Microstructure evolution of the transitional region in isothermal local loading of TA15 titanium alloy

Materials Science and Engineering A 528 (2011) 2694–2703

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Microstructure evolution of the transitional region in isothermal local loading of TA15 titanium alloy X.G. Fan, P.F. Gao, H. Yang ∗ State Key Laboratory of Solidification Processing, School of Materials Science and Engineering, Northwestern Polytechnical University, P.O. Box 542, Xi’an 710072, PR China

a r t i c l e

i n f o

Article history: Received 21 September 2010 Received in revised form 1 December 2010 Accepted 2 December 2010 Available online 9 December 2010 Keywords: Local loading TA15 titanium alloy Transitional region Microstructure evolution

a b s t r a c t The microstructure of the transitional region determines the performance of TA15 component under isothermal local loading forming. To understand the characteristic of microstructure evolution in transitional region, a set of analogue experiments for isothermal local loading were designed and carried out to investigate the microstructure development under different temperatures and complex strain path. It is found that the isothermal local loading does not change the microstructural composition of TA15 alloy with initially equiaxed microstructure in the transitional region, though a small fraction of lamellar ␣ phases appear at temperatures above 950 ◦ C due to reheating and small deformation in the second loading step. Recrystallization takes place in the ␤ matrix but not in the primary equiaxed ␣ phases. The ␤ grains are fine and equiaxed irrespective of strain path due to recrystallization. The existence of lamellar ␣ refines the ␤ grains. Kinked and disordered primary equiaxed ␣ phases are produced under complex strain path. The grain size of primary equiaxed ␣ phases increases slightly due to static coarsening and strain induced grain growth in multi-heat processing. The volume fraction of primary equiaxed ␣ phases is not influenced by strain path but determined by working temperature. © 2010 Elsevier B.V. All rights reserved.

1. Introduction The large-scale integral complex components of titanium alloys (such as bulkhead) which can meet the demand of high performance, light weight and high reliability, have been gaining extensive applications in the aerospace industry. However, it is difficult to form such components due to the complexity in shape, large deformation resistance and low ductility of the material, and high requirement of forming quality. The local loading forming provides a feasible way to form such components. During local loading forming, load is applied to part of the billet and the component is formed by changing loading region. By controlling material flow and reducing loading area at each forming step, the local loading forming can enhance the plasticity and filling ability of the material, and enlarge the size of component to be formed. An example of local loading forming is given in Fig. 1 [1], where the large-scale integral bulkhead was formed by one pass local loading (including two loading steps). The local loading forming is a multi-step hot working process. The alternation of loading region results in the loading region, the unloading region, and the transitional region which connects the loading region and the unloading region (Fig. 2). The transitional region deforms under

∗ Corresponding author. Tel.: +86 029 8849 5632; fax: +86 029 8849 5632. E-mail addresses: [email protected], [email protected] (H. Yang). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.12.006

the constraints of the loading region and the unloading region, and undergoes large uneven plastic deformation and complex thermal processing path. Therefore, the microstructure evolution might be very complex in the transitional region. Consequently, the transitional region is more concerned as it may determine the performance of the component. Extensive investigations have been carried out on the macroscopic deformation [2–9] and microstructural changes [1,10–12] during local loading forming. In Refs. [1,10], the effects of deformation inhomogeneity and processing parameters on the microstructure and mechanical properties of TA15 titanium alloy under local loading were experimentally investigated. However, the effect of strain path variation inside the transitional region on the microstructure needs further investigation. Fan [11] and Li [12] studied the variation of primary equiaxed ␣ grain size in local loading forming using FE simulation. However, the performance of titanium alloys is influenced by the size and morphology of every phase. Therefore, investigations are furthermore needed so as to reveal microstructure changes in the transitional region. In current study, an analogue experiment which can reflect the deformation characteristic of the transitional region was designed and carried out. A FE model of the analogue experiment was established for quantitative determination of deformation behavior in the transitional region. Microstructure evolution under different strain paths and hot working conditions was observed and analyzed. It will provide the basis for microstructure control and

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Fig. 1. A large-scale titanium bulkhead by one pass local loading forming: (a) billet; (b) after the first loading step; (c) after the second loading step [1].

consisted of approximately 60% primary ␣ phases (with a standard deviation of 6%) within the transformed ␤ matrix. The average grain size and aspect ratio of the primary ␣ were approximately 7.6 ␮m and 2.4, respectively. No obvious ␤ grain boundary was observed in the transformed ␤ matrix and the lamellar ␣ was about 1.4 ␮m thick.

Fig. 2. Schematic diagram of the transitional region [1].

performance improvement in local loading forming of titanium alloys. 2. Materials and methods 2.1. Starting material A near-␣ TA15 titanium alloy was used in current study with the chemical composition (wt.%) of 6.06 Al, 2.08 Mo, 1.32 V, 1.86 Zr, 0.30 Fe and balanced Ti, and measured ␤-transus temperature of 990 ◦ C. The TA15 alloy was received in hot rolled plate form with specification 1200 mm × 170 mm × 80 mm and initial microstructure given in Fig. 3. The microstructure of the as-received material

Fig. 3. Microstructure of the as-received TA15 alloy.

Fig. 4. Schematic of the analogue experiment (one pass, two steps): (a) the first loading step; (b) the second loading step; (c) the heating and deformation schedule.

Fig. 5. Photos of the workpiece exploited in the local loading test.

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Fig. 6. FE plot of strain distribution on the cross section in local loading forming to 50% reduction at 950 ◦ C and nominal rate of 0.01 s−1 : (a) after the first loading step; (b) after the second loading step. P1–P4 and p1–p4 are the image locations for microstructure observation.

2.2. Experimental procedure As mentioned in Section 1, the deformation of the transitional region is not finished in a single loading step but accumulated in the

Fig. 7. Strain paths at the image locations.

multi-step processing. In each forming step, the transitional region connects the loading region and unloading region and it undergoes large uneven deformation and complex stress state due to the constraints of loading region and unloading region. The amount of

Fig. 8. Microstructure prior to local loading at (a) 910 ◦ C; (b) 930 ◦ C; (c) 950 ◦ C; (d) 970 ◦ C.

Fig. 9. Microstructure parameters prior to deformation at different temperatures: (a) volume fraction of ␣p ; (b) aspect ratio of ␣p ; (c) grain size of ␣p ; (d) ␤ grain size.

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Fig. 10. Microstructure at different working temperatures after the first local loading step: (a)–(d) 910 ◦ C; (e)–(h) 930 ◦ C; (i)–(l) 950 ◦ C; (m)–(p) 970 ◦ C. At each testing temperature, the images are taken at p1–p4 marked in Fig. 6(a), respectively. The compression axis is vertical.

deformation decreases away from loading region in each loading step. Therefore, the strain path varies with the specific location in transitional region. To simulate the multi-step uneven deformation behavior in transitional region with reduced cost, an analogue experiment was designed in current study, as shown in Fig. 4. 20 mm × 16 mm × 12 mm cuboid specimens were machined from the as-received plate with the height direction parallel to the normal direction of the plate (Fig. 5(a)). There exist two loading steps in one pass of the local loading process. In the first loading step, the specimen was heated to the test temperature at 12 ◦ C/min, held for 15 min to achieve thermal equilibrium prior to testing, compressed along the height direction with a cylindrical anvil and a flat anvil (Fig. 4(a) and (c)) at a constant nominal strain rate of 0.01 s−1 , and then air cooled. In the second loading step, the specimen was reheated to the same test temperature, pressed between two flat anvils to the same reduction rate, and air cooled (Fig. 4(b) and (c)). The processing parameters are shown in Table 1. All the experiments were carried out in the (␣ + ␤) phase field. After the local loading process, the specimen was annealed with the following heat treatment route: 810 ◦ C/1 h/AC. For each test temperature, one specimen was processed only with the first loading step to capture

the microstructure changes throughout processing. The deformed samples are shown in Fig. 5. 2.3. FE modeling The local loading process was simulated using a commercial FE code DEFORM3D to obtain the deformation behavior of the workpiece. Fig. 6 illustrates the strain distribution on the cross section of the workpiece. The strain decreased sharply from middle to both sides of the workpiece after the first loading step (Fig. 6(a)), and a more uniform strain distribution was achieved after the second loading step (Fig. 6(b)). In Fig. 6(b), P1–P4 mark the image locations for microstructure observation after the second loading step. The corresponding locations after the first loading step are marked as

Table 1 Processing parameters. Temperature ( ◦ C) Reduction rate (%) Strain rate (s−1 ) Loading pass

910, 930, 950, 970 50 0.01 1

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Fig. 11. Microstructure parameters after the first loading step: (a) volume fraction of ␣p ; (b) aspect ratio of ␣p ; (c) grain size of ␣p ; (d) ␤ grain size.

p1–p4 in Fig. 6(a) by backward tracing. The strain paths at these locations are shown in Fig. 7. It is found that P1 underwent large deformation in the first loading step but small passive deformation in the second loading step, while P4 underwent small passive deformation in the first loading step but large deformation in the second loading step. P2 and P3 are deformed directly or indirectly in both loading steps. Therefore, the deformation feature of the transitional region can be reflected by the analogue experiment. 2.4. Metallographic evaluation Metallographic preparation of the specimens was carried out by sectioning in half, mechanically grinding and polishing, and etching with a solution of 13%HNO3 , 7%HF and 80%H2 O. Micrographs were taken on a LECIA DFC320 microscope, and examined using quantitative image analysis (Image-pro plus 5.0) to acquire the microstructural parameters. 3. Results and discussion 3.1. Microstructural features prior to local loading Fig. 8 illustrates the microstructure of the specimens which were heated to the test temperatures without deformation and air cooled. The measured microstructural parameters are given in Fig. 9. At the test temperatures, the starting microstructure consisted of primary equiaxed ␣ phases (␣p ) and ␤ matrix. In the lower (␣ + ␤) phase field (910 ◦ C), there existed some fine equiaxed or lamellar ␣ grains in the ␤ matrix (Fig. 8(a)), which confined the growth of ␤ grains. As a result, the ␤ grain size was small (Fig. 9(d)). With increasing temperature, the ␣p and fine ␣ grains transformed to ␤ phase simultaneously, and the fine ␣ grains gradually disappeared (Fig. 8(b) and (c)). Therefore, the ␤ grains grew larger. The ␣p grain size, however, decreased initially and then increased (Fig. 9(c)) which might be caused by the joint effect of ␣ → ␤ phase transformation and thermally assisted grain growth. In the upper (␣ + ␤) phase field (970 ◦ C), the percentage of ␣p decreased sharply

while the grain size increased slightly (Fig. 9(a)). The aspect ratio of ␣p varied little with temperature, which fluctuated slightly around the initial value (Fig. 9(b)). 3.2. Microstructure after the first loading step The microstructure after the first loading step at different test temperatures is given in Fig. 10. After the first local loading step, the microstructure mainly consisted of primary equiaxed ␣ phases and ␤ matrix throughout the specimen which was the same to the microstructure prior to deformation, suggesting that the uneven deformation did not affect the microstructural composition in one heat deformation under the conditions of current study. However, the morphology and size of different phases were changed after deformation. The morphology of ␣p was influenced by working temperature and deformation amount. In the lower (␣ + ␤) phase field, both ␣p and the fine ␣ phases were pancaked transverse to the compression axis at p1 where large compressive principle strain was imposed (Fig. 10(a)). At p2, ␣p which were originally perpendicular to the compression axis tilted due to the uneven deformation (Fig. 10(b)). ␣p were not only elongated but also distorted slightly, because the shear strain increased. As the compression strain decreased, ␣p were relatively undeformed at p3 and p4 (Fig. 10(c) and (d)). However, the distortion of ␣p was obvious due to the increase of shear strain. With increasing temperature, ␣p were less deformed. During subtransus working of wrought two-phase titanium alloys, strain partitioning between ␣p and ␤ matrix takes place as ␣p are harder than ␤ matrix. This phenomenon has been modeled for the twophase Ti–6Al–4V alloy [13] and near-␣ IMI834 alloy [14] using the self-consistent method. The self-consistent modeling suggested that viscosity-parameter ratio for ␣ and ␤ phases increases with temperature (the viscosity-parameter is proportional to the flow stress) and the strain rate ratio for ␣ phases and overall composite ε˙ ˛ /ε˙ decreases. Therefore, the deformation of ␣p decreased with temperature.

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Fig. 12. Microstructure at different working temperatures after the second local loading step: (a)–(d) 910 ◦ C; (e)–(h) 930 ◦ C; (i)–(l) 950 ◦ C; (m)–(p) 970 ◦ C. At each testing temperature, the images are taken at P1–P4 marked in Fig. 6(b), respectively. The compression axis is vertical.

Due to strain partitioning, the ␤ matrix underwent more deformation. However, the ␤ grains were finely equiaxed at p1 where large compression strain was imposed, indicating that the ␤ phase underwent dynamic recrystallization below the ␤ transus temperature. For many two-phase titanium alloys, reported apparent activation energy above ␤ transus temperature is close to the activation energy of self-diffusion in ␤ phase [14–16], indicating dynamic recovery dominates ␤ working [17]. Though dynamic recrystallization is often observed in ␤ working, the volume fraction is generally small [17–20]. However, ␤ recrystallization was greatly accelerated in (␣ + ␤) working. As a soft phase, the actual deformation was much larger than that of the aggregate. Moreover, the existence of ␣p which acted as hard inclusions during deformation, increased the deformation inhomogeneity of ␤ phase and provided more nuclei for recrystallization. As a result, ␤ recrystallization was fast and resulted in equiaxed ␤ grains (R grain), as shown in Fig. 10. Similar behavior has been reported for the near-␣ IMI834 alloy [19]. The morphology of the secondary ␣ phases (␣s ) was influenced by test temperature and uneven deformation. The ␣ precipitates were short and disordered at lower temperatures or large strains. The morphology of ␣s is determined by the nucleation rate of phase transformation. The nucleation cites are mainly at the grain bound-

aries or high energy crystal defects inside the grains [21,22]. A high defect density results in more nuclei for secondary ␣ precipitation, promotes the pipe-diffusion through dislocation core which enhances the growth and coarsening of the precipitates [21]. Consequently short and disordered secondary ␣ plates are produced. As the deformation amount varied on the cross section after the first loading step, the morphology of ␣s was different. At lower temperature, the strain rate imposed on the ␤ phases increased [13] while the thermal restoration was suppressed. Therefore, there were more crystal defects. Meanwhile, there existed more grain boundaries (␣–␤ and ␤–␤). As a result, relatively short and disordered secondary ␣ plates were observed. The measured microstructural indices are shown in Fig. 11. It is found that the volume fraction of ␣p varied little at a certain test temperature though large uneven deformation and the value was close to that prior to deformation, indicating that the deformation heating could be neglected at low strain rate (Fig. 11(a)). In the lower (␣ + ␤) phase field (910 ◦ C), the aspect ratio of ␣p was much larger at large deformations (p1 and p2) than at small deformations (p3 and p4) (Fig. 11(b)). The discrepancy decreased with temperature. Zhu et al. [23] reported the decrease of aspect ratio with the amount of deformation, and suggested that the pancaked ␣p can be broken up by processes such as boundary splitting

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Fig. 13. Microstructure prior to the second loading step at 950 ◦ C. The specimen was heated without deformation after the first loading step and water quenched. The image locations are (a) p1 and (b) p4.

and edge spheroidization, which account for the globularization of colony structure [24,25]. However, in current study the break-up of pancaked ␣p was not significant due to a relatively small deformation (a maximum reduction rate of 80% was used in [23]). Therefore, the variation of aspect ratio was determined by the deformation amount of ␣p , which increased with increasing deformation but decreasing temperature. At lower temperatures (≤950 ◦ C), the grain size of ␣p decreased slightly with the amount of deformation, while in the near-␤ phase field (970 ◦ C), the largest grain size appeared at p1 where maximum deformation occurred, as shown in Fig. 11(c). Dynamic recrystallization has been reported to take place in the ␣ phases of many titanium alloys [26,27], and it accounts for the microstructural changes of TA15 alloy in the temperature range much lower than current study [28,29]. However, dynamic recrystallization of ␣p was not observed in the present work. The variation of ␣p grain size might be caused by static grain growth, strain induced grain growth and strain induced grain refinement [30]. The strain induced grain growth plays an important role in low speed deformation [31].

At lower temperatures, the grain refinement dominated the grain size variation. Therefore, the grain size decreased with deformation amount. At higher temperature, the grain refinement was minor as the ␣p deformation was small. So the grain size increased with deformation amount due to strain induced growth. Quantitative metallography indicated that the ␤ grain size decreased with strain at all test temperatures (Fig. 11(d)), which could be attributed to the effect of ␤ dynamic recrystallization. 3.3. Microstructure after the second loading step Fig. 12 shows the microstructure after the second local loading step at different temperatures. The microstructure mainly consisted of ␣p within ␤ matrix, which was close to that after the first loading step (Fig. 10). The ␣p did not align transversely to the compression axis at P2–P4 because the strain states were complex at these locations after the second deformation. The elongated ␣p at P1 kinked due to a compression strain along the transverse direction. The transformed ␤ matrix was found to be smaller and

Fig. 14. Microstructure parameters after the second loading step: (a) volume fraction of ␣p ; (b) aspect ratio of ␣p ; (c) grain size of ␣p ; (d) ␤ grain size.

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Fig. 15. Microstructure at different working temperatures after annealing: (a)–(d) 910 ◦ C; (e)–(h) 930 ◦ C; (i)–(l) 950 ◦ C; (m)–(p) 970 ◦ C. At each testing temperature, the images are taken at P1–P4 marked in Fig. 6(b), respectively. The compression axis is vertical.

more disordered than that after the first local loading step. This is because the increase of specific surface area increased the cooling rate. At higher temperatures (≥950 ◦ C), a small fraction of lamellar ␣ phases were observed at P1 where the deformation was small in the second loading step (Fig. 12(i) and (m)). The lamellar ␣ gradually disappeared with increasing strain in the second loading step (Fig. 12(j)–(l) and (n)–(p)). During reheating, the secondary ␣ were not fully transformed to ␤ phase so the lamellar ␣ were produced. The lamellar ␣ were observed in the whole transitional region prior to the second deformation, as shown in Fig. 13. However, they might globularize with imposed strain or transform to ␤ phase due to deformation heating. The measured volume fraction of ␣p is illustrated in Fig. 14(a). It is found that the volume fraction of ␣p after the second loading step was not affected by strain path but determined by working temperature. It was close to that after the first loading step. Compared with the results after the first loading step, the aspect ratio of ␣p decreased at P1 and P2 but increased at P3 and P4 at temperatures below 950 ◦ C (Fig. 14(b)). The aspect ratio became larger at P3 and P4 than at P1 and P2. In the second loading step, ␣p would spheroidize during reheating, which decreased the aspect ratio. This process was accelerated by the stored energy of defor-

mation. The large deformation at P1 and P2 in the first loading step promoted the spheroidization. Meanwhile, the deformation in the second loading step was relatively small at P1 and P2. Therefore, the aspect ratio at P1 and P2 decreased. At P3 and P4, the large deformation in the second loading step increased the aspect ratio. In the near-␤ phase field (970 ◦ C), the aspect ratio was close at all image locations though different strain paths. It was smaller than that after the first loading step due to thermally assisted spheroidization. The ␣p grain size was a little larger at P1 than at other image locations at all test temperatures, as shown in Fig. 14(c). This may be caused by the grain growth which was promoted by the stored energy of previous deformation. Compared to the results after the first loading step, the average grain size increased slightly after the second loading step, indicating that static grain growth should be considered in multi-heat deformation. The ␤ grain size was smaller at P3 and P4 than at P1 and P2 at temperatures below 930 ◦ C. As discussed in section 3.2, this is due to the recrystallization of ␤ phase. The existence of lamellar ␣ at P1 above 950 ◦ C also refined ␤ grains. Therefore, the ␤ grain size was smaller though small deformation. The average ␤ grain size decreased slightly after the second loading step due to the recrystallization induced grain refinement.

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Fig. 16. Microstructure parameters after annealing: (a) volume fraction of ␣p ; (b) aspect ratio of ␣p ; (c) grain size of ␣p .

During hot working of single-phase alloys, the amount of deformation has a strong influence on the grain size. There exists a critical deformation range in which abnormal grain growth is prone to take place in the subsequent annealing. The transitional region connects the loading region and unloading region. Though the transitional region might partly fall into the critical deformation range in the first loading step and was reheated in the second loading step, coarse grains were not observed throughout the transitional region.

3.4. Microstructure after annealing The microstructure after annealing is illustrated in Fig. 15. The residual ␤ phases transformed to ␣ phases during annealing, so the lamellar ␣ were thickened. The lamellar ␣ were thicker, shorter and more disordered at locations where the deformation in the second loading step was larger (Fig. 15(d), (h), (l) and (p)). The volume fraction of the ␣p was not influenced by strain path at all test temperatures, as shown in Fig. 16(a). The volume fraction increased by about 6% after annealing. Similar behavior has been reported for TC11 alloy (Ti–6.5Al–3.5Mo–1.5Zr–0.3Si) in the double annealing treatment by Zhou et al. [32], which has been attributed to the mergence of secondary ␣ at grain boundaries of equiaxed ␣ phases. In most cases, the ␣p aspect ratio increased slightly at P1and P2 but decreased at P3 and P4 after annealing. However, the variation of ␣p aspect ratio with strain path showed no regularity, as shown in Fig. 16(b). The measured ␣p grain size is given in Fig. 16(c). It is found that the grain size decreased initially and then increased with increasing deformation in the second loading step. The grain size at all image locations increased after annealing due to the mergence of secondary ␣ into equiaxed ␣. The increasing extent was larger at P3 and P4 which underwent larger deformation in the second loading step. This is because the large stored energy promoted the mergence of secondary ␣. As a result, the grain size at P3 and P4 became larger though it was small prior to annealing.

3.5. Comparison between local loading forming and integral forming The transitional region undergoes multi-heat uneven deformation during isothermal local loading forming. In corresponding integral forming, simple compressive deformation takes place. The thermal processing path, strain state and stress state are more complicated in local loading forming. Though the complexity of local loading forming, the present study suggests that the local loading forming may not alter the microstructural composition of the TA15 alloy under current experimental conditions. The microstructure after local loading forming mainly consists of equiaxed ␣ and transformed ␤ matrix, though a small fraction of lamellar ␣ are produced due to reheating (Fig. 15(i)). The morphology of ␣p after local loading is different from that after integral forming. This is because the strain path is complex and diverse in local loading forming while a simple compressive strain is imposed in integral forming. The ␤ grains remain finely equiaxed due to recrystallization in hot working. The ␣p grain size increases in multi-heat processing. Previous investigations suggested that the coarsening of ␣p in duplex titanium alloy is controlled by bulk diffusion [33] and can be accelerated by low speed deformation [31]. However, the coarsening rate is low. As a result, the growth of ␣p is limited (Fig. 16(c)). The decrease in strength due to coarsening is small [1]. 4. Conclusions The microstructure evolution of the transitional region in subtransus isothermal local loading of TA15 titanium alloy has been investigated. From this work, the following conclusions were drawn: (1) In subtransus working of TA15 alloy with equiaxed microstructure, recrystallization takes place in the ␤ matrix but it is not observed in the primary equiaxed ␣ phases.

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(2) The complex strain path in local loading forming affects the morphology of the primary equiaxed ␣ phases, but not the morphology of ␤ grains. Kinked and disordered primary equiaxed ␣ phases are produced due to the variation of strain path while ␤ grains are equiaxed due to recrystallization. (3) The grain size of primary equiaxed ␣ phases increases slightly due to static coarsening and strain induced grain growth in multi-heat processing. The ␤ grains are fine due to recrystallization. (4) The volume fraction of primary equiaxed ␣ phases is not influenced by strain path at all test temperatures. The volume fraction increases by about 6% after annealing. (5) The isothermal local loading forming does not change the microstructural composition in the transitional region, though a small amount of lamellar ␣ appear in reheating at temperatures above 950 ◦ C. The lamellar ␣ disappear at large strain in the second loading step. Acknowledgements The authors would like to gratefully acknowledge the support of Natural Science Foundation for Key Program of China (No. 50935007), National Basic Research Program of China (No. 2010CB731701), Research Fund of the State Key Laboratory of Solidification Processing (NWPU) (No. 27-TZ-2010) and the 111 Project (B08040). References [1] X.G. Fan, H. Yang, Z.C. Sun, D.W. Zhang, Mater. Sci. Eng. A 527 (2010) 5391–5399. [2] K. Welschof, R. Kopp, Aluminium 63 (1987) 168–172. [3] R. Kopp, L. Schaeffer, G. Schuler, Metall. Plant Technol. Int. 5 (1982) 76–81.

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