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A critical test of twin-induced softening in a magnesium alloy extruded to a strain of 0.7 at room temperature
Available online at www.sciencedirect.com
Scripta Materialia 67 (2012) 1015–1018 www.elsevier.com/locate/scriptamat
A critical test of twin-induced softening in a magnesium alloy extruded to a strain of 0.7 at room temperature Zhen Zhang,a,b,⇑ Pavel Cizeka and Matthew Barnetta a
Institute for Frontier Materials, Geelong Technology Precinct, Deakin University, Geelong, Vic 3220, Australia School of Materials Science and Engineering, Central South University, Changsha 410083, People’s Republic of China
b
Received 12 August 2012; revised 19 September 2012; accepted 22 September 2012 Available online 28 September 2012
Ambient extrusion was used to impart different uniaxial strains to a magnesium alloy. The aim was to test the idea that texture change at high strains can lead to work softening. Tensile tests, optical microscopy, X-ray diffraction and transmission electron microscopy were used to investigate the flow stress, microstructure and texture evolution up to a uniaxial strain of 0.7. The microstructure after extrusion is dominated by contraction twinning, which served to weaken the basal texture. However, no softening was observed; global work hardening persisted. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloys; Twinning; Texture; Work hardening
Single crystals of magnesium display very little ductility when compressed along the c-axis [1–6]. Unfortunately, such deformation is common. Polycrystalline magnesium alloys develop a sharp “basal” texture during wrought processing and this texture is characterized by caxis aligned near to perpendicular to the tooling face [7– 12]. This type of deformation activates f1 0 1 1g contraction twinning and subsequent double twinning [3,4,6]. In the twinned volumes, reorientation of the basal planes into the plane of maximum shear stress triggers flow concentration which is followed by failure [13–17]. In addition to this local texture softening effect, a number of workers (e.g. [16]), including one of us [18], has suggested that it may also lead to global softening. Such has been seen by Couling et al. [19] in cold-rolled sheet and the phenomenon may play an important role in tensile ductility. This idea is critically tested in the current work and is shown to be incorrect for the chosen alloy. The intent of the study is to determine the microstructural evolution and flow stress at similar strain levels and strain paths to those at which local necking and failure occur in a tensile test. To achieve this, ambient extrusion was used to impart uniaxial strains up to 0.68. Uniaxial tensile tests were then conducted on these extruded samples. X-ray diffraction (XRD) and
⇑ Corresponding
author at: School of Materials Science and Engineering, Central South University, Changsha 410083, People’s Republic of China.; e-mail: [email protected]
automated crystal orientation mapping in a transmission electron microscope (TEM) were also conducted to characterize twinning and texture evolution. The material used in the present work was commercially available wrought AZ31 (Mg–3% Al–1% Zn) magnesium alloy, received in the form of 30 and 100 mm diameter extruded bars. The material was given two different annealing treatments: 430 °C for 8 h and 300 °C for 30 min to produce grain sizes of 8 and 30 lm, respectively. Conical extrusion dies with the same initial diameter of 17 mm but with different exit diameters of 12, 13, 14, 15 and 16 mm were employed. During ambient extrusion, fracture was evident only in some of the 30 lm samples when the 13 mm die (53% strain) was employed, and our attempts at ambient extrusion using the 12 mm die (70% strain) failed for both grain sizes. In order to introduce higher strains, 10 mm diameter billets were fitted into aluminum tube-like shells prior to extrusion. This method permitted true strains of 0.64 for the 30 lm and 0.68 for the 8 lm material to be achieved without failure using the 12 mm die. Round tensile samples were cut along the extrusion direction with a gauge length of 11.5 mm and a gauge diameter of 2.9 mm. Uniaxial tensile tests were performed on these samples with different extrusion strain for both grain sizes. True stress–strain curves are shown in Figure 1a and b and it can be seen that the flow stress increases with strain. No global softening (drop in yield stress with increase in strain) is evident, unlike that seen by Couling
1359-6462/$ - see front matter Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.scriptamat.2012.09.021
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Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Figure 1. The true stress–strain curves for the (a) 8 lm and (b) 30 lm materials after ambient uniaxial extrusion to different strains. (c) The reduction in area for the 8 and 30 lm materials after different extrusion strain.
et al. [19] for rolling. The Ludwik–Hollomon relationship (r = ken) is fitted to the initial tensile curve and extrapolated to higher strains. This only slightly overestimates the measured behavior. Interestingly, the reduction in area measured after each test showed very little sensitivity to strain (Fig. 1c). There is even some slight increase with strain at the low extrusion reductions. It is as if the extrusion deformation renders the material slightly more ductile. Microstructures were observed for each strain condition, before tensile testing, on the longitudinal section (parallel to the extrusion direction, ED). Both grain size materials display extensive twinning activity (Fig. 2a and b). The twin volume fraction seen at low strains agrees well with previous work [18] and the twins are consistent in appearance with f1 0 1 1g double twins [3,4,6]. However, orientation information for the present twins was difficult to obtain using electron backscatter diffraction (EBSD) in a scanning electron microscope
Figure 2. Optical microstructure of the 30 lm material after ambient uniaxial extrusion to (a) 10% and (b) 25%. (c) The relationship between twinning volume fraction and extrusion strain for the 8 and 30 lm materials.
(SEM), and therefore in the following XRD and TEM orientation measurements are employed for the material with a grain size of 30 lm. Macrotexture measurements were conducted on a Panalytical X-ray diffractometer in Schulz reflection geometry. All the measurements were performed on the plane 45° to ED; in this way, complete pole figure data could be collected [20]. Considering the small 2h angle (around 34.5° for the basal plane) and the limited sample dimensions, a small spot size (0.1 mm 1 mm) was used. Accordingly, the tube voltage and current was set as 40 kV and 50 mA, and the exposure time was 60 s per step. Defocusing and background were corrected and all the measurements were collected under the same conditions. Normalization was performed as follows. Due to the axial symmetry of the fiber texture in the round extrusion samples, the normalization integration was only carried out along one equal-U line from the radius direction (RD) to the extrusion direction (ED) (see Fig. 3). The normalizing factor F [21] could thus be calculated as (w is the angle from RD), based on which the X-ray Rp intensity was divided to 02 I intensity cos wdw acquire normalized random values of x. The (0002) pole density is plotted along one equal-U line (U = 0) in Figure 3 (where the w = 0° position represents RD). Since the X-ray intensity is extremely low at high w angles (w > 40°), we only show the 35° to +35° w range for the U = 0 line. Symmetrization was applied to the data to give a full description of the distribution of the basal intensity around RD. It is clear in the plot that the texture intensity decreases with increasing strain. It is frequently reported [22–25] that the characteristic basal texture strengthens with imposed strain during plastic processing. The weakening of basal texture seen here is most likely a consequence of the high twin fractions formed. TEM analysis was carried out to support this idea. A JEOL 2100F TEM was employed with an accelerating voltage of 200 kV. The specimens were sectioned perpendicular to ED. Figure 4a shows a TEM bright-field image of a single twin formed in a matrix with the typical h1 0 1 0i==ED orientation. The twinning area is broken up into several individual blocks with different orientations. Similar orientations are outlined in the same color in the figure. The orientations of the different blocks within the twinning area share a common h1 1 2 0i pole with the matrix.
Figure 3. The distribution of (0002) pole density along the U = 0 line.
Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Figure 4. (a) TEM image of a single twin; the zone axis ½2 1 1 0 is 30° from the ED. (b) Inverse pole figure map of twinning band area. Both single twin and twin bands were formed in the matrix, which had typical ½1 0 1 0==ED orientation.
Selected-area diffraction patterns were taken separately from all these areas along their common h1 1 2 0i zone axis. Block D has a misorientation of 42° clockwise along ½2 1 1 0 from the matrix, and block A has a misorientation of 36.2° anticlockwise along ½2 1 1 0 from the matrix. Both areas have a close orientation to the ideal f1 0 1 g–f1 0 1 2g double twin, which is positioned at 38° h1 1 2 0i to the matrix. If block A is considered as a f1 0 1 1g–f1 0 1 2g double twinning area, the first twinning event would have occurred on the ð0 1 1 1Þ plane. Alternatively, if block D is considered as a f1 0 1 1g–f1 0 1 2g double twinning area, the first twinning event would have occurred on the ð0 1 1 1Þ plane. From the diffraction pattern, both ð0 1 1 1Þ and ð0 1 1 1Þ plane traces were established and these are shown in Figure 4a. The expected trace for block D agrees with the observed habit of the primary twin, so block D is identified as the initial double twinning area. Blocks A and B were found to share the typical f1 0 1 2g twinning relationship (86° h1 1 2 0i). Since the boundary between block A and block B coincided well with the ð1 0 1 2Þ plane trace in block D, block A could be considered as a triple twinning region within block D. The resulting orientations in the twinned volume are consistent with the spread of the texture seen in the X-ray data. It is evident that the twinning processes taking place in the present material are quite complex and multiple twinning events, with a
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common h1 1 2 0i axis, can form in one volume. Similar complex contraction twins were previously observed by some of the present authors close to the fracture surface of AZ31 tensile specimens [26,27]. To gain a better view of the twinning reorieintations, automatic crystal orientation measurements were performed using the Nanomegas Astar system. Selected areas were scanned in nanobeam mode using a step size of 3 nm. Diffraction patterns were collected and compared with simulated templates to find the best matching solution [28]. An inverse pole figure (IPF) map of a typical twinned band is shown in Figure 4b. High-angle boundaries, including twin boundaries, are highlighted in black. Low-angle boundaries, <10°, are highlighted in white. The blue area, whose orientation lies close to the h1 0 1 0i corner in the ED IPF, corresponds to the matrix of the original grain. No particular orientation relationship is evident between these twinning bands and matrix, but all orientations share a h1 1 2 0i pole with the matrix. The common h1 1 2 0i axis is both perpendicular to the extrusion direction and to the plane of shear if orientations A–F are indeed twins. (The twinning plane of shear contains the twin plane normal and the twin shear direction. For the present material, the normal to the plane of shear is the direction about which the twin crystal misorientations are typically described.) The twin band in Figure 4b displays a distinct substructure but the matrix does not. The twinning area is most likely a region of some degree of localized slip. Such is expected from the basal Schmid factors in Table 1, which are 10 times higher in the twinned area than in the adjoining matrix. This slip activity is likely to have rotated the band materials away from its original orientation. Inspection of the f1 0 1 1g pole figure reveals that the habit of the twin band is consistent with two f1 0 1 1g planes, III and IV, in the pole figure in Figure 4b. Orientation E is within 8° of the primary f1 0 1 1g twin expected for twinning on the IV ð1 0 1 1Þ plane. Orientation A is within 6° of the primary f1 0 1 1g twin expected for twinning on the III ð1 0 1 1Þ plane. It is thus apparent that Figure 4b portrays a region close to an intersection of two f1 0 1 1g twin systems. According to this view, orientations A and C are part of one family, and orientations B, D, E and F are part of another family. However, we acknowledge that other interpretations are possible. Nevertheless, the result confirms the presence of complex multiple twinning events and their tendency to weaken the texture. Although the present work confirms the weakening of the basal texture, no global softening was observed. The flow stresses at high strains were slightly lower than extrapolated predictions, which might point to an effect of texture. However, it is clear that work hardening dominates. This is manifested by the substructure seen in the twin bands and by the fragmentation resulting from multiple twinning events in the twinned volumes. Both of these serve to reduce the slip length. Although not observed here, voids have been seen to form in twin volumes in previous studies [18]. Such void formation can be expected to lead to instability of flow and to rupture during a tensile test [3,29]. In the present
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Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Table 1. Euler angles and basal slip Schmid factor for matrix and different twinning bands. Area
Euler angle (Bunge)
Misorientation from matrix
Schmid factor for basal slip
Matrix A B C D E F
(203, 87.6, 60) (205, 137.4, 241.2) (200, 60.3, 63.6) (203, 117.4, 241.7) (197, 49, 65.4) (24, 140.2, 246.4) (199, 47.3, 60)
– +49.8 27.6 +29.8 39.1 48.2 40.5
0.04 0.44 0.38 0.36 0.45 0.45 0.43
case, the hydrostatic compressive stress component of extrusion prevented such softening from occurring. Indeed, it may be that the fragmentation in the twinned volumes permitted by these circumstances serves to slow the development of twin-induced voiding in subsequent tensile tests. The reduction in area seen in the present tensile samples was preserved after ambient extrusion, despite the higher flow stress levels. Finally, it remains to suggest that the common h1 1 2 0i pole seen in the multiple twinned volumes may relate to the role of basal dislocations in twin nucleation. It has been shown that incoming basal dislocations can lead to secondary twin nucleation at pre-existing twin interfaces [30]. In such a case the nucleating twin is more likely to share with the existing twin the h1 1 2 0i pole that lies in the twin interface. This is the normal to the “plane of shear” identified above. In conclusion, the present study has successfully employed ambient extrusion to attain uniaxial strains as high as 68%. The structures thus produced display copious complex twinning events, which serve to subtly weaken the texture. Global work hardening persists, unlike the behavior observed in cold-rolled magnesium alloy [19]. It is likely that the contribution of contraction twinning to tensile ductility of extruded material lies in the formation of voids in doubly (or tertiary) twinned volumes rather than in global texture softening. The multiply twinned volumes are frequently seen to share a common h1 1 2 0i pole, due possibly to the role of basal dislocation–twin boundary interactions in twin nucleation. The authors acknowledge Professor Bevis Hutchison for his helpful discussion and technical support in XRD measurements. [1] E.W. Kelley, W.F. Hosford, Trans. Metall. Soc. AIME 242 (1968) 5. [2] W.H. Hartt, R.E. Reed-Hill, Trans. Metall. Soc. AIME 239 (1967) 1511. [3] B.C. Wonsiewicz, W.A. Backofen, Trans. Metall. Soc. AIME 239 (1967) 1422. [4] W.H. Hartt, R.E. Reed-Hill, Trans. Metall. Soc. AIME 242 (1968) 1127.
½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0
[5] E.W. Kelley, W.F. Hosford, Trans. Metall. Soc. AIME 242 (1968) 654. [6] R.E. Reedhill, Trans. Metall. Soc. AIME 218 (1960) 554. [7] S. Biswas, S. Suwas, R. Sikand, A.K. Gupta, Mater. Sci. Eng. A 528 (2011) 3722. [8] B. Beausir, S. Suwas, L.S. To´th, K.W. Neale, J.J. Fundenberger, Acta Mater. 56 (2008) 200. [9] B. Beausir, S. Biswas, D.I. Kim, L.S. To´th, S. Suwas, Acta Mater. 57 (2009) 5061. [10] J.Y. Kang, B. Bacroix, R. Brenner, Scripta Mater. 66 (2012) 654. [11] B. Song, G.S. Huang, H.C. Li, L. Zhang, G.J. Huang, F.S. Pan, J. Alloys Compd. 489 (2010) 475. [12] L. Helis, K. Okayasu, H. Fukutomi, Mater. Sci. Eng. A 430 (2006) 98. [13] M.R. Barnett, M.D. Nave, C.J. Bettles, Mater. Sci. Eng. A 386 (2004) 205. [14] M.D. Nave, M.R. Barnett, Scripta Mater. 51 (2004) 881. [15] G. Omori, Trans. Natl. Res. Inst. Met. 21 (1979) 38. [16] L. Jiang, J.J. Jonas, A.A. Luo, A.K. Sachdev, S. Godetet, Scripta Mater. 54 (2006) 771. [17] J. Koike, Metall. Mater. Trans. A 36 (2005) 1689. [18] M.R. Barnett, Mater. Sci. Eng. A 464 (2007) 8. [19] S.L. Couling, J.F. Pashak, L. Sturkey, Trans. ASM 51 (1959) 94. [20] W.B. Hutchinson, B.A. Wilcox, Met. Sci. J. 7 (1973) 6. [21] V. Randle, O. Engler, Introduction to Texture Analysis: Macrotexture, Microtexture and Orientation Mapping, first ed., Gordon & Breach, Amsterdam, 2000. [22] S.B. Yi, C.H.J. Davies, H.G. Brokmeier, R.E. Bolmaro, K.U. Kainer, J. Homeyer, Acta Mater. 54 (2006) 549. [23] M. Knezevic, A. Levinson, R. Harris, R.K. Mishra, R.D. Doherty, S.R. Kalidindi, Acta Mater. 58 (2010) 6230. [24] L. Jiang, J.J. Jonas, R.K. Mishra, A.A. Luo, A.K. Sachdev, S. Godet, Acta Mater. 55 (2007) 3899. [25] Z. Zhang, M.P. Wang, N. Jiang, S.M. Li, Mater. Sci. Eng. A 527 (2010) 6467. [26] P. Cizek, M.R. Barnett, Scripta Mater. 59 (2008) 959. [27] M.R. Barnett, N. Stanford, P. Cizek, A. Beer, Z. Xuebin, Z. Keshavarz, JOM 61 (2009) 19. [28] E.F. Rauch, M. Ve´ron, J. Portillo, D. Bultreys, Y. Maniette, S. Nicolopoulos, Microsc. Anal. 22 (2008) S5. [29] D. Teirlinck, F. Zok, J.D. Embury, M.F. Ashby, Acta Metall. 36 (1988) 1213. [30] I.J. Beyerlein, J. Wang, M.R. Barnett, C.N. Tome´, Proc. R. Soc. A, http://dx.doi.org/10.1098/rspa.2011.0731.
Scripta Materialia 67 (2012) 1015–1018 www.elsevier.com/locate/scriptamat
A critical test of twin-induced softening in a magnesium alloy extruded to a strain of 0.7 at room temperature Zhen Zhang,a,b,⇑ Pavel Cizeka and Matthew Barnetta a
Institute for Frontier Materials, Geelong Technology Precinct, Deakin University, Geelong, Vic 3220, Australia School of Materials Science and Engineering, Central South University, Changsha 410083, People’s Republic of China
b
Received 12 August 2012; revised 19 September 2012; accepted 22 September 2012 Available online 28 September 2012
Ambient extrusion was used to impart different uniaxial strains to a magnesium alloy. The aim was to test the idea that texture change at high strains can lead to work softening. Tensile tests, optical microscopy, X-ray diffraction and transmission electron microscopy were used to investigate the flow stress, microstructure and texture evolution up to a uniaxial strain of 0.7. The microstructure after extrusion is dominated by contraction twinning, which served to weaken the basal texture. However, no softening was observed; global work hardening persisted. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloys; Twinning; Texture; Work hardening
Single crystals of magnesium display very little ductility when compressed along the c-axis [1–6]. Unfortunately, such deformation is common. Polycrystalline magnesium alloys develop a sharp “basal” texture during wrought processing and this texture is characterized by caxis aligned near to perpendicular to the tooling face [7– 12]. This type of deformation activates f1 0 1 1g contraction twinning and subsequent double twinning [3,4,6]. In the twinned volumes, reorientation of the basal planes into the plane of maximum shear stress triggers flow concentration which is followed by failure [13–17]. In addition to this local texture softening effect, a number of workers (e.g. [16]), including one of us [18], has suggested that it may also lead to global softening. Such has been seen by Couling et al. [19] in cold-rolled sheet and the phenomenon may play an important role in tensile ductility. This idea is critically tested in the current work and is shown to be incorrect for the chosen alloy. The intent of the study is to determine the microstructural evolution and flow stress at similar strain levels and strain paths to those at which local necking and failure occur in a tensile test. To achieve this, ambient extrusion was used to impart uniaxial strains up to 0.68. Uniaxial tensile tests were then conducted on these extruded samples. X-ray diffraction (XRD) and
⇑ Corresponding
author at: School of Materials Science and Engineering, Central South University, Changsha 410083, People’s Republic of China.; e-mail: [email protected]
automated crystal orientation mapping in a transmission electron microscope (TEM) were also conducted to characterize twinning and texture evolution. The material used in the present work was commercially available wrought AZ31 (Mg–3% Al–1% Zn) magnesium alloy, received in the form of 30 and 100 mm diameter extruded bars. The material was given two different annealing treatments: 430 °C for 8 h and 300 °C for 30 min to produce grain sizes of 8 and 30 lm, respectively. Conical extrusion dies with the same initial diameter of 17 mm but with different exit diameters of 12, 13, 14, 15 and 16 mm were employed. During ambient extrusion, fracture was evident only in some of the 30 lm samples when the 13 mm die (53% strain) was employed, and our attempts at ambient extrusion using the 12 mm die (70% strain) failed for both grain sizes. In order to introduce higher strains, 10 mm diameter billets were fitted into aluminum tube-like shells prior to extrusion. This method permitted true strains of 0.64 for the 30 lm and 0.68 for the 8 lm material to be achieved without failure using the 12 mm die. Round tensile samples were cut along the extrusion direction with a gauge length of 11.5 mm and a gauge diameter of 2.9 mm. Uniaxial tensile tests were performed on these samples with different extrusion strain for both grain sizes. True stress–strain curves are shown in Figure 1a and b and it can be seen that the flow stress increases with strain. No global softening (drop in yield stress with increase in strain) is evident, unlike that seen by Couling
1359-6462/$ - see front matter Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.scriptamat.2012.09.021
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Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Figure 1. The true stress–strain curves for the (a) 8 lm and (b) 30 lm materials after ambient uniaxial extrusion to different strains. (c) The reduction in area for the 8 and 30 lm materials after different extrusion strain.
et al. [19] for rolling. The Ludwik–Hollomon relationship (r = ken) is fitted to the initial tensile curve and extrapolated to higher strains. This only slightly overestimates the measured behavior. Interestingly, the reduction in area measured after each test showed very little sensitivity to strain (Fig. 1c). There is even some slight increase with strain at the low extrusion reductions. It is as if the extrusion deformation renders the material slightly more ductile. Microstructures were observed for each strain condition, before tensile testing, on the longitudinal section (parallel to the extrusion direction, ED). Both grain size materials display extensive twinning activity (Fig. 2a and b). The twin volume fraction seen at low strains agrees well with previous work [18] and the twins are consistent in appearance with f1 0 1 1g double twins [3,4,6]. However, orientation information for the present twins was difficult to obtain using electron backscatter diffraction (EBSD) in a scanning electron microscope
Figure 2. Optical microstructure of the 30 lm material after ambient uniaxial extrusion to (a) 10% and (b) 25%. (c) The relationship between twinning volume fraction and extrusion strain for the 8 and 30 lm materials.
(SEM), and therefore in the following XRD and TEM orientation measurements are employed for the material with a grain size of 30 lm. Macrotexture measurements were conducted on a Panalytical X-ray diffractometer in Schulz reflection geometry. All the measurements were performed on the plane 45° to ED; in this way, complete pole figure data could be collected [20]. Considering the small 2h angle (around 34.5° for the basal plane) and the limited sample dimensions, a small spot size (0.1 mm 1 mm) was used. Accordingly, the tube voltage and current was set as 40 kV and 50 mA, and the exposure time was 60 s per step. Defocusing and background were corrected and all the measurements were collected under the same conditions. Normalization was performed as follows. Due to the axial symmetry of the fiber texture in the round extrusion samples, the normalization integration was only carried out along one equal-U line from the radius direction (RD) to the extrusion direction (ED) (see Fig. 3). The normalizing factor F [21] could thus be calculated as (w is the angle from RD), based on which the X-ray Rp intensity was divided to 02 I intensity cos wdw acquire normalized random values of x. The (0002) pole density is plotted along one equal-U line (U = 0) in Figure 3 (where the w = 0° position represents RD). Since the X-ray intensity is extremely low at high w angles (w > 40°), we only show the 35° to +35° w range for the U = 0 line. Symmetrization was applied to the data to give a full description of the distribution of the basal intensity around RD. It is clear in the plot that the texture intensity decreases with increasing strain. It is frequently reported [22–25] that the characteristic basal texture strengthens with imposed strain during plastic processing. The weakening of basal texture seen here is most likely a consequence of the high twin fractions formed. TEM analysis was carried out to support this idea. A JEOL 2100F TEM was employed with an accelerating voltage of 200 kV. The specimens were sectioned perpendicular to ED. Figure 4a shows a TEM bright-field image of a single twin formed in a matrix with the typical h1 0 1 0i==ED orientation. The twinning area is broken up into several individual blocks with different orientations. Similar orientations are outlined in the same color in the figure. The orientations of the different blocks within the twinning area share a common h1 1 2 0i pole with the matrix.
Figure 3. The distribution of (0002) pole density along the U = 0 line.
Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Figure 4. (a) TEM image of a single twin; the zone axis ½2 1 1 0 is 30° from the ED. (b) Inverse pole figure map of twinning band area. Both single twin and twin bands were formed in the matrix, which had typical ½1 0 1 0==ED orientation.
Selected-area diffraction patterns were taken separately from all these areas along their common h1 1 2 0i zone axis. Block D has a misorientation of 42° clockwise along ½2 1 1 0 from the matrix, and block A has a misorientation of 36.2° anticlockwise along ½2 1 1 0 from the matrix. Both areas have a close orientation to the ideal f1 0 1 g–f1 0 1 2g double twin, which is positioned at 38° h1 1 2 0i to the matrix. If block A is considered as a f1 0 1 1g–f1 0 1 2g double twinning area, the first twinning event would have occurred on the ð0 1 1 1Þ plane. Alternatively, if block D is considered as a f1 0 1 1g–f1 0 1 2g double twinning area, the first twinning event would have occurred on the ð0 1 1 1Þ plane. From the diffraction pattern, both ð0 1 1 1Þ and ð0 1 1 1Þ plane traces were established and these are shown in Figure 4a. The expected trace for block D agrees with the observed habit of the primary twin, so block D is identified as the initial double twinning area. Blocks A and B were found to share the typical f1 0 1 2g twinning relationship (86° h1 1 2 0i). Since the boundary between block A and block B coincided well with the ð1 0 1 2Þ plane trace in block D, block A could be considered as a triple twinning region within block D. The resulting orientations in the twinned volume are consistent with the spread of the texture seen in the X-ray data. It is evident that the twinning processes taking place in the present material are quite complex and multiple twinning events, with a
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common h1 1 2 0i axis, can form in one volume. Similar complex contraction twins were previously observed by some of the present authors close to the fracture surface of AZ31 tensile specimens [26,27]. To gain a better view of the twinning reorieintations, automatic crystal orientation measurements were performed using the Nanomegas Astar system. Selected areas were scanned in nanobeam mode using a step size of 3 nm. Diffraction patterns were collected and compared with simulated templates to find the best matching solution [28]. An inverse pole figure (IPF) map of a typical twinned band is shown in Figure 4b. High-angle boundaries, including twin boundaries, are highlighted in black. Low-angle boundaries, <10°, are highlighted in white. The blue area, whose orientation lies close to the h1 0 1 0i corner in the ED IPF, corresponds to the matrix of the original grain. No particular orientation relationship is evident between these twinning bands and matrix, but all orientations share a h1 1 2 0i pole with the matrix. The common h1 1 2 0i axis is both perpendicular to the extrusion direction and to the plane of shear if orientations A–F are indeed twins. (The twinning plane of shear contains the twin plane normal and the twin shear direction. For the present material, the normal to the plane of shear is the direction about which the twin crystal misorientations are typically described.) The twin band in Figure 4b displays a distinct substructure but the matrix does not. The twinning area is most likely a region of some degree of localized slip. Such is expected from the basal Schmid factors in Table 1, which are 10 times higher in the twinned area than in the adjoining matrix. This slip activity is likely to have rotated the band materials away from its original orientation. Inspection of the f1 0 1 1g pole figure reveals that the habit of the twin band is consistent with two f1 0 1 1g planes, III and IV, in the pole figure in Figure 4b. Orientation E is within 8° of the primary f1 0 1 1g twin expected for twinning on the IV ð1 0 1 1Þ plane. Orientation A is within 6° of the primary f1 0 1 1g twin expected for twinning on the III ð1 0 1 1Þ plane. It is thus apparent that Figure 4b portrays a region close to an intersection of two f1 0 1 1g twin systems. According to this view, orientations A and C are part of one family, and orientations B, D, E and F are part of another family. However, we acknowledge that other interpretations are possible. Nevertheless, the result confirms the presence of complex multiple twinning events and their tendency to weaken the texture. Although the present work confirms the weakening of the basal texture, no global softening was observed. The flow stresses at high strains were slightly lower than extrapolated predictions, which might point to an effect of texture. However, it is clear that work hardening dominates. This is manifested by the substructure seen in the twin bands and by the fragmentation resulting from multiple twinning events in the twinned volumes. Both of these serve to reduce the slip length. Although not observed here, voids have been seen to form in twin volumes in previous studies [18]. Such void formation can be expected to lead to instability of flow and to rupture during a tensile test [3,29]. In the present
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Z. Zhang et al. / Scripta Materialia 67 (2012) 1015–1018
Table 1. Euler angles and basal slip Schmid factor for matrix and different twinning bands. Area
Euler angle (Bunge)
Misorientation from matrix
Schmid factor for basal slip
Matrix A B C D E F
(203, 87.6, 60) (205, 137.4, 241.2) (200, 60.3, 63.6) (203, 117.4, 241.7) (197, 49, 65.4) (24, 140.2, 246.4) (199, 47.3, 60)
– +49.8 27.6 +29.8 39.1 48.2 40.5
0.04 0.44 0.38 0.36 0.45 0.45 0.43
case, the hydrostatic compressive stress component of extrusion prevented such softening from occurring. Indeed, it may be that the fragmentation in the twinned volumes permitted by these circumstances serves to slow the development of twin-induced voiding in subsequent tensile tests. The reduction in area seen in the present tensile samples was preserved after ambient extrusion, despite the higher flow stress levels. Finally, it remains to suggest that the common h1 1 2 0i pole seen in the multiple twinned volumes may relate to the role of basal dislocations in twin nucleation. It has been shown that incoming basal dislocations can lead to secondary twin nucleation at pre-existing twin interfaces [30]. In such a case the nucleating twin is more likely to share with the existing twin the h1 1 2 0i pole that lies in the twin interface. This is the normal to the “plane of shear” identified above. In conclusion, the present study has successfully employed ambient extrusion to attain uniaxial strains as high as 68%. The structures thus produced display copious complex twinning events, which serve to subtly weaken the texture. Global work hardening persists, unlike the behavior observed in cold-rolled magnesium alloy [19]. It is likely that the contribution of contraction twinning to tensile ductility of extruded material lies in the formation of voids in doubly (or tertiary) twinned volumes rather than in global texture softening. The multiply twinned volumes are frequently seen to share a common h1 1 2 0i pole, due possibly to the role of basal dislocation–twin boundary interactions in twin nucleation. The authors acknowledge Professor Bevis Hutchison for his helpful discussion and technical support in XRD measurements. [1] E.W. Kelley, W.F. Hosford, Trans. Metall. Soc. AIME 242 (1968) 5. [2] W.H. Hartt, R.E. Reed-Hill, Trans. Metall. Soc. AIME 239 (1967) 1511. [3] B.C. Wonsiewicz, W.A. Backofen, Trans. Metall. Soc. AIME 239 (1967) 1422. [4] W.H. Hartt, R.E. Reed-Hill, Trans. Metall. Soc. AIME 242 (1968) 1127.
½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0 ½1 2 1 0
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