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Microstructure and texture evolution of commercial pure titanium deformed at elevated temperatures
Materials Science and Engineering A 513–514 (2009) 83–90
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and texture evolution of commercial pure titanium deformed at elevated temperatures Zhipeng Zeng ∗ , Yanshu Zhang, Stefan Jonsson Department of Materials Science and Engineering, Royal Institute of Technology, SE-100 44 Stockholm, Sweden
a r t i c l e
i n f o
Article history: Received 17 September 2008 Received in revised form 13 January 2009 Accepted 19 January 2009 Keywords: Commercial pure titanium Texture Schmid factor Misorientation Hot compression Slip system
a b s t r a c t Microstructure and texture evolution of commercial pure titanium were investigated by electron backscattered diffraction (EBSD) after compression tests at elevated temperatures. The basal planes of both the fine and coarse grains in the deformed samples tend to rotate from the initial orientations, perpendicular to the compression axis, to an inclination of 45◦ . The Schmid factor is used to analyse how the individual slip systems activate and how their activities evolve under various deformation conditions. After deformation, the distribution frequency of the misorientation angles shows that the low angle grain boundaries increased dramatically while the high angle grain boundaries decreased. In particular, after deformation at 723 K and 0.1 s−1, a peak around 50–60◦ in the misorientation frequency-distribution is found, which is due to 1 0 1¯ 1 twinning. The analysis of the deformed microstructure indicates that dynamic recovery is the dominant deformation mechanism for commercial pure titanium when subjected to the investigated deformation conditions. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Mechanical properties of metals are dependent on the microstructure and texture formed during processing. In order to control the final properties of the products it is thus necessary to understand the microstructure and texture evolution during plastic deformation. Numerous efforts have been invested to investigate the microstructure and texture development in BCC and FCC metals. In contrast, there is little attention paid to the hexagonal close packed (HCP) metals. As a typical hexagonal close packed metal, commercial pure titanium (CP Ti) is of great importance in many industrial applications due to its highly attractive properties, such as good deformability at high temperature, low density, high biocompatibility and excellent corrosion resistance [1–2]. However, it is hard for CP Ti to strengthen to a level comparable to Ti–6Al–4V [3], which is currently the material of choice for most medical implants. Therefore, to develop higher strength CP Ti is an attractive work for medical applications. In order to increase the strength of CP Ti, large deformation with an accompanying grain refinement is an effective approach. In the present work, the uniaxial compression test was used because it is a convenient method for obtaining large strains and test specimens are easy to prepare. In addition, the deformation process is similar to forging; hence, it has become a commonly used approach
∗ Corresponding author. Tel.: +46 8790 6542; fax: +46 8203107. E-mail address: [email protected] (Z. Zeng). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.01.065
for testing mechanical properties during hot-working processes. However, most of the available investigations focus on severe deformation of CP Ti and its alloys [4–6]. Admittedly, there are some investigations on the deformation mechanisms of CP Ti at elevated temperatures based on the analysis of flow stress–strain curves or optical observations [7–10], but there is no available investigation on the crystallographic evolution during hot deformation supporting the conclusions in the literature. Although optical microscopy is conventionally used to investigate the microstructure, texture analysis is a much more powerful method to investigate the microstructure evolution of deformed materials [11] and to understand the deformation mechanisms [12,13] because it gives more insights and information about orientations and misorientations of grain boundaries. Thus, in the present study, texture evolutions during uniaxial compression tests at 723 and 973 K of CP Ti have been investigated by electron diffraction (EBSD) in a scanning electron microscope (SEM). The two temperatures were chosen because they are widely used in the forming of CP Ti and its alloys [14–15]. All the potential slip systems that may affect the texture evolution during the deformation process are investigated. The mechanism of grain size refinement is analysed based on the distributions of misorientation angles. 2. Experiments The as-received CP Ti, with chemical composition (wt.%) of C, 0.01; Fe, 0.01; H, 0.0005; N, 0.004; O, 0.12; and Ti remaining in balance was supplied by Sandvik Materials Technology, Sandviken,
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Fig. 1. Schematic illustration of as-received material and cutting of samples for the present work.
Sweden in form of a thick-walled extruded tube (Fig. 1). The material was annealed at 1073 K for one hour in argon atmosphere and air cooled in order to obtain new grains and remove residual stresses. With this procedure, an equiaxed microstructure having an average grain size of ∼40 m was obtained which is shown in Fig. 2(a). Cylindrical specimens, 10 mm in diameter and 15 mm in height, were machined from the annealed material. The compression axis of the specimens was always parallel to the extrusion direction of the supplied tube as shown in Fig. 1. Hot compression was conducted using a Gleeble 1500 thermal simulator. Prior to hot compression, each specimen was heated to the deformation temperature for three minutes to ensure a homogenous temperature distribution throughout the specimen. The deformation temperature was measured by a thermocouple, spot welded on the centre of the specimen surface. A tantalum foil with a thickness of 0.05 mm was put between the anvil and the specimen in order to reduce the friction and to prevent sticking. The deformation strain, temperature and strain rate were automatically controlled and recorded by the Gleeble 1500 thermal simulator system. Compression tests were conducted at 723 K with strain rates of 0.01 and 0.1 s−1 and at 973 K with a strain rate of 0.01 s−1 up to a total reduction of 60%. The samples were cooled in water immediately after the compressions were stopped, in order to retain the deformed microstructures. The deformed specimens were cut along their axis and prepared by mechanical grinding using grit papers with different particle sizes from 320 to 1200 mesh. In order to achieve the surface quality required for EBSD examinations, electropolishing was then conducted at 17 V and room temperature in a solution consisting of 25 ml sulphuric acid, 15 ml hydrofluoric acid and 60 ml glacial acetic acid. The prepared samples were analysed with a JSM-7000F FEGSEM operating at 15 kV and equipped with an EBSD orientation imaging system. The Kikuchi patterns were obtained with an ultrasensitive CCD camera and processed with the Channel 5 software from HKL technology. 3. Results and discussions 3.1. Starting material The EBSD image of grain boundaries in Fig. 2(a) shows that the starting material comprised of equiaxed CP Ti with an average grain size of ∼40 m. The colour code, red for (0 0 0 1), green for 1 1 2¯ 0 and blue for 1 0 1¯ 0 , gives the crystallographic direction of each grain that is pointing in the intended compression direction, CD, (previous extrusion direction). The same colour code is used in the inverse pole figure, shown as an insert in Fig. 2(a), and naturally, the high intensity is found in the green and blue regions. The absence of the red colour indicates that (0 0 0 1) is not lying in the compression direction. This is verified by the direct (0 0 0 1) pole figure in Fig. 2(b) showing a band perpendicular to CD. As seen, there are two maxima in the band. The strongest one (8.36 × random) is almost parallel to the radial direction, RD, whereas the other is rotated towards
Fig. 2. The microstructure and texture of the annealed CP Ti revealed by EBSD band contrastmaps in the RD–CD plane. (a) The starting material, (b) the {0 0 0 1} and 1 0 1¯ 0 pole figures and (c) the sample orientation related to the compression axis, CD, and the radial direction, RD with the dominating texture components. The colours of the arrows in (c) match the colours of the grains in (a). (For interpretation of the references to colour in the text, the reader is referred to the web version of the article.)
the normal direction. The first, strongest maximum, results in the three maxima located on the horizontal line of the 1 0 1¯ 0 direct pole figure (Fig. 2(b), whereas the second, rotated maxima, gives an increased intensity in the upper part of the same pole figure. Thus, the original, recrystallized texture can be described as a 1 1 2¯ 0 fibre texture with respect to CD, with the (0 0 0 1) direction perpendicular to CD. A schematic sketch of this texture indicating the strongest texture components is given in Fig. 2(c). It should be pointed out that the EBSD image in Fig. 2(a) evidently shows that the material consists of 100% ˛-phase and that there is no ˇ-phase present. As both the annealing and deformation temperatures were below the transformation temperature (1157 K), the possible influence of ˇ-phase on the deformation behaviour is eliminated. 3.2. Texture evolution in the refined and coarse grains After 60% reduction at 723 and 973 K, the observed microstructures displayed a mixture of refined and coarse grains as shown in
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Fig. 3(a–c). The analysed areas were taken from the central part of the samples subjected to the largest amount of deformation. The coarse grains were elongated in the RD. The pole figures of Fig. 3(d–i) show the orientation distributions in the areas marked A–F, respectively. As seen in the figure, areas A, C and E were taken from fine-grained regions whereas areas B, D and F were taken from coarse grained regions. Naturally, fine and coarse grains have a relative meaning since the EBSD images have different magnification and spatial resolution in order to catch the details of the deformed material. The texture components can be characterized by three Euler angles ϕ1 , ˚ and ϕ2 defined according to the Matthies–Roe convention [16]. As seen, the measured pole figures in Fig. 3(d–i) are compared to ideal texture components reporting the corresponding Euler angles and the angle between the (0 0 0 1) direction and CD, . The refined grains in the areas marked A, C and E were observed to have texture components of A: {ϕ1 = 90◦ , ˚ = 120◦ , ϕ1 = 30◦ }, C:{ϕ1 = 90◦ , ˚ = 130◦ , ϕ1 = 8◦ }, {ϕ1 = 115◦ , ˚ = 80◦ , ϕ1 = 20◦ } and E:{ϕ1 = 90◦ , ˚ = 55◦ , ϕ1 = 30◦ }, respectively. Obviously area C contains a mixture of two texture components. One can be identified as a typical fine-grained texture whereas the other can be identified as a typical coarse one. Typically, for the fine-grained texture components, the angle is between 41 and 46◦ , irrespective of deformation temperature and strain rate. In contrast, the areas with coarse grains, marked B, D and F, had texture components of B:{ϕ1 = 110◦ , ˚ = 90◦ , ϕ1 = 0◦ }, D:{ϕ1 = 65◦ ,
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˚ = 80◦ , ϕ1 = 30◦ }and F:{ϕ1 = 90◦ , ˚ = 125◦ , ϕ1 = 30◦ }, respectively, forming angles of 20.0, 38.4 and 45.8◦ , with the compression axis. These results correspond to 723, 723 and 973 K and to 0.1, 0.01 and 0.01 s−1 , respectively. Thus, the more efficient recovery of the material, the closer the angle falls to 45◦ . The present results indicate that the [0 0 0 1] directions of the deformed grains were preferentially oriented 45◦ away from the compression axis. The possible explanation for this geometrical preference is presented in Fig. 4 indicating the optimum orientation for a base plane subjected to plastic deformation. It is well known that the maximum shear stress is found on a plane 45◦ away from the compression axis. Basal slip has low critically resolved shear stress, CRSS, at elevated temperatures and will easily be activated under compressive loading. During the course of deformation the basal planes will rotate to a more geometrically preferential orientation, i.e. to the aforementioned 45◦ orientation, shown in Fig. 4. Apparently, the efficiency in rotation depends on temperature, strain rate and grain size. For the fine grains, rotation seems very efficient and all of them have orientations close to 45◦ . However, for the coarse grains rotation is more difficult and is promoted by high temperature and low strain rate. The pole figures presented in Figs 2(b) and 3(d–i) indicate that the observed rotation and texture development is caused by basal slip. It is obvious that the basal planes of the grains rotated from the initial 90◦ angle towards a 45◦ angle with the compression direction, CD. Thus, basal slip seems to be an important deformation process during deformation of CP Ti at elevated temperatures.
Fig. 3. The microstructure and texture of CP Ti revealed by EBSD band contrast maps in the RD–CD plane after 60% deformation at (a): 723 K and 0.1 s−1 , (b) 723 K and 0.01 s−1 and (c) 973 K and 0.01 s−1 . The colours of the microstructure and inserted inverse pole figures indicate the orientation relative to CD. The experimental pole figures (d)–(i) correspond to the areas A–F marked in (a)–(c). The accompanying ideal textures indicate the Euler angles and , the misorientation angle with CD. Solid and open circles signify textures for fine and coarse grains, respectively. The white arrows in (a) mark 1 0 1¯ 0 twinning. The scale bars in (a)–(c) indicate lengths of 20, 10 and 100 m, respectively. (For interpretation of the references to colour in the text related to figure, the reader is referred to the web version of the article.)
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Fig. 3. (Continued ).
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the correlations between the slip systems and the Schmid factors under various deformation conditions. Schmid’s law can be written as cr = cos cos
Fig. 4. Schematic illustration of sample and favorable orientation for the [0 0 0 1] direction of deformed grains. The maximum shear stress occurs on the shaded surface forming a 45◦ angle with CD.
Activation of other slip systems cannot be identified directly from the analysis of the pole figures. Therefore, in the following section, the Schmid factor will be used to analyse how the individual slip systems activate and how their activities evolve under various deformation conditions. 3.3. Analysis of slip systems using the Schmid factor In pure hcp titanium, the commonly observed slip systems include the basal slip (0 0 0 1) 1 1 2¯ 0 , the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 , the pyramidal slip 1 0 1¯ 1 1 1 2¯ 0 and the
1st and 2nd order pyramidal slip, viz. 1 0 1¯ 1 1 1 2¯ 3 and 1 1 2¯ 2 1 1 2¯ 3 [17]. The operating slip systems are generally determined by the Von Mises criterion, the Schmid factors, the critical resolved shear stress (CRSS) being dependent on temperature and strain rate, in which the Schmid factors are related to the texture and stress status in the sample. Thus, it is of interest to investigate
(1)
where cr is the CRSS of a given slip system, which is significantly different for the various slip systems in hcp crystals [18]. is the applied stress in uniaxial compression, and are the angles between the loading axis and the shear direction and slip plane normal, respectively. The orientation data of the grains were obtained by EBSD experiments. Based on the orientation data, the distributions of the Schmid factors in the grains can be calculated as the compression direction is given. Fig. 5(a–d) shows the orientationimaging microscopy (OIM) images for the starting material and the deformed samples at 723 K with strain rates of 0.1 and 0.01 s−1 and at 973 K with a strain rate of 0.01 s−1 . The maps are contoured with gray levels of the Schmid factor for basal slip, black corresponding to 0 and white to 0.5. Thus, the brighter the image, the more geometrically favorable orientation for basal slip. The frequency distributions of the Schmid factors for basal slip were calculated and the results are illustrated by the insets of the images. It can be seen that the original orientations of the basal systems are unfavorable because the Schmid factors are relatively low. However, the basal slip became more geometrically favorable after the deformations. This can also be observed directly from the brightness increase from the undeformed grains in Fig. 5(a) to the deformed grains in Fig. 5(b–d). Thus, the analysis indicates that the deformation forces the basal slip to rotate to a more favorable orientation, which is consistent with the conclusion in previous section.
Fig. 5. Schmid factor analyses for the starting material, (a), and the samples deformed at (b) 723 K and 0.1 s−1 (c) 723 K and 0.01 s−1 and (d) 973 K and 0.01 s−1 . The gray levels, black corresponding to 0 and white to 0.5, indicate the Schmid factor for basal slip. The inserts show the frequency distribution of the Schmid factors.
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Fig. 6. Relative frequency of Schmid factor values before and after deformation for (a) basal slip, (b) prismatic slip, (c) pyramidal slip, (d) 1st order pyramidal slip and (e) 2nd order pyramidal slip.
In order to investigate the evolutions of the individual slip systems during the compression tests, Fig. 6(a–e) was constructed showing the frequency distributions of the Schmid factors for the possible slip systems. Each figure shows the Schmid factor distribution of a given slip system before and after deformation. As seen from Fig. 6(a) the initial orientation has a high frequency of low Schmid factors and a low frequency of high Schmid factors for basal slip, which is consistent with the (0 0 0 1) plane being orientated 90◦ to the CD. After deformation, on the other hand, the situation is reversed. The deformed curves show a low frequency of low Schmid factors and a high frequency of high Schmid factors, that deformation in the basal slip sys indicating tem (0 0 0 1) 1 1 2¯ 0 became more geometrically favorable after the deformation. From Fig. 6(b), it can be concluded that the prismatic planes of the starting material are favorably oriented compared to the loading axis. However, after deformations, they have
rotated to less favorable orientations i.e. the high frequency range of the Schmid factors has moved from high to low values. Fig. 6(c) shows that the orientation for pyramidal slip becomes less favored during deformation. In the starting material the highest frequency occurred at 0.4 whereas in the deformed material it has decreased to around 0.35. In contrast to a-type slip, i.e. basal, prismatic and pyramidal slip, discussed above, there is no notable change for the 1st order pyramidal slip with deformation, as shown in Fig. 6(d). It means that the 1st order pyramidal slip always keeps a favorable orientation. For the 2nd order pyramidal slip, reorientations can be both favorable and unfavorable as shown in Fig. 6(e). At 723 K and 0.1 s−1 the texture 1 1 2¯ 2 1 1 2¯ 3 remains in favorable orientation, whereas the other two deformation conditions show a Schmid-factor frequency-peak decreasing from 0.47 to about 0.3 with a considerable broadening.
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Fig. 7. Frequency-misorientation curves in starting material, (a), and after deformation at (b) 723 K and 0.1 s−1 , (c) 723 and 0.01 s−1 and (d) 973 and 0.01 s−1 .
It has been reported that the prismatic slip system
1 0 1¯ 0 1 1 2¯ 0 acts as the primary slip system at room temperature and that the basal slip is activated less frequently [19]. However, the present results suggest that the basal slip system is one of the pronounced deformation modes, which is consistent with results reported by Williams et al. [20] They pointed out that the slip mechanism varies with increasing temperature and that an increased temperature can suppress deformation twinning and promote basal slip. Thus, in the present work at the early stages of deformation, the prismatic planes were favorably oriented for slip having a Schmid factor between 0.4 and 0.5, as shown in Fig. 6(b). Because of that, and that the prismatic slip systems have the lowest CRSS in pure titanium [6], it is most likely that the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 played a vital role in the texture evolution at the early stage of deformation in the present work. After the deformation, the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 becomes inactive under the current stress situation. As the basal, prismatic and pyramidal slip have the common slip direction 1 1 2¯ 0 , being perpendicular to the c-axis, slip on the three systems cannot induce any elongation along the c-axis. In order to accommodate straining in the c-direction, slip on the 1st or 2nd order pyramidal systems, or twinning, must be activated. At low temperatures, deformation twinning is the dominant mechanism allowing inelastic deformation in the [21]. However, c-direction at elevated temperatures, slip in the 1 1 2¯ 3 directions becomes possible, and the slip planes containing this slip direction are the first-order pyramidal 1 0 1¯ 1 and the second-order pyramidal 1 1 2¯ 2 planes. Thus, it can be concluded that activation of slip on these systems is the reason why they keep a favorable orientation during deformation. 3.4. The distribution of the misorientation angle Fig. 7 illustrates the distributions of the misorientation angles in the starting material and the deformed samples. The results indi-
cate that the dominant grain boundaries in the starting material are high angle grain boundaries, HAGBs, with a misorientation angle above 15◦ . After deformation, the frequencies of low angle grain boundaries, LAGBs, ranging from 2 to 15◦ increased dramatically with a simultaneous decrease of the HAGBs, as shown in Fig. 7(b–d). In Fig. 7(b), it can be seen that the distributions of HAGBs between 50 and 60◦ exhibit a relatively high frequency. This is due to the exis tence of the 1 0 1¯ 1 twinning as marked in Fig. 3(a). As reported by Chun et al. [22], the formation of such twins is often observed in pure titanium when deformed above 673 K. The twins form a misorientation angle of 57.26◦ between matrix and twin for CP Ti [23]. This angle is well defined during formation of the twins but is gradually changed due to dislocation pile-ups in the twin boundaries at subsequent deformation. As a result, a broader peak appears rather than a thin spike. In the present work it is about 10◦ wide after deformation at 723 K and 0.1 s−1 . However, at 723 K and 0.01 s−1 , Fig. 7(c) and 973 K and 0.01 s−1 , Fig. 7(d), no obvious peak is found in the same region, indicating that twinning is suppressed by low strain rates and higher temperature, in accordance with Williams et al. [20]. The EBSD results indicate that the grains in CP Ti are effectively refined from equiaxial to inhomogeneous grains. The deformed microstructures consist of mixtures of HAGBs and LAGBs, which cannot be observed or identified properly using optical microscopy or SEM. Only an inhomogeneous mixture of fine and coarse grains can be observed by these methods [24]. However, by using the more sophisticated EBSD-method, as done in the present work, it is possible to separate HAGBs and LAGBs directly from the mapped orientation images, rendering detailed information about the grain structure and misorientations. Thus, by a careful examination of Fig. 3(a–c) it was found that the boundaries of the coarse grains were serrated and elongated along the radial directions of the samples. The wavelengths of these serrations were close to the subgrain sizes inferring that LAGBs gradually turn into HAGBs with increased strain. As
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classic recrystallized nuclei, free of dislocations and equiaxed, were not observed in the present experiments and LAGBs were dominant in the deformed samples, discontinuous dynamic recrystallization (DRX) was obviously not operating. Instead, dynamic recovery appears as the main deformation mechanism. This can be expected from the high stacking fault energy (SFE) of Ti, 300 mJ/m2 [25], as it results in a high degree of recoverability.
4. Conclusions The EBSD technique was employed to investigate the texture and microstructure evolution of CP Ti deformed at 723 K with strain rates of 0.01 and 0.1 s−1 and at 973 K with a strain rate of 0.01 s−1 . The following conclusions can be drawn from the studies: (1) The pre-existing strong texture with (0 0 0 1) base plane perpendicular to the compression axis in the starting material tends to evolve into orientations forming 45◦ angle to the compression axis after deformation. (2) By analysing the Schmid factors of the possible slip systems before and after deformation, it was found that that the prismatic slip system 1 0 1¯ 0 1 1 2¯ 0 dominates during the early stages of deformation. As deformation proceeds, the basal slip system (0 0 0 1) 1 1 2¯ 0 becomes another prevalent deformation mode. Throughout the entire deformation, the pyramidal slip systems keep afavorable for orientation activation. However, the pyramidal 1 0 1¯ 1 1 1 2¯ 0 and the second-order pyramidal slip systems 1 1 2¯ 2 1 1 2¯ 3 tend to become slightly less favored. (3) The distributions of the misorientation angles show that LAGBs are dominant in the deformed samples. At 723 K and 0.1 s−1 ¯ 1 0 1 1 twinning resulted in an increased frequency distribution of the misorientation angle in the range of 50–60◦ . (4) The deformed microstructure implies that dynamic recovery is the dominant deformation mechanism.
Acknowledgements The authors gratefully acknowledge the financial support by the China Scholarship Council. AB Sandvik Materials Technology, Sandviken, Sweden is thanked for supplying the material. References [1] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford, 1996. [2] R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juul Jensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen, A.D. Rollett, Mater. Sci. Eng. A 238 (1997) 219–274. [3] V.V. Stolyarov, Y.T. Zhu, I.V. Alexandrov, T.C. Lowe, R.Z. Valiev, Mater. Sci. Eng. A 343 (2003) 43–50. [4] V.V. Stolyarov, L. Zeipper, B. Mingler, M. Zehetbauer, Mater. Sci. Eng. A 476 (2008) 98–105. [5] K. Morsi, S. Goyal, J. Alloy Compd. 429 (2007) L1–L4. [6] D.H. Shin, I. Kim, J. Kim, Y.S. Kim, S.L. Semiatin, Acta Mater. 51 (2003) 983–996. [7] S. Nemat-Nasser, W.G. Guo, J.Y. Cheng, Acta Mater. 47 (1999) 3705–3720. [8] J.R. Donoso, P.G. Watson, R.E. Reed-Hill, Metall. Trans. A 10 (1979) 1165–1171. [9] S. Venugopal, P. Venugopal, S.L. Mannan, J. Mater. Process. Techol. 202 (2008) 201–215. [10] Senkov, J.J. Jonas, Metall. Mater. Trans. A 27A (1996) 1877–1887. [11] O. Engler, V. Randle, Texture Analysis: Macrotexture, Microtexture and Orientation Mapping, Gordon and Breach, Amsterdam, Holland, 2000. [12] T. Al-Samman, G. Gottstein, Mater. Sci. Eng. A 488 (2008) 406–414. [13] J.A. del Valle, O.A. Ruano, Mater. Sci. Eng. A 487 (2008) 473–480. [14] V.V. Stolyarov, Y.T. Zhu, T.C. Lowe, R.K. Islamgaliev, R.Z. Valiev, Nanostruct. Mater. 11 (1999) 947–954. [15] I.P. Semenova, G.I. Raab, L.R. Saitova, R.Z. Valiev, Mater. Sci. Eng. A 387–389 (2004) 805–808. [16] S. Matthies, G.W. Vinel, Phys. Status Solidi (b) 112 (1982) K111–K114. [17] S. Balasubramanian, L. Anand, Acta Mater. 50 (2002) 133–148. [18] P.G. Partridge, Metall. Rev. 12 (1967) 169–194. [19] M.H. Yoo, Metall. Trans. A12 (1981) 409–418. [20] J.C. Williams, R.G. Baggerly, N.E. Paton, Metall. Mater. Trans. 33A (2002) 837–850. [21] A. Staroselskya, L. Anand, Int. J. Plasticity 19 (2003) 1843–1864. [22] Y.B. Chun, S.H. Yu, S.L. Semiatin, S.K. Hwang, Mater. Sci. Eng. A 398 (2005) 209–219. [23] J.W. Christian, S. Mahajan, Progr. Mater. Sci. 39 (1995) 1–157. [24] L. Jin, D.L. Lin, D.L. Mao, X.Q. Zen, W.J. Ding, Mater. Lett. 59 (2005) 2267–2270. [25] D.J. Bacon, J.W. Martin, Philos. Mag. 4 (1981) 883–894.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and texture evolution of commercial pure titanium deformed at elevated temperatures Zhipeng Zeng ∗ , Yanshu Zhang, Stefan Jonsson Department of Materials Science and Engineering, Royal Institute of Technology, SE-100 44 Stockholm, Sweden
a r t i c l e
i n f o
Article history: Received 17 September 2008 Received in revised form 13 January 2009 Accepted 19 January 2009 Keywords: Commercial pure titanium Texture Schmid factor Misorientation Hot compression Slip system
a b s t r a c t Microstructure and texture evolution of commercial pure titanium were investigated by electron backscattered diffraction (EBSD) after compression tests at elevated temperatures. The basal planes of both the fine and coarse grains in the deformed samples tend to rotate from the initial orientations, perpendicular to the compression axis, to an inclination of 45◦ . The Schmid factor is used to analyse how the individual slip systems activate and how their activities evolve under various deformation conditions. After deformation, the distribution frequency of the misorientation angles shows that the low angle grain boundaries increased dramatically while the high angle grain boundaries decreased. In particular, after deformation at 723 K and 0.1 s−1, a peak around 50–60◦ in the misorientation frequency-distribution is found, which is due to 1 0 1¯ 1 twinning. The analysis of the deformed microstructure indicates that dynamic recovery is the dominant deformation mechanism for commercial pure titanium when subjected to the investigated deformation conditions. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Mechanical properties of metals are dependent on the microstructure and texture formed during processing. In order to control the final properties of the products it is thus necessary to understand the microstructure and texture evolution during plastic deformation. Numerous efforts have been invested to investigate the microstructure and texture development in BCC and FCC metals. In contrast, there is little attention paid to the hexagonal close packed (HCP) metals. As a typical hexagonal close packed metal, commercial pure titanium (CP Ti) is of great importance in many industrial applications due to its highly attractive properties, such as good deformability at high temperature, low density, high biocompatibility and excellent corrosion resistance [1–2]. However, it is hard for CP Ti to strengthen to a level comparable to Ti–6Al–4V [3], which is currently the material of choice for most medical implants. Therefore, to develop higher strength CP Ti is an attractive work for medical applications. In order to increase the strength of CP Ti, large deformation with an accompanying grain refinement is an effective approach. In the present work, the uniaxial compression test was used because it is a convenient method for obtaining large strains and test specimens are easy to prepare. In addition, the deformation process is similar to forging; hence, it has become a commonly used approach
∗ Corresponding author. Tel.: +46 8790 6542; fax: +46 8203107. E-mail address: [email protected] (Z. Zeng). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.01.065
for testing mechanical properties during hot-working processes. However, most of the available investigations focus on severe deformation of CP Ti and its alloys [4–6]. Admittedly, there are some investigations on the deformation mechanisms of CP Ti at elevated temperatures based on the analysis of flow stress–strain curves or optical observations [7–10], but there is no available investigation on the crystallographic evolution during hot deformation supporting the conclusions in the literature. Although optical microscopy is conventionally used to investigate the microstructure, texture analysis is a much more powerful method to investigate the microstructure evolution of deformed materials [11] and to understand the deformation mechanisms [12,13] because it gives more insights and information about orientations and misorientations of grain boundaries. Thus, in the present study, texture evolutions during uniaxial compression tests at 723 and 973 K of CP Ti have been investigated by electron diffraction (EBSD) in a scanning electron microscope (SEM). The two temperatures were chosen because they are widely used in the forming of CP Ti and its alloys [14–15]. All the potential slip systems that may affect the texture evolution during the deformation process are investigated. The mechanism of grain size refinement is analysed based on the distributions of misorientation angles. 2. Experiments The as-received CP Ti, with chemical composition (wt.%) of C, 0.01; Fe, 0.01; H, 0.0005; N, 0.004; O, 0.12; and Ti remaining in balance was supplied by Sandvik Materials Technology, Sandviken,
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Fig. 1. Schematic illustration of as-received material and cutting of samples for the present work.
Sweden in form of a thick-walled extruded tube (Fig. 1). The material was annealed at 1073 K for one hour in argon atmosphere and air cooled in order to obtain new grains and remove residual stresses. With this procedure, an equiaxed microstructure having an average grain size of ∼40 m was obtained which is shown in Fig. 2(a). Cylindrical specimens, 10 mm in diameter and 15 mm in height, were machined from the annealed material. The compression axis of the specimens was always parallel to the extrusion direction of the supplied tube as shown in Fig. 1. Hot compression was conducted using a Gleeble 1500 thermal simulator. Prior to hot compression, each specimen was heated to the deformation temperature for three minutes to ensure a homogenous temperature distribution throughout the specimen. The deformation temperature was measured by a thermocouple, spot welded on the centre of the specimen surface. A tantalum foil with a thickness of 0.05 mm was put between the anvil and the specimen in order to reduce the friction and to prevent sticking. The deformation strain, temperature and strain rate were automatically controlled and recorded by the Gleeble 1500 thermal simulator system. Compression tests were conducted at 723 K with strain rates of 0.01 and 0.1 s−1 and at 973 K with a strain rate of 0.01 s−1 up to a total reduction of 60%. The samples were cooled in water immediately after the compressions were stopped, in order to retain the deformed microstructures. The deformed specimens were cut along their axis and prepared by mechanical grinding using grit papers with different particle sizes from 320 to 1200 mesh. In order to achieve the surface quality required for EBSD examinations, electropolishing was then conducted at 17 V and room temperature in a solution consisting of 25 ml sulphuric acid, 15 ml hydrofluoric acid and 60 ml glacial acetic acid. The prepared samples were analysed with a JSM-7000F FEGSEM operating at 15 kV and equipped with an EBSD orientation imaging system. The Kikuchi patterns were obtained with an ultrasensitive CCD camera and processed with the Channel 5 software from HKL technology. 3. Results and discussions 3.1. Starting material The EBSD image of grain boundaries in Fig. 2(a) shows that the starting material comprised of equiaxed CP Ti with an average grain size of ∼40 m. The colour code, red for (0 0 0 1), green for 1 1 2¯ 0 and blue for 1 0 1¯ 0 , gives the crystallographic direction of each grain that is pointing in the intended compression direction, CD, (previous extrusion direction). The same colour code is used in the inverse pole figure, shown as an insert in Fig. 2(a), and naturally, the high intensity is found in the green and blue regions. The absence of the red colour indicates that (0 0 0 1) is not lying in the compression direction. This is verified by the direct (0 0 0 1) pole figure in Fig. 2(b) showing a band perpendicular to CD. As seen, there are two maxima in the band. The strongest one (8.36 × random) is almost parallel to the radial direction, RD, whereas the other is rotated towards
Fig. 2. The microstructure and texture of the annealed CP Ti revealed by EBSD band contrastmaps in the RD–CD plane. (a) The starting material, (b) the {0 0 0 1} and 1 0 1¯ 0 pole figures and (c) the sample orientation related to the compression axis, CD, and the radial direction, RD with the dominating texture components. The colours of the arrows in (c) match the colours of the grains in (a). (For interpretation of the references to colour in the text, the reader is referred to the web version of the article.)
the normal direction. The first, strongest maximum, results in the three maxima located on the horizontal line of the 1 0 1¯ 0 direct pole figure (Fig. 2(b), whereas the second, rotated maxima, gives an increased intensity in the upper part of the same pole figure. Thus, the original, recrystallized texture can be described as a 1 1 2¯ 0 fibre texture with respect to CD, with the (0 0 0 1) direction perpendicular to CD. A schematic sketch of this texture indicating the strongest texture components is given in Fig. 2(c). It should be pointed out that the EBSD image in Fig. 2(a) evidently shows that the material consists of 100% ˛-phase and that there is no ˇ-phase present. As both the annealing and deformation temperatures were below the transformation temperature (1157 K), the possible influence of ˇ-phase on the deformation behaviour is eliminated. 3.2. Texture evolution in the refined and coarse grains After 60% reduction at 723 and 973 K, the observed microstructures displayed a mixture of refined and coarse grains as shown in
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Fig. 3(a–c). The analysed areas were taken from the central part of the samples subjected to the largest amount of deformation. The coarse grains were elongated in the RD. The pole figures of Fig. 3(d–i) show the orientation distributions in the areas marked A–F, respectively. As seen in the figure, areas A, C and E were taken from fine-grained regions whereas areas B, D and F were taken from coarse grained regions. Naturally, fine and coarse grains have a relative meaning since the EBSD images have different magnification and spatial resolution in order to catch the details of the deformed material. The texture components can be characterized by three Euler angles ϕ1 , ˚ and ϕ2 defined according to the Matthies–Roe convention [16]. As seen, the measured pole figures in Fig. 3(d–i) are compared to ideal texture components reporting the corresponding Euler angles and the angle between the (0 0 0 1) direction and CD, . The refined grains in the areas marked A, C and E were observed to have texture components of A: {ϕ1 = 90◦ , ˚ = 120◦ , ϕ1 = 30◦ }, C:{ϕ1 = 90◦ , ˚ = 130◦ , ϕ1 = 8◦ }, {ϕ1 = 115◦ , ˚ = 80◦ , ϕ1 = 20◦ } and E:{ϕ1 = 90◦ , ˚ = 55◦ , ϕ1 = 30◦ }, respectively. Obviously area C contains a mixture of two texture components. One can be identified as a typical fine-grained texture whereas the other can be identified as a typical coarse one. Typically, for the fine-grained texture components, the angle is between 41 and 46◦ , irrespective of deformation temperature and strain rate. In contrast, the areas with coarse grains, marked B, D and F, had texture components of B:{ϕ1 = 110◦ , ˚ = 90◦ , ϕ1 = 0◦ }, D:{ϕ1 = 65◦ ,
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˚ = 80◦ , ϕ1 = 30◦ }and F:{ϕ1 = 90◦ , ˚ = 125◦ , ϕ1 = 30◦ }, respectively, forming angles of 20.0, 38.4 and 45.8◦ , with the compression axis. These results correspond to 723, 723 and 973 K and to 0.1, 0.01 and 0.01 s−1 , respectively. Thus, the more efficient recovery of the material, the closer the angle falls to 45◦ . The present results indicate that the [0 0 0 1] directions of the deformed grains were preferentially oriented 45◦ away from the compression axis. The possible explanation for this geometrical preference is presented in Fig. 4 indicating the optimum orientation for a base plane subjected to plastic deformation. It is well known that the maximum shear stress is found on a plane 45◦ away from the compression axis. Basal slip has low critically resolved shear stress, CRSS, at elevated temperatures and will easily be activated under compressive loading. During the course of deformation the basal planes will rotate to a more geometrically preferential orientation, i.e. to the aforementioned 45◦ orientation, shown in Fig. 4. Apparently, the efficiency in rotation depends on temperature, strain rate and grain size. For the fine grains, rotation seems very efficient and all of them have orientations close to 45◦ . However, for the coarse grains rotation is more difficult and is promoted by high temperature and low strain rate. The pole figures presented in Figs 2(b) and 3(d–i) indicate that the observed rotation and texture development is caused by basal slip. It is obvious that the basal planes of the grains rotated from the initial 90◦ angle towards a 45◦ angle with the compression direction, CD. Thus, basal slip seems to be an important deformation process during deformation of CP Ti at elevated temperatures.
Fig. 3. The microstructure and texture of CP Ti revealed by EBSD band contrast maps in the RD–CD plane after 60% deformation at (a): 723 K and 0.1 s−1 , (b) 723 K and 0.01 s−1 and (c) 973 K and 0.01 s−1 . The colours of the microstructure and inserted inverse pole figures indicate the orientation relative to CD. The experimental pole figures (d)–(i) correspond to the areas A–F marked in (a)–(c). The accompanying ideal textures indicate the Euler angles and , the misorientation angle with CD. Solid and open circles signify textures for fine and coarse grains, respectively. The white arrows in (a) mark 1 0 1¯ 0 twinning. The scale bars in (a)–(c) indicate lengths of 20, 10 and 100 m, respectively. (For interpretation of the references to colour in the text related to figure, the reader is referred to the web version of the article.)
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Fig. 3. (Continued ).
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the correlations between the slip systems and the Schmid factors under various deformation conditions. Schmid’s law can be written as cr = cos cos
Fig. 4. Schematic illustration of sample and favorable orientation for the [0 0 0 1] direction of deformed grains. The maximum shear stress occurs on the shaded surface forming a 45◦ angle with CD.
Activation of other slip systems cannot be identified directly from the analysis of the pole figures. Therefore, in the following section, the Schmid factor will be used to analyse how the individual slip systems activate and how their activities evolve under various deformation conditions. 3.3. Analysis of slip systems using the Schmid factor In pure hcp titanium, the commonly observed slip systems include the basal slip (0 0 0 1) 1 1 2¯ 0 , the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 , the pyramidal slip 1 0 1¯ 1 1 1 2¯ 0 and the
1st and 2nd order pyramidal slip, viz. 1 0 1¯ 1 1 1 2¯ 3 and 1 1 2¯ 2 1 1 2¯ 3 [17]. The operating slip systems are generally determined by the Von Mises criterion, the Schmid factors, the critical resolved shear stress (CRSS) being dependent on temperature and strain rate, in which the Schmid factors are related to the texture and stress status in the sample. Thus, it is of interest to investigate
(1)
where cr is the CRSS of a given slip system, which is significantly different for the various slip systems in hcp crystals [18]. is the applied stress in uniaxial compression, and are the angles between the loading axis and the shear direction and slip plane normal, respectively. The orientation data of the grains were obtained by EBSD experiments. Based on the orientation data, the distributions of the Schmid factors in the grains can be calculated as the compression direction is given. Fig. 5(a–d) shows the orientationimaging microscopy (OIM) images for the starting material and the deformed samples at 723 K with strain rates of 0.1 and 0.01 s−1 and at 973 K with a strain rate of 0.01 s−1 . The maps are contoured with gray levels of the Schmid factor for basal slip, black corresponding to 0 and white to 0.5. Thus, the brighter the image, the more geometrically favorable orientation for basal slip. The frequency distributions of the Schmid factors for basal slip were calculated and the results are illustrated by the insets of the images. It can be seen that the original orientations of the basal systems are unfavorable because the Schmid factors are relatively low. However, the basal slip became more geometrically favorable after the deformations. This can also be observed directly from the brightness increase from the undeformed grains in Fig. 5(a) to the deformed grains in Fig. 5(b–d). Thus, the analysis indicates that the deformation forces the basal slip to rotate to a more favorable orientation, which is consistent with the conclusion in previous section.
Fig. 5. Schmid factor analyses for the starting material, (a), and the samples deformed at (b) 723 K and 0.1 s−1 (c) 723 K and 0.01 s−1 and (d) 973 K and 0.01 s−1 . The gray levels, black corresponding to 0 and white to 0.5, indicate the Schmid factor for basal slip. The inserts show the frequency distribution of the Schmid factors.
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Fig. 6. Relative frequency of Schmid factor values before and after deformation for (a) basal slip, (b) prismatic slip, (c) pyramidal slip, (d) 1st order pyramidal slip and (e) 2nd order pyramidal slip.
In order to investigate the evolutions of the individual slip systems during the compression tests, Fig. 6(a–e) was constructed showing the frequency distributions of the Schmid factors for the possible slip systems. Each figure shows the Schmid factor distribution of a given slip system before and after deformation. As seen from Fig. 6(a) the initial orientation has a high frequency of low Schmid factors and a low frequency of high Schmid factors for basal slip, which is consistent with the (0 0 0 1) plane being orientated 90◦ to the CD. After deformation, on the other hand, the situation is reversed. The deformed curves show a low frequency of low Schmid factors and a high frequency of high Schmid factors, that deformation in the basal slip sys indicating tem (0 0 0 1) 1 1 2¯ 0 became more geometrically favorable after the deformation. From Fig. 6(b), it can be concluded that the prismatic planes of the starting material are favorably oriented compared to the loading axis. However, after deformations, they have
rotated to less favorable orientations i.e. the high frequency range of the Schmid factors has moved from high to low values. Fig. 6(c) shows that the orientation for pyramidal slip becomes less favored during deformation. In the starting material the highest frequency occurred at 0.4 whereas in the deformed material it has decreased to around 0.35. In contrast to a-type slip, i.e. basal, prismatic and pyramidal slip, discussed above, there is no notable change for the 1st order pyramidal slip with deformation, as shown in Fig. 6(d). It means that the 1st order pyramidal slip always keeps a favorable orientation. For the 2nd order pyramidal slip, reorientations can be both favorable and unfavorable as shown in Fig. 6(e). At 723 K and 0.1 s−1 the texture 1 1 2¯ 2 1 1 2¯ 3 remains in favorable orientation, whereas the other two deformation conditions show a Schmid-factor frequency-peak decreasing from 0.47 to about 0.3 with a considerable broadening.
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Fig. 7. Frequency-misorientation curves in starting material, (a), and after deformation at (b) 723 K and 0.1 s−1 , (c) 723 and 0.01 s−1 and (d) 973 and 0.01 s−1 .
It has been reported that the prismatic slip system
1 0 1¯ 0 1 1 2¯ 0 acts as the primary slip system at room temperature and that the basal slip is activated less frequently [19]. However, the present results suggest that the basal slip system is one of the pronounced deformation modes, which is consistent with results reported by Williams et al. [20] They pointed out that the slip mechanism varies with increasing temperature and that an increased temperature can suppress deformation twinning and promote basal slip. Thus, in the present work at the early stages of deformation, the prismatic planes were favorably oriented for slip having a Schmid factor between 0.4 and 0.5, as shown in Fig. 6(b). Because of that, and that the prismatic slip systems have the lowest CRSS in pure titanium [6], it is most likely that the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 played a vital role in the texture evolution at the early stage of deformation in the present work. After the deformation, the prismatic slip 1 0 1¯ 0 1 1 2¯ 0 becomes inactive under the current stress situation. As the basal, prismatic and pyramidal slip have the common slip direction 1 1 2¯ 0 , being perpendicular to the c-axis, slip on the three systems cannot induce any elongation along the c-axis. In order to accommodate straining in the c-direction, slip on the 1st or 2nd order pyramidal systems, or twinning, must be activated. At low temperatures, deformation twinning is the dominant mechanism allowing inelastic deformation in the [21]. However, c-direction at elevated temperatures, slip in the 1 1 2¯ 3 directions becomes possible, and the slip planes containing this slip direction are the first-order pyramidal 1 0 1¯ 1 and the second-order pyramidal 1 1 2¯ 2 planes. Thus, it can be concluded that activation of slip on these systems is the reason why they keep a favorable orientation during deformation. 3.4. The distribution of the misorientation angle Fig. 7 illustrates the distributions of the misorientation angles in the starting material and the deformed samples. The results indi-
cate that the dominant grain boundaries in the starting material are high angle grain boundaries, HAGBs, with a misorientation angle above 15◦ . After deformation, the frequencies of low angle grain boundaries, LAGBs, ranging from 2 to 15◦ increased dramatically with a simultaneous decrease of the HAGBs, as shown in Fig. 7(b–d). In Fig. 7(b), it can be seen that the distributions of HAGBs between 50 and 60◦ exhibit a relatively high frequency. This is due to the exis tence of the 1 0 1¯ 1 twinning as marked in Fig. 3(a). As reported by Chun et al. [22], the formation of such twins is often observed in pure titanium when deformed above 673 K. The twins form a misorientation angle of 57.26◦ between matrix and twin for CP Ti [23]. This angle is well defined during formation of the twins but is gradually changed due to dislocation pile-ups in the twin boundaries at subsequent deformation. As a result, a broader peak appears rather than a thin spike. In the present work it is about 10◦ wide after deformation at 723 K and 0.1 s−1 . However, at 723 K and 0.01 s−1 , Fig. 7(c) and 973 K and 0.01 s−1 , Fig. 7(d), no obvious peak is found in the same region, indicating that twinning is suppressed by low strain rates and higher temperature, in accordance with Williams et al. [20]. The EBSD results indicate that the grains in CP Ti are effectively refined from equiaxial to inhomogeneous grains. The deformed microstructures consist of mixtures of HAGBs and LAGBs, which cannot be observed or identified properly using optical microscopy or SEM. Only an inhomogeneous mixture of fine and coarse grains can be observed by these methods [24]. However, by using the more sophisticated EBSD-method, as done in the present work, it is possible to separate HAGBs and LAGBs directly from the mapped orientation images, rendering detailed information about the grain structure and misorientations. Thus, by a careful examination of Fig. 3(a–c) it was found that the boundaries of the coarse grains were serrated and elongated along the radial directions of the samples. The wavelengths of these serrations were close to the subgrain sizes inferring that LAGBs gradually turn into HAGBs with increased strain. As
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classic recrystallized nuclei, free of dislocations and equiaxed, were not observed in the present experiments and LAGBs were dominant in the deformed samples, discontinuous dynamic recrystallization (DRX) was obviously not operating. Instead, dynamic recovery appears as the main deformation mechanism. This can be expected from the high stacking fault energy (SFE) of Ti, 300 mJ/m2 [25], as it results in a high degree of recoverability.
4. Conclusions The EBSD technique was employed to investigate the texture and microstructure evolution of CP Ti deformed at 723 K with strain rates of 0.01 and 0.1 s−1 and at 973 K with a strain rate of 0.01 s−1 . The following conclusions can be drawn from the studies: (1) The pre-existing strong texture with (0 0 0 1) base plane perpendicular to the compression axis in the starting material tends to evolve into orientations forming 45◦ angle to the compression axis after deformation. (2) By analysing the Schmid factors of the possible slip systems before and after deformation, it was found that that the prismatic slip system 1 0 1¯ 0 1 1 2¯ 0 dominates during the early stages of deformation. As deformation proceeds, the basal slip system (0 0 0 1) 1 1 2¯ 0 becomes another prevalent deformation mode. Throughout the entire deformation, the pyramidal slip systems keep afavorable for orientation activation. However, the pyramidal 1 0 1¯ 1 1 1 2¯ 0 and the second-order pyramidal slip systems 1 1 2¯ 2 1 1 2¯ 3 tend to become slightly less favored. (3) The distributions of the misorientation angles show that LAGBs are dominant in the deformed samples. At 723 K and 0.1 s−1 ¯ 1 0 1 1 twinning resulted in an increased frequency distribution of the misorientation angle in the range of 50–60◦ . (4) The deformed microstructure implies that dynamic recovery is the dominant deformation mechanism.
Acknowledgements The authors gratefully acknowledge the financial support by the China Scholarship Council. AB Sandvik Materials Technology, Sandviken, Sweden is thanked for supplying the material. References [1] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford, 1996. [2] R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. Juul Jensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen, A.D. Rollett, Mater. Sci. Eng. A 238 (1997) 219–274. [3] V.V. Stolyarov, Y.T. Zhu, I.V. Alexandrov, T.C. Lowe, R.Z. Valiev, Mater. Sci. Eng. A 343 (2003) 43–50. [4] V.V. Stolyarov, L. Zeipper, B. Mingler, M. Zehetbauer, Mater. Sci. Eng. A 476 (2008) 98–105. [5] K. Morsi, S. Goyal, J. Alloy Compd. 429 (2007) L1–L4. [6] D.H. Shin, I. Kim, J. Kim, Y.S. Kim, S.L. Semiatin, Acta Mater. 51 (2003) 983–996. [7] S. Nemat-Nasser, W.G. Guo, J.Y. Cheng, Acta Mater. 47 (1999) 3705–3720. [8] J.R. Donoso, P.G. Watson, R.E. Reed-Hill, Metall. Trans. A 10 (1979) 1165–1171. [9] S. Venugopal, P. Venugopal, S.L. Mannan, J. Mater. Process. Techol. 202 (2008) 201–215. [10] Senkov, J.J. Jonas, Metall. Mater. Trans. A 27A (1996) 1877–1887. [11] O. Engler, V. Randle, Texture Analysis: Macrotexture, Microtexture and Orientation Mapping, Gordon and Breach, Amsterdam, Holland, 2000. [12] T. Al-Samman, G. Gottstein, Mater. Sci. Eng. A 488 (2008) 406–414. [13] J.A. del Valle, O.A. Ruano, Mater. Sci. Eng. A 487 (2008) 473–480. [14] V.V. Stolyarov, Y.T. Zhu, T.C. Lowe, R.K. Islamgaliev, R.Z. Valiev, Nanostruct. Mater. 11 (1999) 947–954. [15] I.P. Semenova, G.I. Raab, L.R. Saitova, R.Z. Valiev, Mater. Sci. Eng. A 387–389 (2004) 805–808. [16] S. Matthies, G.W. Vinel, Phys. Status Solidi (b) 112 (1982) K111–K114. [17] S. Balasubramanian, L. Anand, Acta Mater. 50 (2002) 133–148. [18] P.G. Partridge, Metall. Rev. 12 (1967) 169–194. [19] M.H. Yoo, Metall. Trans. A12 (1981) 409–418. [20] J.C. Williams, R.G. Baggerly, N.E. Paton, Metall. Mater. Trans. 33A (2002) 837–850. [21] A. Staroselskya, L. Anand, Int. J. Plasticity 19 (2003) 1843–1864. [22] Y.B. Chun, S.H. Yu, S.L. Semiatin, S.K. Hwang, Mater. Sci. Eng. A 398 (2005) 209–219. [23] J.W. Christian, S. Mahajan, Progr. Mater. Sci. 39 (1995) 1–157. [24] L. Jin, D.L. Lin, D.L. Mao, X.Q. Zen, W.J. Ding, Mater. Lett. 59 (2005) 2267–2270. [25] D.J. Bacon, J.W. Martin, Philos. Mag. 4 (1981) 883–894.