- Email: [email protected]
Strain rate dependence of twinning at 450 °C and its effect on microstructure of an extruded magnesium alloy
Materials Science & Engineering A 559 (2013) 314–318
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Strain rate dependence of twinning at 450 1C and its effect on microstructure of an extruded magnesium alloy Q. Ma a,n, B. Li a, A.L. Oppedal a, W.R. Whittington a, S.J. Horstemeyer a, E.B. Marin a, P.T. Wang a, M.F. Horstemeyer a,b a b
Center for Advanced Vehicular Systems, Mississippi State University, Starkville, MS 39759, USA Department of Mechanical Engineering, Mississippi State University, Starkville, MS 39762, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 February 2012 Received in revised form 9 July 2012 Accepted 22 August 2012 Available online 30 August 2012
Deformation twinning in magnesium alloys at elevated temperatures has received relatively little attention because it is generally deemed that dislocation slip dominates plastic deformation. In this work, twinning at 450 1C in an extruded Mg–Al–Mn magnesium alloy (AM30) was studied by interrupted compression tests at various strain rates within a practical range for lab-scale extrusion ( o 1.0 s 1). Microstructure and texture evolution were examined by electron backscatter diffraction (EBSD) at different strain levels. The results show that sporadic twins started to appear at strain rate of 0.1 s 1, whereas profuse twinning was activated at strain rates of 0.5 and 0.8 s 1. The deformation twins quickly lost original morphology because of dynamic recrystallization. These results show that deformation twinning has a significant effect on microstructural and texture evolution of wrought Mg alloys at elevated temperatures within practical strain rate range. & 2012 Elsevier B.V. All rights reserved.
Keywords: Twinning Magnesium Texture Dynamic recrystallization
1. Introduction Due to their stiffness-to-weight ratios, hexagonal close-packed (hcp) magnesium (Mg) alloys are of particular interest for improved fuel efficiency in automotive and aerospace applications. In sharp contrast to dislocation-controlled deformation in high symmetry crystal structures, twinning plays a dominant role in plastic deformation of Mg alloys, because these low symmetry hcp crystal structures have a limited number of easy slip systems [1,2]. The f1012g o 1011 4 twin is the most commonly observed twinning mode p inffiffiffihcp metals, irrespective of the c/a ratio. For c/a ratios less than 3 this twinning mode creates an extension along the c-axis of the crystal and hence can accommodate a tensile strain [2]. In Mg, another twinning mode, i.e., f1011g o 1012 4 , can accommodate compression strain along the c-axis and plays an important role at strain levels near fracture [3,4]. Double twinning f1011g þ f1012g was also reported in Mg and Mg alloys [5–7]. Generally, twinning is favored at low temperatures and/or high strain rates [1,2,8]. f1012g o1011 4 twins are rarely observed in Mg at temperatures higher than 350 1C and strain rates less than 0.01 s 1 [9–11]. However, there are experimental results indicative of twinning at temperatures above 350 1C. Through transmission electron microscopy (TEM), Myshlyaev
n
Corresponding author. Tel.: þ1 662 325 5457; fax: þ 1 662 325 5433. E-mail address: [email protected] (Q. Ma).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.104
et al. [9] identified occasional twin-like substructures in an AZ31 alloy deformed at 450 1C. Lou et al. [12] observed twin bands in AZ80 at 450 1C by TEM and at a relatively high strain rate of 10 s 1. Using optical microscopy, Ishikawa et al. [13] uncovered twins in AZ91 at 400 1C and at a high strain rate of e_ ¼103 s 1. Liu and Ding [14] observed microtwin bands in AZ91 at 400 1C under TEM, but they could not find twins via optical microscopy between 350 and 450 1C, regardless of the strain rate. Despite these previous reports [9–11,15–20], EBSD characterization of deformation twinning in Mg at elevated temperatures in the range of strain rates for lab-scale extrusion of Mg ( o1.0 s 1) is not well documented in the literature. In practice, the low workability of Mg alloys requires thermomechanical processing at elevated temperatures; for example, extrusion is normally performed above 350 1C [21]. Strain rates, however, may vary and depend upon the geometry of the work piece and tooling fixture as well as the speed of the process (e.g., ram speed in extrusion). Even in a single work piece that is undergoing thermomechanical processing, the strain rate can be inhomogeneous from one location to another. Strain rate exerts important influence on deformation mechanism in metals [8], thus the change in strain rate can lead to mechanistic complications in deformation of Mg alloys during thermomechanical processing, i.e., the activation of deformation modes in both slip and twinning can be different. The purpose of this work is to characterize with clarity deformation twinning of an Mg alloy (AM30) at 450 1C and
Q. Ma et al. / Materials Science & Engineering A 559 (2013) 314–318
various strain rates within the practical range and, in turn, the effect of both twinning and dynamic crystallization (DRX) on the alloy’s microstructure.
2. Experimental According to Atwell and Barnett [21], the average strain rate in extrusion can be expressed as e_ 9:6R0:6 V R =DB , where e_ is the extrusion average strain rate, R the extrusion ratio, VR the ram speed, and DB the billet diameter. In our lab-scale Mg extrusion practice, the extrusion ratio is usually less than 25, DB is 31.5 mm and the ram speed may vary between 0–30 mm/min [22]. Hence, the average strain rate for extrusion falls in the range less than 1.0 s 1. Accordingly, we selected for our experimental purposes the following strain rates: 0.001 s 1, 0.1 s 1, 0.5 s 1 and 0.8 s 1. An extruded commercial AM30 magnesium alloy (mass: Al 2.54%, Mn 0.40%, Mg balance) was selected as the experimental material because this material has been used for validation of extrusion modeling in our laboratory. Cylindrical samples 10 mm in diameter and 10 mm in height were cut by electrical-discharge machining (EDM), with the axial direction parallel to the extrusion direction (ED). Before uniaxial compression, the sample was heated to 450 1C in the furnace that houses the sample stage of INSTRON 5869 testing machine, and the temperature was maintained at 450 1C during testing. Then the samples were compressed inside the furnace at strain rates of 0.001 s 1, 0.1 s 1, 0.5 s 1 or 0.8 s 1, respectively. For a given strain rate, a series of samples were compressed to different strain levels that are indicated in Fig. 1, and the test was stopped and followed by immediate water quench of the sample to preserve the microstructure. Uninterrupted tests at each strain rate were also performed to collect overall stress–strain responses shown in Fig. 1. The deformed samples were cold mounted and then subsequently ground with a series of SiC paper (down to a grit number 4000), and then polished using alumina (down to 0.05 mm) suspended in ethylene glycol. The polished samples were then etched using a solution of 10 ml HNO3, 30 ml acetic acid, 40 ml H2O and 120 ml ethanol for 5 s. EBSD scans were performed on the etched samples with step sizes of 0.5–1 mm in a field emission gun scan electron microscopy (FEG-EBSD) equipped with TSL OIM data collection and analysis software.
Fig. 1. The uniaxial compression true stress–strain curves of AM30 at 450 1C and strain rates: 0.001 s 1, 0.1 s 1, 0.5 s 1 and 0.8 s 1. The curves were obtained from uninterrupted tests. The arrows denote the EBSD scans of interrupted tests at different strain levels and strain rates.
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3. Results and discussion The true stress–strain curves of the AM30 at 450 1C and at different strain rates are presented in Fig. 1. The yield stress and the flow stress increase with increasing strain rate. At strain rates of 0.001 s 1 and 0.1 s 1, the strain hardening is minimal. The stress–strain curve of 0.5 s 1 shows barely recognizable hardening, followed by a slight softening. When the strain rate increases to 0.8 s 1, however, the stress–strain behavior exhibits an apparent sigmoidal shape of appreciable hardening, with a peak stress followed by a pronounced softening. The stress–strain curves at 0.001 s 1 and 0.1 s 1 present a parabolic shape in which dislocation slip dominates the plastic flow [23,24]. The texture evolution during compression in the ED at various strain rates is shown in Fig. 2; the inverse pole figure (IPF) maps are shown as well. Fig. 2a displays the initial microtexture and microstructure of the untested sample. The initial microstructure consists of a mixture of coarse grains and dynamically recrystallized fine grains created during extrusion. In the pole figure, the basal planes are roughly parallel to the ED (or o1010 4 O ED), which is typical of extruded Mg alloys. In Fig. 2b, homogenous, equiaxed grains at 0.001 s 1 and e ¼ 0.15 indicates a full recrystallization, i.e., nucleation and ripening in comparison to the initial state (Fig. 2a). Fig. 2c presents the microstructure at 0.1 s 1 and e ¼ 0.15 in which lenticular twins (in red) start emerging. Twinning is confirmed by the 901 rotation in the {0001} pole figure, which is caused by the presence of f1012g o1011 4 twins. In general, at a given strain, the twin volume fraction increases as the strain rate increases, and this can be seen from the basal texture intensity in the {0001} pole figure in Fig. 2. But the intensity at 0.5 s 1 and e ¼ 0.04 is slightly higher than that of 0.8 s 1 and e ¼ 0.04. This interesting phenomenon is worth further investigations. When the strain rate increases to 0.5 s 1, as shown in Fig. 2d, profuse, large f1012g o1011 4 twins can be observed in a coarse, nonrecrystallized grain in addition to the small twins in the fine grains. A strong basal texture can be observed in the pole figure. In the meantime, dynamically recrystallized, ‘‘necklace’’ grains nucleate and grow, surrounding the parent grains. Some new grains invade into the matrix along the twin boundaries. Fig. 2e–g show the microstructure at 0.8 s 1, e ¼ 0.04, 0.15, and 0.60, respectively. The f1012g o 1011 4 twinning dominates the deformation and gives rise to an enhanced basal texture, and then the basal texture weakens due to dynamic recrystallization. At e ¼ 0.04, high density twins can be observed. Meanwhile, fine ‘‘necklace’’ grains surrounding parent grains nucleate and grow. At e ¼ 0.15, the twins grow significantly, coalesce and consume most of the original grains. The dynamically recrystallized grains are sandwiched between the impinging twins. At e ¼ 0.60, new grains also nucleate inside the twins and grow, eventually segmenting the large twins into irregular islands. No f1011g o 1012 4 contraction twins or f1011g þ f1012g double twins were found in this study. Grain size distributions under different processing conditions are presented in Fig. 3. The as-received, extruded sample presents a bimodal structure containing large parent grains and small, recrystallized grains, as seen in the double-peak distribution. At the quasi-static strain rate of 0.001 s 1 and e ¼ 0.15, only a single peak is observed because grain growth dominated the DRX. At 0.1 s 1 and e ¼ 0.15, we also see a single peak in the grain size distribution, but the grain size is finer than that of 0.001 s 1. In contrast, for the high strain rate of 0.8 s 1, at strains of e ¼ 0.04 and e ¼ 0.15, large grains and fine, dynamically recrystallized grains coexist due to twinning and DRX. As the strain was increased to e ¼ 0.60, grain growth and DRX inside the twins and at the twin boundaries changed the microstructure structure and the high peak corresponding to the large grains
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Fig. 2. EBSD inverse pole figure map and micro-texture of AM30 under uniaxial compression at 450 1C at (a) initial status; (b) 0.001 s 1, e ¼ 0.15, grains coarsen and become equiaxed; (c) 0.1 s 1, e ¼ 0.15, visible thin f1012g o 1011 4 twins appear; (d) 0.5 s 1, e ¼ 0.04, thick twins locate in parent grain and ‘‘necklace’’ grains appear, some of the latter grow along twin interfaces; (e) 0.8 s 1, e ¼ 0.04, twins propagate in both parent grains and recrystallized grains; (f) 0.8 s 1, e ¼ 0.15, twins sweep almost all parent grains and result in a very strong basal texture, and DRX mainly occurs at grain boundaries; and (g) 0.8 s 1, e ¼ 0.60, DRX grains grow at the expense of twinned grains, resulting in irregular twin remains. The pole figures are calculated based on the corresponding EBSD maps. The inverse pole figure represents the compression direction, namely the extrusion direction (ED). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Q. Ma et al. / Materials Science & Engineering A 559 (2013) 314–318
disappears. It should be noted that DRX and grain growth are time dependent. At different strain rates, the amounts of time that the samples were exposed to the high temperature varied from tens of seconds to a fraction of a second, contributing to the variations in grain size distribution and in the microstructure. Our EBSD analysis (Fig. 2) reveals the following: (1) twinning occurs when the strain rate reaches and exceeds 0.1 s 1 at 450 1C; (2) twinning strongly depends upon strain rate at 450 1C, and twin volume fraction increases rapidly as the strain rate exceeds a threshold value; (3) twinning and DRX interact and influence the ultimate microstructure. The strain rate and temperature influence the critical resolved shear stress (CRSS) of non-basal slip markedly, however their influence on CRSS of twinning is comparatively minimal [1,24]. At the quasi-static strain rate of 0.001 s 1, the dominant deformation mechanism is dislocation slip, including non-basal slip as the CRSS of the non-basal slip decreases as the temperature increases [25]. As the strain rate increases, twinning starts to play a marked role in the deformation. In general, dislocation slip is both temperature and strain rate dependent. An increase in the temperature leads to a decrease in the CRSS of slip, thus
Fig. 3. Grain size distribution of AM30 at various strain rates and strain levels compressed at 450 1C. The initial grain size shows a bimodal structure (double peaks) containing large parent grains and small, recrystallized grains. At 10 3 s 1 and e ¼ 0.15, the microstructure was fully recrystallized. At 0.8 s 1, e ¼ 0.04 and e ¼ 0.15, large parent grains and dynamically recrystallized grain coexists, whereas at e ¼ 0.60, the twinned parent grains were also dynamically recrystallized and the corresponding high peak disappears.
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promoting dislocation activity; however, an increase in strain rate strongly increases the CRSS of slips, as shown in Fig. 1, where the flow stress significantly increases with strain rate. This is consistent with hardening models proposed by Mecking and Kocks [26], Beyerlein and Tome´ [27], and Oppedal et al. [28]. Using a full-constraint Taylor model, Barnett [29] calculated the CRSS for basal, prismatic, second-order pyramidal slip, and f1012g twinning. It was suggested that the CRSS for basal slip and f1012g twinning is relatively temperature independent. Ulacia et al. [30] did extensive studies on the mechanical behavior of an AZ31 Mg alloy over a wide range of strain rates and temperatures, and they confirmed that the CRSS for f1012g twinning remained relatively constant with temperature up to 400 1C. Based on these previous studies in the literature, it could be generally concluded that the CRSS of basal slip and twinning are less sensitive to the temperature and strain rate in Mg alloys [8,24], while the CRSS of nonbasal slips are strongly sensitive to temperature and strain rate. Although Barnett’s calculation [29] shows that the CRSS for basal slip is about 2.5 MPa, and the CRSS for f1012g twinning is about 32 MPa in an AZ31 Mg alloy, the most recent measurement [31] of the CRSS for f1012g twinning in pure single crystal Mg is only 2.4 MPa, close to that of basal slip. These calculations and measurements indicate that the CRSS of f1012g twinning can actually be very low. Therefore, as strain rate increases, a higher CRSS for dislocation slip promotes f1012g twinning, as shown in the present experimental work. The mechanical testing (Fig. 1) reflects the competition between mechanical twinning and DRX, as shown in previous studies [9,10,12–15,17–19]. Our EBSD analysis of samples examined after interrupted mechanical tests further clarifies the significance of twinning in the moderate strain rates. As Fig. 2c depicts, twinning does occur concomitantly with DRX when the strain rate reaches 0.1 s 1. But this concurrence of twinning and DRX is not reflected in the stress–strain curve (Fig. 1), because DRX dominates the microstructure evolution. Although twinning is occurring as observed in Fig. 2, its contribution to hardening is overwhelmed by the softening from DRX so that the stress–strain curve remains parabolic in contrast to the cases of the higher strain rates (0.5 s 1 and 0.8 s 1). This indicates that the ‘‘sigmoidal’’ hardening is indeed a signature of twinning, but twinning may have already occurred when such ‘‘sigmoidal’’ hardening is not present in stress–strain curves. Hence, twinning starts to occur at a lower strain rate than previously thought [9–11]. When the strain rate increases to 0.5 s 1 and 0.8 s 1, twinning with a higher volume fraction corresponding to the increased basal texture intensity in the {0001} pole figure in Fig. 2 is activated and this generates an observable effect in strain hardening.
Fig. 4. Schematic illustration of twinning behavior and microstructure evolution in the AM30 Mg alloy compression at 450 1C and various strain rates: (a) low strain rate 0.001 s 1, the microstructure evolution is controlled by dynamic recrystallization (DRX) and grain growth; (b) intermediate strain rate 0.1 s 1, sporadic twinning occurs and twins are completely consumed by DRX; (c) high strain rate 0.8 s 1, profuse twinning is activated; DRX occurs inside the twins and at the twin boundaries; the twins are then invaded by new grains, lose their original morphology, and become unrecognizable in optical microscopy.
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Our EBSD analysis indicates that the interaction of DRX and twinning strongly influences the microstructure and texture at different strain levels. When twinning becomes more pronounced as the strain rate increases, twins with relatively coherent twin boundaries can be observed in optical microscopy (Fig. 2d and e) at low strain levels. As the plastic flow proceeds, however, the twins are invaded by the DRX grains and therefore are no longer recognizable in optical microscopy; hence, EBSD is preferred for microstructure analysis. A few DRX grains nucleate inside the twins (Fig. 2f), because the twins orientations favor pyramidal slip, which facilitates DRX inside the twins [18]. Meanwhile, the twin boundaries are the preferred nucleation sites for DRX grains. As a result, the growth of the DRX grains consumes parts of or entire twins, leading to highly irregular, incoherent twin boundaries. Consequently, caution should be exercised when optical microscopy is used to examine the microstructure of samples deformed at elevated temperatures. The effect of strain rate on twinning and microstructure of the AM30 compressed at 450 1C is schematically illustrated in Fig. 4. At quasi-static strain rates, slip dominates in deformation. Nucleation and grain growth by DRX determine the microstructure and texture evolution. Due to the low strain rate, the recrystallized grains quickly ripen, as shown in Fig. 4a. At a moderate, intermediate strain rate (0.1 s 1), twinning is activated as the CRSS for dislocation slip increases, but the twinning competes with DRX and has little contribution to hardening (Fig. 4b). At high strain rates (40.5 s 1), profuse twinning dominates the plastic deformation as fine grains nucleate along parent grain boundaries; concomitantly, DRX occurs at the twin boundaries and inside the twins (Fig. 4c). As deformation continues at the higher strain rates, the twins quickly lose their original morphology because of DRX, but twinning has an important effect on the texture.
4. Conclusions Using EBSD, we investigated the effect of strain rate ( o1.0 s 1) on twinning and microstructure at 450 1C in an extruded AM30 alloy. At quasi-static strain rates, DRX and grain growth dominate the microstructure and texture evolution. Although sporadic twins begin to appear at strain rate 0.1 s 1, their contribution to hardening is insignificant and no sigmoidal hardening was observed. DRX plays a dominant role in the softening. As the strain rate increases to 0.5 s 1, the twin volume fraction increases significantly and makes an observable
contribution to the hardening. At strain rate of 0.8 s 1, profuse twinning is activated and a peak stress followed by softening due to DRX is observed. The twins quickly lose their original morphology because of DRX.
Acknowledgments The authors are grateful to the financial support from the Department of Energy, Contract no. DE-FC-26-06NT42755, and the Center for Advanced Vehicular Systems (CAVS) at Mississippi State University. Thanks are extended to H.R. Wainwright for her assistance with the preparation of the manuscript. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
J.W. Christian, S. Mahajan, Prog. Mater. Sci. 39 (1995) 1–157. M.H. Yoo, Metall. Trans. A 12 (1981) 409–418. M.R. Barnett, Mater. Sci. Eng. A 464 (2007) 8–16. Q. Ma, H. El Kadiri, A.L. Oppedal, J.C. Baird, B. Li, M.F. Horstemeyer, S.C. Vogel, Int. J. Plasticity 29 (2012) 60–76. R.E. Reed-Hill, Trans. Metall. Soc. AIME 218 (1960) 554–558. A.G. Crocker, Philos. Mag. 7 (1962) 1901–1924. Q. Ma, H. El Kadiri, A.L. Oppedal, J.C. Baird, M.F. Horstemeyer, M. Cherkaoui, Scripta Mater. 64 (2011) 813–816. ¨ M.A. Meyers, O. Vohringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025–4039. M.M. Myshlyaev, H.J. McQueen, A. Mwembela, E. Konopleva, Mater. Sci. Eng. A 337 (2002) 121–133. T. Al-Samman, G. Gottstein, Mater. Sci. Eng. A 490 (2008) 411–420. A.G. Beer, M.R. Barnett, Metall. Mater. Trans. A 38A (2007) 1856–1867. Y. Lou, L. Li, J. Zhou, L. Na, Mater. Charact. 62 (2011) 346–353. K. Ishikawa, H. Watanabe, T. Mukai, Mater. Lett. 59 (2005) 1511–1515. L. Liu, H. Ding, J. Alloys Compd. 484 (2009) 949–956. O. Sitdikov, R. Kaibyshev, Mater. Trans. 42 (2001) 1928–1937. R. Kaibyshev, O. Sitdikov, Phys. Met. Metall. 89 (2000) 70–77. S.E. Ion, F.J. Humphreys, S.H. White, Acta Metall. 30 (1982) 1909–1919. A. Galiyev, R. Kaibyshev, G. Gottstein, Acta Mater. 49 (2001) 1199–1207. M.R. Barnett, J. Light Met. 1 (2001) 167–177. M. Wang, R. Xin, B. Wang, Q. Liu, Mater. Sci. Eng. A 528 (2011) 2941–2951. D.L. Atwell, M.R. Barnett, Metall. Mater. Trans. A 38A (2007) 3032–3041. Q. Ma, B. Li, E.B. Marin, S.J. Horstemeyer, Scripta Mater. 65 (2011) 823–826. M.R. Barnett, C.H.J. Davies, X. Ma, Scripta Mater. 52 (2005) 627–632. A. Jain, S.R. Agnew, Mater. Sci. Eng. A 462 (2007) 29–36. A. Chapuis, J.H. Driver, Acta Mater. 59 (2011) 1986–1994. H. Mecking, U.F. Kocks, Acta Metall. 29 (1981) 1865–1875. I.J. Beyerlein, C.N. Tome´, Int. J. Plasticity 24 (2008) 867–895. A.L. Oppedal, H. El Kadiri, C.N. Tome´, G.C. Kaschner, S.C. Vogel, J.C. Baird, M.F. Horstemeyer, Int. J. Plasticity 30–31 (2012) 41–61. M.R. Barnett, Metall. Mater. Trans. A 34A (2003) 1799–1806. I. Ulacia, N.V. Dudamell, F. Ga´lvez, S. Yi, M.T. Pe´rez-Prado, I. Hurtado, Acta Mater. 58 (2010) 2988–2998. Q. Yu, J. Zhang, Y. Jiang, Phil. Mag. Lett. 91 (2011) 757–765.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Strain rate dependence of twinning at 450 1C and its effect on microstructure of an extruded magnesium alloy Q. Ma a,n, B. Li a, A.L. Oppedal a, W.R. Whittington a, S.J. Horstemeyer a, E.B. Marin a, P.T. Wang a, M.F. Horstemeyer a,b a b
Center for Advanced Vehicular Systems, Mississippi State University, Starkville, MS 39759, USA Department of Mechanical Engineering, Mississippi State University, Starkville, MS 39762, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 February 2012 Received in revised form 9 July 2012 Accepted 22 August 2012 Available online 30 August 2012
Deformation twinning in magnesium alloys at elevated temperatures has received relatively little attention because it is generally deemed that dislocation slip dominates plastic deformation. In this work, twinning at 450 1C in an extruded Mg–Al–Mn magnesium alloy (AM30) was studied by interrupted compression tests at various strain rates within a practical range for lab-scale extrusion ( o 1.0 s 1). Microstructure and texture evolution were examined by electron backscatter diffraction (EBSD) at different strain levels. The results show that sporadic twins started to appear at strain rate of 0.1 s 1, whereas profuse twinning was activated at strain rates of 0.5 and 0.8 s 1. The deformation twins quickly lost original morphology because of dynamic recrystallization. These results show that deformation twinning has a significant effect on microstructural and texture evolution of wrought Mg alloys at elevated temperatures within practical strain rate range. & 2012 Elsevier B.V. All rights reserved.
Keywords: Twinning Magnesium Texture Dynamic recrystallization
1. Introduction Due to their stiffness-to-weight ratios, hexagonal close-packed (hcp) magnesium (Mg) alloys are of particular interest for improved fuel efficiency in automotive and aerospace applications. In sharp contrast to dislocation-controlled deformation in high symmetry crystal structures, twinning plays a dominant role in plastic deformation of Mg alloys, because these low symmetry hcp crystal structures have a limited number of easy slip systems [1,2]. The f1012g o 1011 4 twin is the most commonly observed twinning mode p inffiffiffihcp metals, irrespective of the c/a ratio. For c/a ratios less than 3 this twinning mode creates an extension along the c-axis of the crystal and hence can accommodate a tensile strain [2]. In Mg, another twinning mode, i.e., f1011g o 1012 4 , can accommodate compression strain along the c-axis and plays an important role at strain levels near fracture [3,4]. Double twinning f1011g þ f1012g was also reported in Mg and Mg alloys [5–7]. Generally, twinning is favored at low temperatures and/or high strain rates [1,2,8]. f1012g o1011 4 twins are rarely observed in Mg at temperatures higher than 350 1C and strain rates less than 0.01 s 1 [9–11]. However, there are experimental results indicative of twinning at temperatures above 350 1C. Through transmission electron microscopy (TEM), Myshlyaev
n
Corresponding author. Tel.: þ1 662 325 5457; fax: þ 1 662 325 5433. E-mail address: [email protected] (Q. Ma).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.104
et al. [9] identified occasional twin-like substructures in an AZ31 alloy deformed at 450 1C. Lou et al. [12] observed twin bands in AZ80 at 450 1C by TEM and at a relatively high strain rate of 10 s 1. Using optical microscopy, Ishikawa et al. [13] uncovered twins in AZ91 at 400 1C and at a high strain rate of e_ ¼103 s 1. Liu and Ding [14] observed microtwin bands in AZ91 at 400 1C under TEM, but they could not find twins via optical microscopy between 350 and 450 1C, regardless of the strain rate. Despite these previous reports [9–11,15–20], EBSD characterization of deformation twinning in Mg at elevated temperatures in the range of strain rates for lab-scale extrusion of Mg ( o1.0 s 1) is not well documented in the literature. In practice, the low workability of Mg alloys requires thermomechanical processing at elevated temperatures; for example, extrusion is normally performed above 350 1C [21]. Strain rates, however, may vary and depend upon the geometry of the work piece and tooling fixture as well as the speed of the process (e.g., ram speed in extrusion). Even in a single work piece that is undergoing thermomechanical processing, the strain rate can be inhomogeneous from one location to another. Strain rate exerts important influence on deformation mechanism in metals [8], thus the change in strain rate can lead to mechanistic complications in deformation of Mg alloys during thermomechanical processing, i.e., the activation of deformation modes in both slip and twinning can be different. The purpose of this work is to characterize with clarity deformation twinning of an Mg alloy (AM30) at 450 1C and
Q. Ma et al. / Materials Science & Engineering A 559 (2013) 314–318
various strain rates within the practical range and, in turn, the effect of both twinning and dynamic crystallization (DRX) on the alloy’s microstructure.
2. Experimental According to Atwell and Barnett [21], the average strain rate in extrusion can be expressed as e_ 9:6R0:6 V R =DB , where e_ is the extrusion average strain rate, R the extrusion ratio, VR the ram speed, and DB the billet diameter. In our lab-scale Mg extrusion practice, the extrusion ratio is usually less than 25, DB is 31.5 mm and the ram speed may vary between 0–30 mm/min [22]. Hence, the average strain rate for extrusion falls in the range less than 1.0 s 1. Accordingly, we selected for our experimental purposes the following strain rates: 0.001 s 1, 0.1 s 1, 0.5 s 1 and 0.8 s 1. An extruded commercial AM30 magnesium alloy (mass: Al 2.54%, Mn 0.40%, Mg balance) was selected as the experimental material because this material has been used for validation of extrusion modeling in our laboratory. Cylindrical samples 10 mm in diameter and 10 mm in height were cut by electrical-discharge machining (EDM), with the axial direction parallel to the extrusion direction (ED). Before uniaxial compression, the sample was heated to 450 1C in the furnace that houses the sample stage of INSTRON 5869 testing machine, and the temperature was maintained at 450 1C during testing. Then the samples were compressed inside the furnace at strain rates of 0.001 s 1, 0.1 s 1, 0.5 s 1 or 0.8 s 1, respectively. For a given strain rate, a series of samples were compressed to different strain levels that are indicated in Fig. 1, and the test was stopped and followed by immediate water quench of the sample to preserve the microstructure. Uninterrupted tests at each strain rate were also performed to collect overall stress–strain responses shown in Fig. 1. The deformed samples were cold mounted and then subsequently ground with a series of SiC paper (down to a grit number 4000), and then polished using alumina (down to 0.05 mm) suspended in ethylene glycol. The polished samples were then etched using a solution of 10 ml HNO3, 30 ml acetic acid, 40 ml H2O and 120 ml ethanol for 5 s. EBSD scans were performed on the etched samples with step sizes of 0.5–1 mm in a field emission gun scan electron microscopy (FEG-EBSD) equipped with TSL OIM data collection and analysis software.
Fig. 1. The uniaxial compression true stress–strain curves of AM30 at 450 1C and strain rates: 0.001 s 1, 0.1 s 1, 0.5 s 1 and 0.8 s 1. The curves were obtained from uninterrupted tests. The arrows denote the EBSD scans of interrupted tests at different strain levels and strain rates.
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3. Results and discussion The true stress–strain curves of the AM30 at 450 1C and at different strain rates are presented in Fig. 1. The yield stress and the flow stress increase with increasing strain rate. At strain rates of 0.001 s 1 and 0.1 s 1, the strain hardening is minimal. The stress–strain curve of 0.5 s 1 shows barely recognizable hardening, followed by a slight softening. When the strain rate increases to 0.8 s 1, however, the stress–strain behavior exhibits an apparent sigmoidal shape of appreciable hardening, with a peak stress followed by a pronounced softening. The stress–strain curves at 0.001 s 1 and 0.1 s 1 present a parabolic shape in which dislocation slip dominates the plastic flow [23,24]. The texture evolution during compression in the ED at various strain rates is shown in Fig. 2; the inverse pole figure (IPF) maps are shown as well. Fig. 2a displays the initial microtexture and microstructure of the untested sample. The initial microstructure consists of a mixture of coarse grains and dynamically recrystallized fine grains created during extrusion. In the pole figure, the basal planes are roughly parallel to the ED (or o1010 4 O ED), which is typical of extruded Mg alloys. In Fig. 2b, homogenous, equiaxed grains at 0.001 s 1 and e ¼ 0.15 indicates a full recrystallization, i.e., nucleation and ripening in comparison to the initial state (Fig. 2a). Fig. 2c presents the microstructure at 0.1 s 1 and e ¼ 0.15 in which lenticular twins (in red) start emerging. Twinning is confirmed by the 901 rotation in the {0001} pole figure, which is caused by the presence of f1012g o1011 4 twins. In general, at a given strain, the twin volume fraction increases as the strain rate increases, and this can be seen from the basal texture intensity in the {0001} pole figure in Fig. 2. But the intensity at 0.5 s 1 and e ¼ 0.04 is slightly higher than that of 0.8 s 1 and e ¼ 0.04. This interesting phenomenon is worth further investigations. When the strain rate increases to 0.5 s 1, as shown in Fig. 2d, profuse, large f1012g o1011 4 twins can be observed in a coarse, nonrecrystallized grain in addition to the small twins in the fine grains. A strong basal texture can be observed in the pole figure. In the meantime, dynamically recrystallized, ‘‘necklace’’ grains nucleate and grow, surrounding the parent grains. Some new grains invade into the matrix along the twin boundaries. Fig. 2e–g show the microstructure at 0.8 s 1, e ¼ 0.04, 0.15, and 0.60, respectively. The f1012g o 1011 4 twinning dominates the deformation and gives rise to an enhanced basal texture, and then the basal texture weakens due to dynamic recrystallization. At e ¼ 0.04, high density twins can be observed. Meanwhile, fine ‘‘necklace’’ grains surrounding parent grains nucleate and grow. At e ¼ 0.15, the twins grow significantly, coalesce and consume most of the original grains. The dynamically recrystallized grains are sandwiched between the impinging twins. At e ¼ 0.60, new grains also nucleate inside the twins and grow, eventually segmenting the large twins into irregular islands. No f1011g o 1012 4 contraction twins or f1011g þ f1012g double twins were found in this study. Grain size distributions under different processing conditions are presented in Fig. 3. The as-received, extruded sample presents a bimodal structure containing large parent grains and small, recrystallized grains, as seen in the double-peak distribution. At the quasi-static strain rate of 0.001 s 1 and e ¼ 0.15, only a single peak is observed because grain growth dominated the DRX. At 0.1 s 1 and e ¼ 0.15, we also see a single peak in the grain size distribution, but the grain size is finer than that of 0.001 s 1. In contrast, for the high strain rate of 0.8 s 1, at strains of e ¼ 0.04 and e ¼ 0.15, large grains and fine, dynamically recrystallized grains coexist due to twinning and DRX. As the strain was increased to e ¼ 0.60, grain growth and DRX inside the twins and at the twin boundaries changed the microstructure structure and the high peak corresponding to the large grains
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Fig. 2. EBSD inverse pole figure map and micro-texture of AM30 under uniaxial compression at 450 1C at (a) initial status; (b) 0.001 s 1, e ¼ 0.15, grains coarsen and become equiaxed; (c) 0.1 s 1, e ¼ 0.15, visible thin f1012g o 1011 4 twins appear; (d) 0.5 s 1, e ¼ 0.04, thick twins locate in parent grain and ‘‘necklace’’ grains appear, some of the latter grow along twin interfaces; (e) 0.8 s 1, e ¼ 0.04, twins propagate in both parent grains and recrystallized grains; (f) 0.8 s 1, e ¼ 0.15, twins sweep almost all parent grains and result in a very strong basal texture, and DRX mainly occurs at grain boundaries; and (g) 0.8 s 1, e ¼ 0.60, DRX grains grow at the expense of twinned grains, resulting in irregular twin remains. The pole figures are calculated based on the corresponding EBSD maps. The inverse pole figure represents the compression direction, namely the extrusion direction (ED). (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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disappears. It should be noted that DRX and grain growth are time dependent. At different strain rates, the amounts of time that the samples were exposed to the high temperature varied from tens of seconds to a fraction of a second, contributing to the variations in grain size distribution and in the microstructure. Our EBSD analysis (Fig. 2) reveals the following: (1) twinning occurs when the strain rate reaches and exceeds 0.1 s 1 at 450 1C; (2) twinning strongly depends upon strain rate at 450 1C, and twin volume fraction increases rapidly as the strain rate exceeds a threshold value; (3) twinning and DRX interact and influence the ultimate microstructure. The strain rate and temperature influence the critical resolved shear stress (CRSS) of non-basal slip markedly, however their influence on CRSS of twinning is comparatively minimal [1,24]. At the quasi-static strain rate of 0.001 s 1, the dominant deformation mechanism is dislocation slip, including non-basal slip as the CRSS of the non-basal slip decreases as the temperature increases [25]. As the strain rate increases, twinning starts to play a marked role in the deformation. In general, dislocation slip is both temperature and strain rate dependent. An increase in the temperature leads to a decrease in the CRSS of slip, thus
Fig. 3. Grain size distribution of AM30 at various strain rates and strain levels compressed at 450 1C. The initial grain size shows a bimodal structure (double peaks) containing large parent grains and small, recrystallized grains. At 10 3 s 1 and e ¼ 0.15, the microstructure was fully recrystallized. At 0.8 s 1, e ¼ 0.04 and e ¼ 0.15, large parent grains and dynamically recrystallized grain coexists, whereas at e ¼ 0.60, the twinned parent grains were also dynamically recrystallized and the corresponding high peak disappears.
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promoting dislocation activity; however, an increase in strain rate strongly increases the CRSS of slips, as shown in Fig. 1, where the flow stress significantly increases with strain rate. This is consistent with hardening models proposed by Mecking and Kocks [26], Beyerlein and Tome´ [27], and Oppedal et al. [28]. Using a full-constraint Taylor model, Barnett [29] calculated the CRSS for basal, prismatic, second-order pyramidal slip, and f1012g twinning. It was suggested that the CRSS for basal slip and f1012g twinning is relatively temperature independent. Ulacia et al. [30] did extensive studies on the mechanical behavior of an AZ31 Mg alloy over a wide range of strain rates and temperatures, and they confirmed that the CRSS for f1012g twinning remained relatively constant with temperature up to 400 1C. Based on these previous studies in the literature, it could be generally concluded that the CRSS of basal slip and twinning are less sensitive to the temperature and strain rate in Mg alloys [8,24], while the CRSS of nonbasal slips are strongly sensitive to temperature and strain rate. Although Barnett’s calculation [29] shows that the CRSS for basal slip is about 2.5 MPa, and the CRSS for f1012g twinning is about 32 MPa in an AZ31 Mg alloy, the most recent measurement [31] of the CRSS for f1012g twinning in pure single crystal Mg is only 2.4 MPa, close to that of basal slip. These calculations and measurements indicate that the CRSS of f1012g twinning can actually be very low. Therefore, as strain rate increases, a higher CRSS for dislocation slip promotes f1012g twinning, as shown in the present experimental work. The mechanical testing (Fig. 1) reflects the competition between mechanical twinning and DRX, as shown in previous studies [9,10,12–15,17–19]. Our EBSD analysis of samples examined after interrupted mechanical tests further clarifies the significance of twinning in the moderate strain rates. As Fig. 2c depicts, twinning does occur concomitantly with DRX when the strain rate reaches 0.1 s 1. But this concurrence of twinning and DRX is not reflected in the stress–strain curve (Fig. 1), because DRX dominates the microstructure evolution. Although twinning is occurring as observed in Fig. 2, its contribution to hardening is overwhelmed by the softening from DRX so that the stress–strain curve remains parabolic in contrast to the cases of the higher strain rates (0.5 s 1 and 0.8 s 1). This indicates that the ‘‘sigmoidal’’ hardening is indeed a signature of twinning, but twinning may have already occurred when such ‘‘sigmoidal’’ hardening is not present in stress–strain curves. Hence, twinning starts to occur at a lower strain rate than previously thought [9–11]. When the strain rate increases to 0.5 s 1 and 0.8 s 1, twinning with a higher volume fraction corresponding to the increased basal texture intensity in the {0001} pole figure in Fig. 2 is activated and this generates an observable effect in strain hardening.
Fig. 4. Schematic illustration of twinning behavior and microstructure evolution in the AM30 Mg alloy compression at 450 1C and various strain rates: (a) low strain rate 0.001 s 1, the microstructure evolution is controlled by dynamic recrystallization (DRX) and grain growth; (b) intermediate strain rate 0.1 s 1, sporadic twinning occurs and twins are completely consumed by DRX; (c) high strain rate 0.8 s 1, profuse twinning is activated; DRX occurs inside the twins and at the twin boundaries; the twins are then invaded by new grains, lose their original morphology, and become unrecognizable in optical microscopy.
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Our EBSD analysis indicates that the interaction of DRX and twinning strongly influences the microstructure and texture at different strain levels. When twinning becomes more pronounced as the strain rate increases, twins with relatively coherent twin boundaries can be observed in optical microscopy (Fig. 2d and e) at low strain levels. As the plastic flow proceeds, however, the twins are invaded by the DRX grains and therefore are no longer recognizable in optical microscopy; hence, EBSD is preferred for microstructure analysis. A few DRX grains nucleate inside the twins (Fig. 2f), because the twins orientations favor pyramidal slip, which facilitates DRX inside the twins [18]. Meanwhile, the twin boundaries are the preferred nucleation sites for DRX grains. As a result, the growth of the DRX grains consumes parts of or entire twins, leading to highly irregular, incoherent twin boundaries. Consequently, caution should be exercised when optical microscopy is used to examine the microstructure of samples deformed at elevated temperatures. The effect of strain rate on twinning and microstructure of the AM30 compressed at 450 1C is schematically illustrated in Fig. 4. At quasi-static strain rates, slip dominates in deformation. Nucleation and grain growth by DRX determine the microstructure and texture evolution. Due to the low strain rate, the recrystallized grains quickly ripen, as shown in Fig. 4a. At a moderate, intermediate strain rate (0.1 s 1), twinning is activated as the CRSS for dislocation slip increases, but the twinning competes with DRX and has little contribution to hardening (Fig. 4b). At high strain rates (40.5 s 1), profuse twinning dominates the plastic deformation as fine grains nucleate along parent grain boundaries; concomitantly, DRX occurs at the twin boundaries and inside the twins (Fig. 4c). As deformation continues at the higher strain rates, the twins quickly lose their original morphology because of DRX, but twinning has an important effect on the texture.
4. Conclusions Using EBSD, we investigated the effect of strain rate ( o1.0 s 1) on twinning and microstructure at 450 1C in an extruded AM30 alloy. At quasi-static strain rates, DRX and grain growth dominate the microstructure and texture evolution. Although sporadic twins begin to appear at strain rate 0.1 s 1, their contribution to hardening is insignificant and no sigmoidal hardening was observed. DRX plays a dominant role in the softening. As the strain rate increases to 0.5 s 1, the twin volume fraction increases significantly and makes an observable
contribution to the hardening. At strain rate of 0.8 s 1, profuse twinning is activated and a peak stress followed by softening due to DRX is observed. The twins quickly lose their original morphology because of DRX.
Acknowledgments The authors are grateful to the financial support from the Department of Energy, Contract no. DE-FC-26-06NT42755, and the Center for Advanced Vehicular Systems (CAVS) at Mississippi State University. Thanks are extended to H.R. Wainwright for her assistance with the preparation of the manuscript. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
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