)/a-Ti composite (b).

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Fig. 3. BSEM micrographs of the matrix (a and c) and (TiB þ TiC þ Y2O3)/a-Ti composite (b and d) subjected to b forging (a and b) and (aþb) forging (c and d). The compression axis is pointed out in a.

axis). Specially, there were many equiaxed b grains like a necklace distributed at the prior b grain boundary. The equiaxed b grain was marked with black dotted line while the prior b grain boundary was marked with yellow dotted area in Fig. 3 a and b. According to Huang [23], those equiaxed b grains were identified as discontinuous dynamic recrystallization (dDRX) grains. Compared Fig. 3 a and b, the dDRX behavior of the (TiB þ TiC þ Y2O3)/a-Ti composite was more intense than that in matrix. In addition, the (TiB þ TiC þ Y2O3)/a-Ti composite was relatively more homogeneous and the a owns a lower length to width ratio. Besides, the (TiB þ TiC þ Y2O3) reinforcements are distributed along the metal flow direction, especially near those newly formed dDRX grains. However, the remarkable microstructure of (aþb) deformation at 950  C was characteristic of kinked a colonies (Fig. 3c and d). It was reported that the a lamellar had a tendency to along the metal flow direction under subtransus deformation and the kinked morphology feature was a result of flow softening that associated with plastic buckling of a laths and the loss of Hall-Petch strengthening from the a/b interfaces [24]. Thus for those a colonies parallel to metal flow direction had little morphological change while those a colonies vertical to metal flow direction tended to become kinked morphology. Obviously, the interface between a and b laths were no longer straight and complete after deformation. As one can see in Fig. 3e, an embedded picture of Fig. 3d, the thin b lath was broken into pieces and the adjacent two a lamellar connected together, which was associated with a series of globularization. The b transformed microstructure was sporadic distributed at certain regions. The fractured (TiB þ TiC þ Y2O3) reinforcements were also observed in Fig. 3d that was a result of local stress.

3.2. The dDRX of the b deformation 3.2.1. PSN by (TiB þ TiC þ Y2O3) reinforcements Preliminary statistical results indicated that the average dDRX

grain size of composite was significant finer than that of matrix. What’s more, the volume fraction of dDRX region of composite is 50% that is far more than that in matrix (Table 1). In other words, the second phase such as (TiB þ TiC þ Y2O3) reinforcement can significantly accelerate DRX process. Qiu [7] and Qu [25] both reported the similar phenomena in TMCs. According to dynamic recrystallization theory [23], the formation of dDRX b grains was associated with nucleation and growth process. The TiB and TiC particle could cause dislocation pile-up and promote the nucleation of equiaxed grain, thus made the microstructure nearby the reinforcement much finer. Huang [9] systematically summarized this mechanism which referred as particle stimulated nucleation (PSN). Ma [26] and Chen [27] confirmed this in their latest research. The dispersed TiB near the grain boundary can retard or pressure on the moving (sub)grain boundaries by Zener pinning force [28]. It is widely accepted that most of second phase particles can accelerate both the nucleation and growth stage of recrystallization by Zener pinning effect [9]. Therefore, with the help of (TiB þ TiC þ Y2O3) reinforcements, the composite in present study showed an excellent DRX behavior that was effective to refine and homogenize microstructure. 3.2.2. Variant selection of a in (TiB þ TiC þ Y2O3)/a-Ti composite Fig. 4 shows the EBSD maps of matrix and (TiB þ TiC þ Y2O3)/aTi composite subjected to b deformation. The high quality maps of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite revealed that the EBSD results was very credible (Fig. 4 a and c). Specially, the TiB and

Table 1 Average dDRX grain size and fraction of dDRX of the matrix and (TiB þ TiC þ Y2O3)/ a-Ti composite subjected to b deformation. Martials

Average dDRX grain/mm

Fraction of dDRX region/%

Matrix (TiB þ TiC þ Y2O3)/a-Ti

46.0 17.0

18.8 50.0

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Fig. 4. EBSD image quality maps of matrix (a) and (TiB þ TiC þ Y2O3)/a-Ti composite (c) subjected to b deformation; IPF and grain boundary map of matrix (b) and (TiB þ TiC þ Y2O3)/a-Ti composite (d) subjected to b deformation; (e) and (f) is the {0001}a pole figure of the (b) and (d), respectively; (g) crystallographic directions and corresponding reference colors in the EBSD IPF maps in (b) and (d); (h) pole plot of {0001}a in (e) and (f). Note that the low angle boundary (LABs) is pointed out by thin white line while high angle boundary (HABs) is pointed out by crude black line in (b) and (d). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

TiC particle were marked with red and green regions, respectively. Fig. 4 b and d were inverse pole figure and grain boundary map (IPF þ GB map). Fig. 4 g marked the sample orientation and crystallographic orientation of a-Ti, TiB and TiC, respectively. Due to a poor resolution, the effect of Y2O3 was not included in this study. Obviously, the orientation of a- Ti in dDRX region was diversity (blue dotted square in Fig. 4a and c) than prior coarse a colony. The prior coarse a colony usually had one orientation while the newly formed dDRX grain not only had fine grain size but also an abundant orientation. Things got more apparent in composite that even in one dDRX grain the a orientation was more colorful that revealed a much rich variant selection of a phase. The {0001}a pole figure further demonstrated this result. Fig. 4e and f depicted the {0001}a pole figure of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite. Furthermore, the distribution of above {0001}a in matrix and composite was showed in Fig. 4h. The distribution of {0001}a pole figure in composite was more average and dispersed that was a results of more a variant selection. Anyway, the c-axis of a lamellar was almost deviated from 90 to Z axis that meaned the c-axis of most of a laths parallel to metal flow direction. This was similar with Germain [29]and Obasi [30]. It was worthy to note that b transformed microstructure of the region without undergoing DRX usually owned only one orientation that inherited form prior a orientation during a/b/a phase transformation [31]. Thus a considerable region had the same a orientation such like both sides of dDRX region in Fig. 4 a which in turn deteriorated final mechanical properties. In titanium alloy, the Widmanstatten a colony (aWGB) always grows with parent b grain with Burgers orientation relationship (BOR, i.e.{0001}a//{110}b and <11e20>a//<111>b) which is associated with lowest interface energy [32]. Fig. 5 coincidentally showed that both two adjacent a lamellar were bound to BOR to parent b grain in matrix. The high resolution TEM maps and Fourier

transform maps (FFT maps) were also further identified the certain crystallographic orientation relationship. What’s more, the misorientation between a1 and a2 is 60 by grid method. It is feasible that this two a grains are bound to BOR with the same b grain at the circumstance where the two a grains parallel to different habit plane such as (110)b and (011)b which misorientation is exactly 60 . However, in (TiB þ TiC þ Y2O3)/a-Ti composite, the as direct precipitates in b grain without forming aGB at b grain boundary because of the constitutional undercooling and heterogeneous nucleation caused by (TiB þ TiC þ Y2O3), thus promotes the formation of non-Burgers related a grain which weaken a texture at last [22,33]. Fig. 6 provided insight into the local crystallographic relationship between a and b phase. There must be a statement that only a lamellar orientation can be recorded clearly while the b laths is not so lucky to be identified. This is because the b phase is thin and easy to be eroded in electrolytic polishing. But the experiment results still can qualitative reflect an interesting results. As one can see, in matrix, the transformed a colonies were on the whole bounded to BOR in newly DRX b grains; while in composite, there was almost none BOR boundary. On the one hand, the only way to minimize interface energy during b/a phase transformation in matrix was to grow with BOR boundary. On the other hand, the reinforcement such as TiB whiskers, could provide additional energy to compensate the additional interface energy for nucleation and the reinforcement also could supply additional site for heterogeneous nucleation of a phase in which results the formation of non-Burgers related a grain within b grain. Therefore, the {0001}a pole figure of (TiB þ TiC þ Y2O3)/a-Ti composite depicted more diverse and random compared to matrix, which in turn weaken the texture of composite. Such characteristic made it have a great advantage in preventing Macon-zone that referred as large a regions with strong local textures.

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Fig. 5. (a) Bright-field STEM image of a precipitates subjected to b deformation in matrix; (b) high resolve STEM of region A in (a); (c), (d), (e), (f), (g) and (h) are FFT of certain region of B, D, C, F, G and H. region C and G corresponding to a1 and a2 grain, respectively; region D corresponding to b grain; region B and H corresponding to (a1þb) and (a2þb), respectively, both a1 and a2 grain are maintain BOR with b grain; region F corresponding to (a1þa2), By grid method, we can speculate that the misorientation between a1 and a2 is 60 .

Fig. 6. EBSD image quality maps with highlighted boundaries linked to the Burgers orientation relationship in matrix (a) and (TiB þ TiC þ Y2O3)/a-Ti composite (b), respectively. Note that the red line corresponding to strictly maintain BOR, the yellow line corresponding to deviated BOR less than 10 and the white line corresponding to deviated BOR more than 10 . (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

3.3. Kinked a colonies of the (aþb) deformation In general, DP microstructure, which is known for its excellent performance in balancing between strength and ductility, involving series of thermo-mechanical processing steps at temperatures below the b-transus. It was reported that such microstructure was desirable for superplasticity which has a considerable potential in

manufacturing processes [34,35]. But larger deformation resistance is usually accompanied by deformation in (aþb) field, one of typical characteristic is kinked a colonies. As mentioned in 3.1, the kinked a colonies is seen in everywhere that is a directly evidence to high local shear stress. Here are emphasis on kinked a colonies and its deformation behavior.

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3.3.1. Grain boundaries in (TiB þ TiC þ Y2O3)/a-Ti composite Fig. 7 shows the representative EBSD maps of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation. As one can see, the high quality maps of matrix and composite clearly restored the original microstructure that revealed the EBSD results were reliable (Fig. 7 a and b). The colonies morphology of matrix was overall unbroken except certain region where the orientation of a parallel to metal flow direction (yellow circle in Fig. 7 a). Those a lamellar was undergone a series globalization thus broken down to equiaxed a grains. However, the microstructure of composites showed somehow confused that the unconspicuous colony morphology was hardly to recognize. As mentioned in 3.1, the boundary of a and b lamellar was not straight anymore that revealed a lots of substructures were formed. As we all know, a laths have smooth and flat interfaces, most of boundaries were high angle in deformation condition; only sporadic low angle boundaries (LABs; i.e.15 ) were found within some a colonies [36,37]. However, things were absolutely different which could be demonstrated by IPF þ GB maps (Fig. 7 c and d). In those maps, LABs and high angle boundaries (HABs; i.e.>15 ) are depicted as white thin lines and thick black lines, respectively. The misorientation distribution of a grains were also illustrated in Fig. 7 e and f. In matrix, the LABs were observed in the region where a colony most likely kinked that reveals that certain region underwent a shear strain. Moreover, the LABs were also distributed in inter-lamellar that revealed there was an uncoordinated deformation between a lamellar and a lamellar in certain a colony. But the HABs were mainly distributed in grain boundaries and a colonies boundaries where the deformation was severe and heterogeneous. In general, the high energy defects such as dislocations are most likely to aggregate at grain boundaries and colony boundaries0 thus forms HABs which is a typical process of globularization. When it comes to (TiB þ TiC þ Y2O3)/a-Ti composite, the

circumstances becomes more severe. Compared to matrix, the grain boundaries were extremely abundant. As introduced above, the colony morphology was no longer obvious instead the typical near basketweave microstructure. According to Luo [38], the interleaved a lamellae could restrain the rotation of neighboring a lamellae and promote the formation of LABs. Fig. 7 e and f shows the misorientation distribution of a grains of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite. The percentage of LABs of composite was 76.2%, which was 1.3 times of that in matrix (58.4%). On the one hand, the interleaved a lamellae restrained the deformation of neighboring a lamellae and promoted the accumulation of dislocations in interleaved a lamellae junction. On the other hand, (TiB þ TiC þ Y2O3) reinforcements themselves were also responsible for the high density dislocations due to their inhibition to deformation. As one can see, LABs were observed at both the inter-lamellar and intra-lamellar. And the thermal groove and subsequent formation of HABs were evolved from those LABs. Those above processes are essential to globularization. The globularization will be detail discuss in 3.4. 3.3.2. Misorientation profile along and transverse the a lamellar Apparently, the IPF maps of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite (Fig. 7c and d) indicated that there was a considerable orientation gradient within certain a lamellar and a colony. Specially, Fig. 8 and Fig. 9 gives the misorientation profiles of typical line in Fig. 7: i.e. along the a lamellar (A1 and A2) and transverse to the a lamellar (T1 and T2). The point-to-point misorientation profiles along a lamellar of matrix (A1, Fig. 8 a) was smooth within a change in 3 . Actually, most of point-to-point changed less than 2 , which was referred as the criterion to an undeformed material within the angular resolution of EBSD [37]. Ridiculously, the cumulative misorientation profile of A1 was up to the maximum of 19 that was even higher than usually defined LABs. The cumulative

Fig. 7. EBSD image quality maps (a and b), IPF þ GB maps (c and d) and a lamellar misorientation distribution maps (e and f) of matrix (a, c and e) and (TiB þ TiC þ Y2O3)/a-Ti composite (b, d and f) subjected to (aþb) deformation. Note that the low angle boundary (LABs) is pointed out by thin white line while high angle boundary (HABs) is pointed out by crude black line in (b) and (d).

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Fig. 8. Misorientation profiles obtained from typical line in Fig. 7 (A1 and A2) along the a lamellar of matrix (a, c) and (TiB þ TiC þ Y2O3)/a-Ti composite (b, d): (a) point-to-point misorientation profiles of A1 of matrix; (b) point-to-point misorientation profiles of A2 of (TiB þ TiC þ Y2O3)/a-Ti composite; (c) cumulative misorientation profile of A1 of matrix; (d) cumulative misorientation profile of A2 of (TiB þ TiC þ Y2O3)/a-Ti composite.

misorientation appeared to increase with the deformation reduction. Mironov [37] found this value increased from 8 with 25% reduction to 28 with 70% reduction. Besides, the temperature was also another key factor to influence deformation behavior. In present study, the alloy was deformation with a reduction of 85% at 950  C, it is reasonable that the cumulative misorientation raise to such value. However, the point-to-point misorientation profiles along a lamellar of composite (A2, Fig. 8 b) was sharp that the maximum angle could reach to 7, which was far more than that in matrix. Conservatively speaking, the normal float was around 4 in Fig. 8 b. According to Romanov [39], the internal stress s is proportional to the Gu, where G is the shear modulus and u is the cumulative rotation angle (Frank vector). Consequently, the internal stress of the composite is much higher than that in matrix. Therefore, it was unsurprising that the maximum cumulative misorientation angle was almost to 40 . In other words, such high misorientation could be considered as two different a grains. As for transverse to a lamellar (T1 and T2) aspect, the point-topoint misorientation profiles in both matrix and (TiB þ TiC þ Y2O3)/ a-Ti composite changed sharp. And the maximum point-to-point angle of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite could reach to 40 and 90 , respectively (Fig. 9 a and b). In general, the orientation gradient transverse to the a lamellar direction was much high than that of along the a lamellar direction that indicated that the a/b phase boundary had a significant role in high temperature deformation. Compared to matrix, the orientation gradient changed more frequent that almost every adjacent lamellar had distinct large misorientation. As a result, the cumulative misorientation in matrix and composite was high and

changeable that revealed a violent deformation had happened in a/ b phase boundary thus resulted a series of HABs. Especially in (TiB þ TiC þ Y2O3)/a-Ti composite, the deformation of interlamellar was serious than intra-lamellar. Most of HABs were observed in Fig. 7 d that demonstrated a common dynamic globularization in intra-lamellar during deformation. 3.4. Globularization during (aþb) deformation 3.4.1. Globularization depend on orientation during (aþb) deformation In general, these high angle cumulative misorientation of both along and transverse the a lamellar was mainly contributed by the crystallographic rotation of neighboring a lamellar which nature associated with the slip of a lamellar and a colony [37]. Due to nearly equivalent critical resolved shear stress (CRSS), the major slip type of a titanium at high temperature (700  C) is both basal slip system ({0001} <11e20>), prism slip system ({10-10} <11e20>) and pyramidal slip system ({10e11} <11e20>) [3,40]. Systematically, Roy [41,42] classified five type of a orientation (include//TD,//ND,//RD, aligned within 45 to the straining directions (ND/RD) and aligned within 45 to TD), and found that for the first three type of a lamellar no basal slip but only limited prism slip; while for last two type of a lamellar both basal slip and prism slip could be activated. Therefore, those a lamellar whose c-axis parallel to RD would become kinked while others a lamellar whose c-axis parallel to ND or TD would become elongated morphology. This was consistent with further finding that those a parallel to compression axis would conduct bent a colonies while those a perpendicular to compression axis

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Fig. 9. Misorientation profiles obtained from typical line in Fig. 7 (T1 and T2) transverse to the a lamellar of matrix (a, c) and (TiB þ TiC þ Y2O3)/a-Ti composite (b, d): (a) point-topoint misorientation profiles of T1 of matrix; (b) point-to-point misorientation profiles of T2 of (TiB þ TiC þ Y2O3)/a-Ti composite; (c) cumulative misorientation profile of T1 of matrix; (d) cumulative misorientation profile of T2 of (TiB þ TiC þ Y2O3)/a-Ti composite.

(parallel to metal flow direction) would conduct elongated a colonies. Schmid factor (SF), as a common parameters, is a key factor to judge the slip system activity and how easy it is to slip. Fig. 10 shows the SF value of three typical slip system of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation. Intuitively, the distribution of SF value was distinct between matrix and (TiB þ TiC þ Y2O3)/a-Ti composite. The SF value of matrix was varies with the a colony, but it was more homogeneous in composite. In matrix, as for certain a colony, it almost had the same SF value for certain slip system. For example, the a colony located in the upper left of Fig. 10 a who had a lower basal slip SF value, which was kinked ultimately after deformation. but the a colony located in the upper right of Fig. 10 a who had a highest SF value, which was easy to operate slip that retained complete colony morphology. Not surprisingly, the SF value of prism/pyramidal slip system was indeed higher than that of basal slip system which was consistent with Roy [41]. Therefore, the prism/pyramidal slip system were prone to operate compared to basal slip at upper (aþb) deformation. However, the distribution SF value was homogeneous and random in (TiB þ TiC þ Y2O3)/a-Ti composite that revealed a random orientation of a laths. In other words, the random orientation of a lamellar that signified a different SF value between adjacent a lamellar which at last leaded to incompatibility slip. Thus further developed to pill-up of dislocation at inter-lamellar which was a significant process of globularization. This phenomenon could be further demonstrated by a large amount of HABs between inter-lamellar in Fig. 7 d and sharp changeable point-topoint misorientation of transverse to the a lamellar in Fig. 9 b.

Hence, there was a priority in composite to globularization. Besides, the globularization of a grains also gave a priority to kinked a colony even though it has a low SF value. As we all known, some certain a colony with hard orientation (low SF value) would yield and kinked during deformation. The large crystallographic rotation in those kinked a colonies was usually associated with a large local shear stress and high density dislocations. Ito [36] recently pointed out that the heterogeneous of slip in a lamellar thus generated two situations at last: discontinuous changes in orientation, continuous and smooth changes in orientation. For those discontinuous change in orientation point, the rotate angle increases with increasing strain and results in thermal groove formation on the lamellar interfaces by interconnecting boundaries within lamellae, which leads to a subdivision of a grains. While for those continuous change in orientation, whether further deformation or subsequent annealing that to accumulate dislocations to reach to geometrically necessary dislocations (GNDs) thus to form geometrically necessary boundaries (GNBs) [43]. As depicted in Fig. 8, the higher misorientation the a lamellar was, the larger local stress and crystallographic rotation it obtained. Thus the boundary splitting was happened at the point where substructure density was high. Moreover, it was a process controlled by the density of high defects. As a result, the kinked a colonies usually first accomplished globularization. Furthermore, there was also another globularization mechanism i.e. continuous dynamic recrystallization (cDRX) of a grain. It was common in lamellar a at relatively low temperature that associated with greater torsion of a lamellar which resulted in HABs at last [38,44,45]. Gao [46] and Zhou [47] found that the process of “tangled cell structure / subgrain bounded by LABs / HABs”

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Fig. 10. Schmid factor maps of different slip system of matrix (a, b, c) and (TiB þ TiC þ Y2O3)/a-Ti composite (d, e and f) subjected to (aþb) deformed; (a) and (d) SF of basal slip system; (b) and (e) SF of prism slip system; (c) and (f) SF of pyramidal slip system. Note that the color corresponding to actual Schmid factor value as shown in bowl of the picture. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

leaded to fragmentation of a grains during deformation. In the present study, lots of HABs (equiaxed a grains) were observed in Fig. 7 c and d, their orientation were rich and colorful that was associated with cDRX thus accelerated the fragments of a grains. 3.4.2. Globularization promoted by (TiB þ TiC þ Y2O3) reinforcements Reinforced with (TiB þ TiC þ Y2O3) reinforcements, the composite displayed different deformation behavior. Fig. 11 a and b shows the local dislocation distribution of matrix and (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation.

Similar with SF value, the matrix had a heterogeneous local dislocation which revealed a huge stress concentration at certain region. The deeper red color indicates greater local stress. The highest local stress region was located at grain boundaries (b grain boundaries and a colony boundaries) and kinked a colonies. However, the local dislocation distribution of (TiB þ TiC þ Y2O3)/a-Ti composite was homogeneous and had a relatively lower local stress. According to our usual understanding, the composite reinforced with those (TiB þ TiC þ Y2O3) usually conducted a high local stress. Ridiculously, the result challenged our cognize. Roy [15,48] pointed out that the biggest difference between the matrix and (TiB þ TiC þ Y2O3)/a-Ti

Fig. 11. Local dislocation distribution map of matrix (a) and (TiB þ TiC þ Y2O3)/a-Ti composite (b) subjected to (aþb) deformation; (c) phase distribution and grain boundary map of (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation.

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Fig. 12. (a) Bright-field STEM image of morphology of (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation; (b) selected area diffraction (SAD) of TiB whisker in the yellow circle of (a); (c) pill-up dislocations in interleaved a lamellar; (d) substructure in a lamellar near the TiB whisker; (d) pill-up dislocations in certain point of a lamellar where is prone to thermal groove. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

composite was that the volume fraction of grain boundary in composite was higher than that of matrix owing to the refinement of reinforcements. Thus the adequate grain boundary slip was happened in composite. But the deformation behavior of matrix was still controlled by substructure process that was associated with dislocation motion. Therefore, the local dislocation of matrix was high and mainly distributed at grain boundary while the local stress of composite was relatively lower and homogeneous due to grain boundary slip. Fig. 11c shows the phase distribution of (TiB þ TiC þ Y2O3)/a-Ti composite subjected to (aþb) deformation. Specially, the TiB and TiC particle were marked with red and green regions, respectively. Coincidently, the a grains near the reinforcement was indeed the region where the a grains suffered large local shear stress. As a result, a large density of dislocations were piled up near the reinforcements which was the necessary requirement for globularization. Fig. 12 shows the TEM morphology of (TiB þ TiC þ Y2O3)/aTi composite. Fig. 12a depicted typical microstructure of (TiB þ TiC þ Y2O3)/a-Ti composite. The interleaved a lamellar, equiaxed a grains and TiB particle were observed. And the selected area diffraction (SAD) of TiB particle in Fig. 12a (yellow circle) was showed in Fig. 12b. In addition, the typical of pill-up dislocation of a lamellar near TiB was observed in Fig. 12d there were lots of substructures that divided this certain a lamellar into several parts. Some substructure even was formed with dislocation wall (Fig. 12e). Besides, the pill-up dislocations was also be found in interleaved a lamellar (Fig. 12c) that was similar with Luo [38]. In general, the thermal groove controlled by boundary splitting mechanism was very active at such point where high density of

defects was. And the solute migration rate was highly depended on the density of dislocations. On the one hand, the random orientation of adjacent a lamellar in (TiB þ TiC þ Y2O3)/a-Ti composite accelerated their uncoordinated deformation thus leaded to high density dislocation accumulation; on the other hand, the (TiB þ TiC þ Y2O3) itself could obstruct dislocation that caused pillup dislocation in near a lamellar. Hence, the (TiB þ TiC þ Y2O3)/a-Ti composite had an apparent dynamic globularization behavior compared to matrix. There must be stated that the globularization of a grain would completely finish after annealing at deformation temperature for several hour. Furthermore, the globularization kinetics of composite is also faster compared to matrix under identical annealing conditions [49]. Apart from that, the reinforcement also had a series of influence on deformation behavior in the composite. Gaisin [50] emphasized that the boron-modified titanium had a better recrystallization behavior during hot deformation; Roy [15] pointed out that the softening temperature of titanium alloy could decreased by ~50  C with trace boron addition; and Srinivasan [51] also confirmed that the titanium alloy with litter boron addition could improves the processability of the cast ingots and enables direct rolling of the cast ingots. Therefore, the (TiB þ TiC þ Y2O3)/a-Ti composite had a great advantage in optimizing microstructure especially globularization that made the composite of great application prospect. 4. Conclusions In present study, the deformation behavior and deformation mechanisms were investigated in (TiB þ TiC þ Y2O3)/a-Ti

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composite by EBSD and TEM analysis to evaluate the effect of (TiB þ TiC þ Y2O3) reinforcements on deformation behavior. The main results were as follow: (1) During b deformation, the dDRX grain was observed mainly at prior b grain boundaries; and the (TiB þ TiC þ Y2O3) reinforcements could accelerate this process by PSN mechanism. (2) The reinforcement promoted the heterogeneous nucleation of a phase thus conducted the {0001}a pole figure of composite after b deformation more widely distributed compared to matrix. (3) During (aþb) deformation, the microstructure was characteristic of kinked a colonies morphology. Moreover, compared to matrix, the (TiB þ TiC þ Y2O3)/a-Ti composite depicted a more homogeneous microstructure and larger crystallography rotation both along or transverse to the a lamellar. (4) The SF value of matrix during (aþb) deformation was varies with the a colony that revealed the deformation behavior was heterogeneous, but it was more homogeneous in composite. (5) Reinforced with (TiB þ TiC þ Y2O3) reinforcements, on the one hand, the random SF value of composite conducted the incompatibility slip between adjacent a lamellar thus at last leaded the formation of HABs in inter-lamellar; on the other hand, the reinforcement itself could obstruct dislocation that accelerated the dynamic globularization behavior. (6) The (TiB þ TiC þ Y2O3)/a-Ti composite had a great advantage in optimizing microstructure especially globularization that made the composite of great application prospect.

Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Jianhui Yang: Investigation, Writing - original draft, Writing review & editing. Shulong Xiao: Conceptualization, Methodology. Yuyong Chen: Supervision, Funding acquisition. Lijuan Xu: Validation, Visualization. Xiaopeng Wang: Formal analysis. Jing Tian: Resources. Dongdong Zhang: Data curation. Zhuangzhuang Zheng: Data curation. Acknowledgements The authors thank the financial support provided by the National Key R&D Project of China (2017YFE0123500) and Sichuan science and technology program (2019YFG0082). References [1] I. Weiss, S.L. Semiatin, Thermomechanical processing of alpha titanium alloysdan overview, Mater. Sci. Eng., A 263 (2) (1999) 243e256. [2] D. Banerjee, J.C. Williams, Perspectives on titanium science and technology, Acta Mater. 61 (3) (2013) 844e879. [3] G. Lütjering, J.C. Williams, Titanium[M], Springer Berlin Heidelberg, New York, 2007. [4] J.P. Qu, C.J. Zhang, J.C. Han, et al., Microstructural evolution and mechanical properties of near a-Ti matrix composites reinforced by hybrid (TiBþY2O3) with bimodal size, Vacuum 144 (2017) 203e206. [5] Y. Jiao, L. Huang, L. Geng, Progress on discontinuously reinforced titanium matrix composites, J. Alloys Compd. 767 (2018) 1196e1215.

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