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Dislocation-modified high index twinning mechanisms in deformed magnesium alloy
Journal of Alloys and Compounds 787 (2019) 423e428
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Dislocation-modified high index twinning mechanisms in deformed magnesium alloy B. Zhou, J. Tai, Y. Lu, M.L. Sui* Beijing Key Laboratory of Microstructure and Properties of Solids, Institute of Microstructure and Properties of Advanced Materials, Beijing University of Technology, Beijing 100124, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 November 2018 Received in revised form 14 January 2019 Accepted 6 February 2019 Available online 7 February 2019
A combination of high resolution transmission electron microscopy (HRTEM) and electron backscatter diffraction (EBSD) was applied to study the formation mechanism associated with high index twins in a deformed magnesium alloy. Utilizing the information obtained about the formation of high index twins by identifying their structural characteristics and interfacial atomic arrangement, we found that high index twins were obtained from dislocation-modified double twins. With increasing deformation, misorientations between the matrix and deformation twin increased, resulting in the evolution of double twins into a variety of structures. This discovery about high index twin formation can greatly improve the ability to identify various types of double twins in EBSD data. This mechanism provides a new insight into the formation of high index twins and aids in understanding the material fractures caused by double twinning. © 2019 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloy Double twinning High index twin Dislocation EBSD HRTEM
1. Introduction Magnesium alloys are lightweight, environmentally friendly materials that are widely used in a variety of areas [1e4]. Applications, especially in the aerospace and automotive industries, require a deep understanding of the deformation mechanism in these alloys. Currently, no magnesium alloys sufficiently meet the Von Mises criteria for slip systems [5]. This slippage deficiency leads to poor plasticity at room temperature. Twinning plays a significant role in mediating the strain along the c-axis and promotes the activation of other slip systems in magnesium alloys. The major types of twins include primary compression twins f1011g and tension twins f1012g, both of which are commonly observed and often induce anisotropic behavior during plastic deformation [6]. High index twins, such as f3034g and f1015g, have also been observed [7]. To explain the formation of these high index twins, previous researchers proposed a shear-shuffle mechanism [8,9]. However, according to shear-shuffle twinning theory, the larger the twinning shear at the high index twin plane, the greater the amount of energy required; thus, the formation of high index twins would not be energetically favored [6,10]. Yoshinaga et al. [7,11]
* Corresponding author. E-mail address: [email protected] (M.L. Sui). https://doi.org/10.1016/j.jallcom.2019.02.079 0925-8388/© 2019 Elsevier B.V. All rights reserved.
observed that well developed f3034g twins were either an array of f1013g twins along the f3034g habit plane or developed from slightly rotated f1011g twins. Reed-Hill et al. [9] also observed the f3034g twin band being formed by tightly grouped clusters of small double twins rather than a simple twinning system. More recently, Barnett et al. [12] found that the poles of f1011g f1012g double twin boundary planes were around the f3034g poles of the matrix. By using TEM selected area electron diffraction (SAED) pattern analyses, Wang et al. [13] identified a group of side-by-side high index twins of f3032g and f1015g. Moreover, large numbers of double twins were discovered in the fracture of magnesium alloys during plastic deformation [14e17], a finding which indicated the important role of double twins in the mechanical properties of magnesium alloys. Although some studies have focused on high index twins in Mg alloys, the formation mechanism of high index twins is still not well understood. In particular, there is a lack of direct observations of high index twin boundaries at the atomic scale. In addition, EBSD technology is usually used to collect largescale statistical data about the deformation structures. However, it has generally been difficult to identify double twins from EBSD patterns, and only small fragments can be indexed [18,19]. Consequently, it is necessary to identify the reason for the low indexing rate of double twins and as a result improve the identification ability of EBSD data.
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The current research focused on the formation mechanism associated with high index twins in magnesium. The high index twins, including their orientation relationship and the atomic arrangement of the twin boundaries, were comprehensively characterized by HRTEM and SAED analyses. These revealed that the twin boundary planes were not parallel to the high index twinning planes but were nearly parallel to the f1011g plane. Moreover, the atomic structure on the two sides of the twin interface did not have mirror symmetry. Based on the experimental results, we proposed a mechanism by which high index twins develop from dislocationmodified f1011g-f1012g double twins. This formation mechanism helps elucidate the nature of high index twins that have a twin orientation relationship with the matrix but have no mirror image. As a result, we not only explained why deformation bands cannot be identified by conventional EBSD methods but also suggested a method to greatly improve the identification ability of EBSD data. 2. Experimental A commercial rolled sheet of AZ31 magnesium alloy with the following chemical composition (wt%) was used: 3.3% Al, 0.68% Zn, 0.27% Mn, 0.095% Si, 0.018% S, 0.003% P, Mg balance. Specimens measuring 20 15 10 mm3 were cut from the as-received plate and annealed at 375 C for 12 h for complete recrystallization. Then the specimens were rolled at room temperature to different reductions in thickness, which was then defined as the rolling deformation strain, and the strain rate was estimated to be about 103 s1. The cold-rolled samples were ground using a series of diamond abrasive papers (grit numbers 200e5000) and subsequently electrolytically polished at 15 V in a solution of 5% perchloric acid and 95% ethyl alcohol to remove the deformed layer introduced by mechanical polishing. EBSD measurements were characterized by means of a JEOL JSM-6500F field emission scanning electron microscope (FE-SEM) equipped with a TSL EBSD system. An accelerating voltage of 30 kV and a working distance of 21 mm were selected to get the maximum signals. The step size was 0.2 mm and the average confidence index (CI) value of the measured points was 0.48. Finally, the EBSD data were analyzed using TSLOIM software. The TEM samples were cut along the rolling direction (RD), ground to 50 mm in thickness, and then electropolished at 15 V with a solution of 10% nitric acid, 30% glycerol, and 60% methanol at 253 K. The HRTEM was performed in an aberration-corrected FEI Titan environmental TEM (ETEM) operated at 300 kV with a point resolution of 0.68 Å. 3. Results and discussion In a magnesium alloy sample with 10% deformation, many basal < a > dislocations can be observed within and around a twin band with a typical deformed microstructure, as shown in Fig. 1a, while the orientation of the matrix was along the ½1210 zone axis, as shown in Fig. 1b. Interestingly, there was an interface inside the deformation twin band, as indicated by the dashed blue line in Fig. 1a, which segmented the twin band into two parts. The SAED pattern taken at the interface between the matrix and the right part (Twin 2) of the twin band had a perfect orientation relationship of a f1015g twin with a misorientation angle/axis pair of 41 < 1210 >, as shown in Fig. 1c. However, the twin boundary (TB) (marked by the dashed red line in Fig. 1a) is obviously not parallel to the f1015g plane. Due to the high energy required for shearing along the f1015g plane [6], such high index twinning is uncommon in magnesium alloys. Fig. 1d shows the SAED pattern, which presents a f1011g twin relationship but with about a 5 deviation of the left twin (Twin 1) from the matrix (circled d in Fig. 1a).
Correspondingly, the TB (marked by the dashed green line in Fig. 1a) is parallel to the f1011g twinning plane. Fig. 1e is the SAED pattern of the interface between Twin 1 and Twin 2, showing a f1012g twin relationship with about a 1.5 deviation and the TB parallel to the f1012g twinning plane (see the dashed blue line in Fig. 1a). It is reasonable to deduce that Twin 2 is a secondary twin within the twin pair that, for the most part, accords with the shape of the primary f1011g twin (Fig. 1a). In addition, the small deviation from the perfect twin relationship (Fig. 1d and e) can be attributed to the accumulation and entanglement of dislocations around the twin band (Fig. 1a and Fig. S1 in the Supplementary materials). Hence, we propose that the f1015g high index twin comes from the f1011g f1012g double twinning system accompanied by the dislocation behavior. In the SAED pattern in Fig. 1c, the f1015g diffraction spots for the matrix and the twin match perfectly, indicating a f1015g high index twin. However, the twin boundary marked by the dashed red line in Fig. 1a is not parallel to the f1015g twinning plane, which is inconsistent with twinning theory [6]. In order to clearly understand the microstructure of this special deformation twin, an HRTEM observation and analysis were performed. The interface atomic arrangements were clearly presented by HRTEM images taken along ½1210. As shown in Fig. 2a, the interface of the matrix and Twin 2, marked as f1015g TB, obviously deviates from the ð1015Þ plane but is nearly parallel to the ð1011Þ plane, and the lattices of the matrix and the twin do not show a mirror relationship with respect to the ð1015Þ twin plane. This indicates that the f1015g high index twin was not formed by a real twinning shear along the f1015g twin plane, even though there is an apparent f1015g twinning relationship. However, in Fig. 2b, the lattices of the matrix and Twin 1 clearly exhibit a perfect symmetrical relationship with respect to the twin boundary of the ð1011Þ plane, which confirms that Twin 1 is a standard f1011g contraction twin. In addition, a large number of stacking faults (SFs) on the basal plane were contained in Twin 1, which is common for the f1011g contraction twin as a primary twin in Mg alloys [20]. Similarly, Fig. 2c shows that the basal planes of Twin 1 and Twin 2 have a symmetrical relationship about the f1012g twin plane, which reveals that Twin 2 is a typical f1012g tension twin relative to Twin 1, i.e. that these are a f1011g f1012g double twin. Therefore, the f1015g twin relationship demonstrated by the SAED pattern in Fig. 1c resulted from the orientation relationship between the f1011g f1012g double twin and the matrix, with a specific geometric relationship of 41. Based on the HRTEM observations, we deduced that the so-called high index twin came from the f1011g f1012g double twin and that real high index twinning systems are energetically unfavored in Mg alloys with a 10% deformation. Mg alloy samples with different rolling deformation strains were intensively investigated in our research, as shown in Fig. 3, and showed that twin bands with a high index twinning relationship can be primarily found in these deformed alloys. Fig. 3a shows the dark field TEM micrograph of a twin band in the 5% deformed Mg alloy. The corresponding SAED pattern (Fig. 3b) shows the rotation angle to be 37.5 in the ½1210 zone axis for the matrix and twin, which is a critical orientation relationship of f1011g f1012g double twinning identified by EBSD [19,21,22]. In this case, in Fig. 3b the diffraction spots of f1015g do not overlap. However, a twin band (Fig. 3c), presented as a f3032g high index twinning relationship, was observed in the Mg alloy with an 8% deformation strain, as shown in Fig. 3d, in which the f3032g diffraction spots can be seen to overlap and the angle difference between the matrix and twin increases to 39 . Moreover, when the deformation increased to 22%, the angle between them increased even more to 46.3 , as shown in Fig. 3e and f. In addition, we can
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Fig. 1. TEM micrograph of 10% deformed samples. (a) TEM micrograph of the deformation twin, in which the white arrow represents the rolling direction and the blue arrows indicate the dislocations. (b)e(d) Selected area diffraction patterns (SAED) corresponding to the marked areas in (a) along a ½1210 zone axis. The yellow, green, and red labels represent the matrix, Twin 1, and Twin 2, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Fig. 2. HRTEM images corresponding to the twin boundaries. (a) f1015g twin boundary; (b) f1011g twin boundary; (c) f1012g twin boundary.
observe many basal < a > dislocations around the twin boundaries and see that the dislocation density increases with the increase in deformation. More data about different rolling deformation strains can be seen in Fig. S2. Hence, a variety of high index twins with specific orientation relationships will be formed at the ½1210 zone axis, as the simulated SAED patterns show in Fig. 3gei. It should be pointed out that high index twins that are not real twins around high index planes but are, rather, some deformation structures of f1011g f1012g double twins with a specific geometric relationship are frequently observed in a series of TEM and SAED analyses. Based on the misorientation between matrices and f1011g f1012g double twins, studies have confirmed four possible variants of the f1011g f1012g double twin, namely variants SA, SB, SC, and SD with the misorientation angle/axis pairs of 37.5 < 1210 > , 30.1 < 1210 >, 66.5 < 5943 > , and 69.9 < 2421 > [14,15,18,19,21,23,24]. In our study, the variant SA was frequently observed because it has the least accommodation strain within the parent grain and the greatest potential for growth. For the variant SA, once the f1012g secondary twin nucleates in the f1011g primary twin, the secondary twin grows rapidly along the primary twin boundary, and this segment of the primary twin is transferred into the f1011g f1012g double twin with misorientation angle/ axis pairs of 37.5 < 1210 > [14,25]. To sum up the observed results, the formation mechanism of high index twins can be illustrated as in Fig. 4. In polycrystalline AZ31 magnesium alloys, in some grains f1011g compression twins may be formed as primary twins by the compression strain induced
by rolling deformation, as shown in Fig. 4a. Because the critical resolved shear stress (CRSS) value for the f1011g compression twin is much larger than that for the dislocation slip on the basal plane in magnesium alloys [5,26], a large number of basal plane dislocations will be simultaneously activated in the matrix, resulting in an accumulation of bB dislocations at the twin boundary. As a result, the lattice of the f1011g primary twin will deviate from the perfect f1011g twin relationship. In the process of deformation, the more the dislocations accumulated, the more the lattice deviated (see the SAED pattern in Fig. 1d). Meanwhile, the secondary twin of f1012g was likely to nucleate and grow in the f1011g primary twin due to the higher internal strain for primary twins and the low CRSS value for f1012g twins, as shown in Fig. 4b. Additionally, in the presence of a large number of basal dislocations in the primary twin (Fig. 1a), the lattice of the secondary twin will deflect from the f1011g-f1012g double twin relationship under the action of dislocations and will eventually be similar to high index twins with a specific orientation relationship, as shown in Fig. 4c. Obviously, the secondary twinning process will absorb a considerable amount of the stored strain energy that was inside the primary twins. Therefore, high index twins are formed by double twinning combined with a dislocation reaction, a process which absorbs a large number of dislocations. It should be noted that the twin boundaries for double twins and high index twins do not have a coherent structure, so the twin boundaries will behave as ordinary large angle boundaries. Therefore, these twin bands with large angle boundaries and nanometer-size thicknesses will easily result in a
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Fig. 3. TEM micrographs of rolled samples and corresponding SAED patterns with different deformation degrees. (a, b) 5%; (c, d) 8%; (e, f) 22%. (g)e(i) Simulated SAED patterns for various high index twins f1016g, f3032g, and f1015g, respectively.
Fig. 4. (a)e(c) Schematic illustration of the f1015g high index twin formation mechanism.
significant concentration of strain, eventually leading to the failure of the material. This mechanism provides a new pathway for explaining the formation of particular high index twins during the deformation process and may also be very helpful for understanding the fact that large numbers of double twins have been discovered in the fracture of magnesium alloys during plastic deformation [14e17]. Finding that the high index twin formation process results from a significant concentration of strain that eventually leads to the failure of the material is of great significance for research into deformation twinning in magnesium alloys.
EBSD technology is often used to analyze the evolution of deformation structures and is especially useful for identifying twinning relationships. Fig. 5a and b are the EBSD inverse pole figure (IPF) maps showing the microstructures of AZ31 magnesium alloys with 10% and 22% strain deformation, respectively. It is obvious that some zonal grains formed in the original large grains, but EBSD could only identify a very small proportion of these zonal grains as deformation twinning relationships, as shown in Fig. 5a and b indicated by the colored twin boundaries of the f1011g compression twin (white), the f1012g tension twin (black), and the
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Fig. 5. Microstructure evolution of AZ31 alloys upon rolling deformation. (a, b) The EBSD IPF maps with the conventional EBSD identification, (c, d) the EBSD IPF maps with the modified EBSD identification, and (e, f) the corresponding point-to-origin misorientation distribution profiles of the samples with rolling strains of 10% and 22%, respectively. In (a)e(d), the white lines indicate the f1011g compression twin boundaries, the black lines indicate the f1012g tension twin boundaries, and the red and the blue lines indicate the f1011g-f1012g double twin boundaries of variant SA and variant SD, respectively. In (e)e(f), the dashed gray lines indicate standard misorientation angle/axis pairs of different deformation twins. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
f1011g-f1012g double twin variants (red for SA and blue for SD), respectively. Comparing Fig. 5a and b, with an increase in deformation, more zonal grains were observed, but fewer twinning relationships were identified. Among these very few identifiable twins, the number of f1011g-f1012g double twins increased relative to the total number of twins. It is well known that twinning deformation plays a significant role in mediating strain along the caxis in magnesium alloys. However, these zonal grains, which likely result from twinning deformation, commonly cannot be recognized as deformation twins by using conventional EBSD identification (Fig. 5a and b). Based on the above SAED and HRTEM results, many f1011g f1012g double twins deviated from their standard misorientation angle/axis pairs of four variants as a result of a combination of dislocations. The dislocation-modified f1011g f1012g double twins could present as any of a variety
of high index twins with a specific geometrical orientation, such as the misorientation of the angle/axis pair from the standard 37.5 to 41 < 1210 > or more. Also, the lattices of these high index twins were confirmed as dislocation-modified f1011g f1012g double twins. Therefore, some deformation twins were unidentified in Fig. 5a and b, probably because they deviated from the double twin parameters. With an increase in deformation, more dislocations will be activated and react with the twins, changing the deformation twinning relationship even more. Obviously, it is not appropriate to identify the deformation twins and double twins using the traditional preset parameters, as the traditional method causes the indexing rate to be very low. The deviation from the double twin relationship can be checked for each angle/axis pair only by measuring the relative point-to-origin misorientation (taking the first point as a reference origin) across the twin boundaries. In
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Fig. 5e and f, it can be seen that the orientation differences of these unidentified parts differed slightly from their standard angles (the dashed gray lines). Fig. 5e shows the relative point-to-origin misorientation profiles measured along four straight lines in Fig. 5c, drawn across the zonal grains and one inner boundary. It is clear that T1 has about a 53 misorientation with the matrix and can be identified as a f1011g twin, while T*2 has a misorientation of ~90 with T1 and corresponds to a f1012g twin relationship. Moreover, T2 and T*2 respectively have 41 and 70 misorientations with the matrix, which indicates that they are modified f1011g-f1012g noncoaxial variants, namely variants SA and SD. Variant SA requires the least accommodation strain, and variant SD is favored by the highest CRSS, which obeys Schmid's law [19]. When taking the dislocation-modified double twin relationship, such as a f1015g twin with an angle/axis pair of 41 < 1210 >, as the preset parameter for the EBSD identification, the zonal grains can almost completely be identified. For example, by using the new parameter, the indexing rate increased from 23% in Fig. 5a to 98% in Fig. 5c. Similarly, Fig. 5d shows the EBSD IPF map of the 22% deformed sample, in which the indexed twin boundaries were identified by using the modified preset parameters (f1015g high index twin); then, the indexing rate increased obviously compared with that in Fig. 5b. The dislocation-modified double twins of T4 and T*4 are recognized based on their corresponding misorientation profiles in Fig. 5d. Note that the dislocation-modified double twins came to dominate as a result of the deformation twinning behavior at greater rolling deformation strain. Clearly, the high index twins observed by TEM are highly consistent with the unidentified twin structures in traditional EBSD maps, and using the modified preset parameters to identify the double twinning modified by dislocation greatly improves the identification rates obtained using EBSD maps. This modified double twinning model not only shed light on the formation mechanism of so-called high index twins with unfavorable energy but also significantly improved the identification in EBSD maps of various types of double twins, such as f1013g f1012g double twins [14] and f1012g f1012g double twins [27,28], which have formed at severe plastic deformations. 4. Conclusions In summary, the present results demonstrate the formation mechanism associated with high index twins in magnesium alloys. By using HRTEM to identify the structural characteristics and interfacial atomic arrangement, we found that high index twins do not have a coherent structure but rather some deformation structures of f1011g f1012g double twins with specific geometric relationships. With an increase in deformation, the misorientation between the matrix and a deformation twin becomes greater, resulting in the evolution of double twins into different structures. The process of dislocation-modified double twinning promotes the formation of a variety of high index twins. These results provide a new way to identify twins in EBSD data. Moreover, these results greatly aid the understanding of the role of double twinning in fractures resulting from deformations and, thus, can play a significant role in improving the properties of magnesium alloys. Acknowledgements The authors are grateful for financial support from the National Natural Science Foundation of China (Grants Nos. 11374028, U1330112 and 51621003) and the Scientific Research Key Program
of Beijing Municipal Commission of Education (KZ201310005002). (Grants No.11374028). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2019.02.079. References [1] T.M. Pollock, Weight loss with magnesium alloys, Science 328 (2010) 986e987. [2] B.-C. Suh, M.-S. Shim, K.S. Shin, N.J. Kim, Current issues in magnesium sheet alloys: where do we go from here? Scripta Mater. 84 (2014) 1e6. [3] W.J. Joost, P.E. Krajewski, Towards magnesium alloys for high-volume automotive applications, Scripta Mater. 128 (2017) 107e112. [4] G.S. Frankel, Magnesium alloys: ready for the road, Nat. Mater. 14 (2015) 1189e1190. [5] M.H. Yoo, Slip, twinning, and fracture in hexagonal close-packed metals, Metall. Trans. A 12 (1981) 409e418. [6] J.W. Christian, S. Mahajan, Deformation twinning, Prog. Mater. Sci. 39 (1995) 1e157. [7] H. Yoshinaga, R. Horiuchi, Deformation mechanisms in magnesium single crystals compressed in the direction parallel to hexagonal axis, Trans. Jpn. Inst. Met. 4 (1963) 1e8. [8] R.E. Reedhill, A study of the {1011} and {1013} twinning modes in magnesium, Trans. Am. Inst. Min. Metall. Eng. 218 (1960) 554e558. [9] R.E. Reedhill, W.D. Robertson, Additional modes of deformation twinning in magnesium, Acta Metall. 5 (1957) 717e727. [10] B. Li, E. Ma, Atomic shuffling dominated mechanism for deformation twinning in magnesium, Phys. Rev. Lett. 103 (2009), 035503. [11] H. Yoshinaga, T. Obara, S. Morozumi, Twinning deformation in magnesium compressed along the C-axis, Mater. Sci. Eng. 12 (1973) 255e264. [12] M.R. Barnett, Twinning and the ductility of magnesium alloys: Part II.“Contraction” twins, Mater. Sci. Eng. A 464 (2007) 8e16. [13] T. Wang, W. Sun, Z. Wang, X. Yang, Y. Feng, Formation and structural features of a type of uncommon deformation twin band in Mg-Y-Nd-Zn alloy, Chin. J. Rare Met. 8 (2017) 864e869. [14] P. Cizek, M.R. Barnett, Characteristics of the contraction twins formed close to the fracture surface in Mge3Ale1Zn alloy deformed in tension, Scripta Mater. 59 (2008) 959e962. [15] S. Niknejad, S. Esmaeili, N.Y. Zhou, The role of double twinning on transgranular fracture in magnesium AZ61 in a localized stress field, Acta Mater. 102 (2016) 1e16. [16] G. Liu, R. Xin, X. Shu, C. Wang, Q. Liu, The mechanism of twinning activation and variant selection in magnesium alloys dominated by slip deformation, J. Alloy. Comp. 687 (2016) 352e359. [17] F. Mokdad, D.L. Chen, D.Y. Li, Twin-twin interactions and contraction twin formation in an extruded magnesium alloy subjected to an alteration of compressive direction, J. Alloy. Comp. 737 (2018) 549e560. [18] M.R. Barnett, Z. Keshavarz, A.G. Beer, X. Ma, Non-Schmid behaviour during secondary twinning in a polycrystalline magnesium alloy, Acta Mater. 56 (2008) 5e15. Martin, L. Capolungo, L. Jiang, J.J. Jonas, Variant selection during secondary [19] E. twinning in Mge3%Al, Acta Mater. 58 (2010) 3970e3983. [20] Q. Sun, Q. Zhang, B. Li, X. Zhang, L. Tan, Q. Liu, Non-dislocation-mediated basal stacking faults inside {1011} twins, Scripta Mater. 141 (2017) 85e88. _ Martin, The role of [21] J.J. Jonas, S. Mu, T. Al-Samman, G. Gottstein, L. Jiang, E. strain accommodation during the variant selection of primary twins in magnesium, Acta Mater. 59 (2011) 2046e2056. [22] Z.-Z. Shi, Secondary twin variant selection in Mg alloy after a strain-path change, J. Alloy. Comp. 696 (2017) 510e515. [23] J. Luo, Q. Liu, W. Liu, A. Godfrey, Influence of rolling temperature on the {1011}-{1012} twinning in rolled AZ31 magnesium alloy sheets, Acta Metall. Sin. 48 (2012) 717e724. [24] S. Mu, J.J. Jonas, G. Gottstein, Variant selection of primary, secondary and tertiary twins in a deformed Mg alloy, Acta Mater. 60 (2012) 2043e2053. [25] M.D. Nave, M.R. Barnett, Microstructures and textures of pure magnesium deformed in plane-strain compression, Scripta Mater. 51 (2004) 881e885. [26] Y.T. Zhu, X.Z. Liao, X.L. Wu, Deformation twinning in nanocrystalline materials, Prog. Mater. Sci. 57 (2012) 1e62. [27] R. Xin, C. Guo, J.J. Jonas, G. Chen, Q. Liu, Variant selection of {1012}-{1012} double twins during the tensile deformation of an AZ31 Mg alloy, Mater. Sci. Eng. A 700 (2017) 226e233. [28] H. Qiao, X.Q. Guo, S.G. Hong, P.D. Wu, Modeling of {1012}-{1012} secondary twinning in pre-compressed Mg alloy AZ31, J. Alloy. Comp. 725 (2017) 96e107.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Dislocation-modified high index twinning mechanisms in deformed magnesium alloy B. Zhou, J. Tai, Y. Lu, M.L. Sui* Beijing Key Laboratory of Microstructure and Properties of Solids, Institute of Microstructure and Properties of Advanced Materials, Beijing University of Technology, Beijing 100124, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 November 2018 Received in revised form 14 January 2019 Accepted 6 February 2019 Available online 7 February 2019
A combination of high resolution transmission electron microscopy (HRTEM) and electron backscatter diffraction (EBSD) was applied to study the formation mechanism associated with high index twins in a deformed magnesium alloy. Utilizing the information obtained about the formation of high index twins by identifying their structural characteristics and interfacial atomic arrangement, we found that high index twins were obtained from dislocation-modified double twins. With increasing deformation, misorientations between the matrix and deformation twin increased, resulting in the evolution of double twins into a variety of structures. This discovery about high index twin formation can greatly improve the ability to identify various types of double twins in EBSD data. This mechanism provides a new insight into the formation of high index twins and aids in understanding the material fractures caused by double twinning. © 2019 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloy Double twinning High index twin Dislocation EBSD HRTEM
1. Introduction Magnesium alloys are lightweight, environmentally friendly materials that are widely used in a variety of areas [1e4]. Applications, especially in the aerospace and automotive industries, require a deep understanding of the deformation mechanism in these alloys. Currently, no magnesium alloys sufficiently meet the Von Mises criteria for slip systems [5]. This slippage deficiency leads to poor plasticity at room temperature. Twinning plays a significant role in mediating the strain along the c-axis and promotes the activation of other slip systems in magnesium alloys. The major types of twins include primary compression twins f1011g and tension twins f1012g, both of which are commonly observed and often induce anisotropic behavior during plastic deformation [6]. High index twins, such as f3034g and f1015g, have also been observed [7]. To explain the formation of these high index twins, previous researchers proposed a shear-shuffle mechanism [8,9]. However, according to shear-shuffle twinning theory, the larger the twinning shear at the high index twin plane, the greater the amount of energy required; thus, the formation of high index twins would not be energetically favored [6,10]. Yoshinaga et al. [7,11]
* Corresponding author. E-mail address: [email protected] (M.L. Sui). https://doi.org/10.1016/j.jallcom.2019.02.079 0925-8388/© 2019 Elsevier B.V. All rights reserved.
observed that well developed f3034g twins were either an array of f1013g twins along the f3034g habit plane or developed from slightly rotated f1011g twins. Reed-Hill et al. [9] also observed the f3034g twin band being formed by tightly grouped clusters of small double twins rather than a simple twinning system. More recently, Barnett et al. [12] found that the poles of f1011g f1012g double twin boundary planes were around the f3034g poles of the matrix. By using TEM selected area electron diffraction (SAED) pattern analyses, Wang et al. [13] identified a group of side-by-side high index twins of f3032g and f1015g. Moreover, large numbers of double twins were discovered in the fracture of magnesium alloys during plastic deformation [14e17], a finding which indicated the important role of double twins in the mechanical properties of magnesium alloys. Although some studies have focused on high index twins in Mg alloys, the formation mechanism of high index twins is still not well understood. In particular, there is a lack of direct observations of high index twin boundaries at the atomic scale. In addition, EBSD technology is usually used to collect largescale statistical data about the deformation structures. However, it has generally been difficult to identify double twins from EBSD patterns, and only small fragments can be indexed [18,19]. Consequently, it is necessary to identify the reason for the low indexing rate of double twins and as a result improve the identification ability of EBSD data.
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The current research focused on the formation mechanism associated with high index twins in magnesium. The high index twins, including their orientation relationship and the atomic arrangement of the twin boundaries, were comprehensively characterized by HRTEM and SAED analyses. These revealed that the twin boundary planes were not parallel to the high index twinning planes but were nearly parallel to the f1011g plane. Moreover, the atomic structure on the two sides of the twin interface did not have mirror symmetry. Based on the experimental results, we proposed a mechanism by which high index twins develop from dislocationmodified f1011g-f1012g double twins. This formation mechanism helps elucidate the nature of high index twins that have a twin orientation relationship with the matrix but have no mirror image. As a result, we not only explained why deformation bands cannot be identified by conventional EBSD methods but also suggested a method to greatly improve the identification ability of EBSD data. 2. Experimental A commercial rolled sheet of AZ31 magnesium alloy with the following chemical composition (wt%) was used: 3.3% Al, 0.68% Zn, 0.27% Mn, 0.095% Si, 0.018% S, 0.003% P, Mg balance. Specimens measuring 20 15 10 mm3 were cut from the as-received plate and annealed at 375 C for 12 h for complete recrystallization. Then the specimens were rolled at room temperature to different reductions in thickness, which was then defined as the rolling deformation strain, and the strain rate was estimated to be about 103 s1. The cold-rolled samples were ground using a series of diamond abrasive papers (grit numbers 200e5000) and subsequently electrolytically polished at 15 V in a solution of 5% perchloric acid and 95% ethyl alcohol to remove the deformed layer introduced by mechanical polishing. EBSD measurements were characterized by means of a JEOL JSM-6500F field emission scanning electron microscope (FE-SEM) equipped with a TSL EBSD system. An accelerating voltage of 30 kV and a working distance of 21 mm were selected to get the maximum signals. The step size was 0.2 mm and the average confidence index (CI) value of the measured points was 0.48. Finally, the EBSD data were analyzed using TSLOIM software. The TEM samples were cut along the rolling direction (RD), ground to 50 mm in thickness, and then electropolished at 15 V with a solution of 10% nitric acid, 30% glycerol, and 60% methanol at 253 K. The HRTEM was performed in an aberration-corrected FEI Titan environmental TEM (ETEM) operated at 300 kV with a point resolution of 0.68 Å. 3. Results and discussion In a magnesium alloy sample with 10% deformation, many basal < a > dislocations can be observed within and around a twin band with a typical deformed microstructure, as shown in Fig. 1a, while the orientation of the matrix was along the ½1210 zone axis, as shown in Fig. 1b. Interestingly, there was an interface inside the deformation twin band, as indicated by the dashed blue line in Fig. 1a, which segmented the twin band into two parts. The SAED pattern taken at the interface between the matrix and the right part (Twin 2) of the twin band had a perfect orientation relationship of a f1015g twin with a misorientation angle/axis pair of 41 < 1210 >, as shown in Fig. 1c. However, the twin boundary (TB) (marked by the dashed red line in Fig. 1a) is obviously not parallel to the f1015g plane. Due to the high energy required for shearing along the f1015g plane [6], such high index twinning is uncommon in magnesium alloys. Fig. 1d shows the SAED pattern, which presents a f1011g twin relationship but with about a 5 deviation of the left twin (Twin 1) from the matrix (circled d in Fig. 1a).
Correspondingly, the TB (marked by the dashed green line in Fig. 1a) is parallel to the f1011g twinning plane. Fig. 1e is the SAED pattern of the interface between Twin 1 and Twin 2, showing a f1012g twin relationship with about a 1.5 deviation and the TB parallel to the f1012g twinning plane (see the dashed blue line in Fig. 1a). It is reasonable to deduce that Twin 2 is a secondary twin within the twin pair that, for the most part, accords with the shape of the primary f1011g twin (Fig. 1a). In addition, the small deviation from the perfect twin relationship (Fig. 1d and e) can be attributed to the accumulation and entanglement of dislocations around the twin band (Fig. 1a and Fig. S1 in the Supplementary materials). Hence, we propose that the f1015g high index twin comes from the f1011g f1012g double twinning system accompanied by the dislocation behavior. In the SAED pattern in Fig. 1c, the f1015g diffraction spots for the matrix and the twin match perfectly, indicating a f1015g high index twin. However, the twin boundary marked by the dashed red line in Fig. 1a is not parallel to the f1015g twinning plane, which is inconsistent with twinning theory [6]. In order to clearly understand the microstructure of this special deformation twin, an HRTEM observation and analysis were performed. The interface atomic arrangements were clearly presented by HRTEM images taken along ½1210. As shown in Fig. 2a, the interface of the matrix and Twin 2, marked as f1015g TB, obviously deviates from the ð1015Þ plane but is nearly parallel to the ð1011Þ plane, and the lattices of the matrix and the twin do not show a mirror relationship with respect to the ð1015Þ twin plane. This indicates that the f1015g high index twin was not formed by a real twinning shear along the f1015g twin plane, even though there is an apparent f1015g twinning relationship. However, in Fig. 2b, the lattices of the matrix and Twin 1 clearly exhibit a perfect symmetrical relationship with respect to the twin boundary of the ð1011Þ plane, which confirms that Twin 1 is a standard f1011g contraction twin. In addition, a large number of stacking faults (SFs) on the basal plane were contained in Twin 1, which is common for the f1011g contraction twin as a primary twin in Mg alloys [20]. Similarly, Fig. 2c shows that the basal planes of Twin 1 and Twin 2 have a symmetrical relationship about the f1012g twin plane, which reveals that Twin 2 is a typical f1012g tension twin relative to Twin 1, i.e. that these are a f1011g f1012g double twin. Therefore, the f1015g twin relationship demonstrated by the SAED pattern in Fig. 1c resulted from the orientation relationship between the f1011g f1012g double twin and the matrix, with a specific geometric relationship of 41. Based on the HRTEM observations, we deduced that the so-called high index twin came from the f1011g f1012g double twin and that real high index twinning systems are energetically unfavored in Mg alloys with a 10% deformation. Mg alloy samples with different rolling deformation strains were intensively investigated in our research, as shown in Fig. 3, and showed that twin bands with a high index twinning relationship can be primarily found in these deformed alloys. Fig. 3a shows the dark field TEM micrograph of a twin band in the 5% deformed Mg alloy. The corresponding SAED pattern (Fig. 3b) shows the rotation angle to be 37.5 in the ½1210 zone axis for the matrix and twin, which is a critical orientation relationship of f1011g f1012g double twinning identified by EBSD [19,21,22]. In this case, in Fig. 3b the diffraction spots of f1015g do not overlap. However, a twin band (Fig. 3c), presented as a f3032g high index twinning relationship, was observed in the Mg alloy with an 8% deformation strain, as shown in Fig. 3d, in which the f3032g diffraction spots can be seen to overlap and the angle difference between the matrix and twin increases to 39 . Moreover, when the deformation increased to 22%, the angle between them increased even more to 46.3 , as shown in Fig. 3e and f. In addition, we can
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Fig. 1. TEM micrograph of 10% deformed samples. (a) TEM micrograph of the deformation twin, in which the white arrow represents the rolling direction and the blue arrows indicate the dislocations. (b)e(d) Selected area diffraction patterns (SAED) corresponding to the marked areas in (a) along a ½1210 zone axis. The yellow, green, and red labels represent the matrix, Twin 1, and Twin 2, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
Fig. 2. HRTEM images corresponding to the twin boundaries. (a) f1015g twin boundary; (b) f1011g twin boundary; (c) f1012g twin boundary.
observe many basal < a > dislocations around the twin boundaries and see that the dislocation density increases with the increase in deformation. More data about different rolling deformation strains can be seen in Fig. S2. Hence, a variety of high index twins with specific orientation relationships will be formed at the ½1210 zone axis, as the simulated SAED patterns show in Fig. 3gei. It should be pointed out that high index twins that are not real twins around high index planes but are, rather, some deformation structures of f1011g f1012g double twins with a specific geometric relationship are frequently observed in a series of TEM and SAED analyses. Based on the misorientation between matrices and f1011g f1012g double twins, studies have confirmed four possible variants of the f1011g f1012g double twin, namely variants SA, SB, SC, and SD with the misorientation angle/axis pairs of 37.5 < 1210 > , 30.1 < 1210 >, 66.5 < 5943 > , and 69.9 < 2421 > [14,15,18,19,21,23,24]. In our study, the variant SA was frequently observed because it has the least accommodation strain within the parent grain and the greatest potential for growth. For the variant SA, once the f1012g secondary twin nucleates in the f1011g primary twin, the secondary twin grows rapidly along the primary twin boundary, and this segment of the primary twin is transferred into the f1011g f1012g double twin with misorientation angle/ axis pairs of 37.5 < 1210 > [14,25]. To sum up the observed results, the formation mechanism of high index twins can be illustrated as in Fig. 4. In polycrystalline AZ31 magnesium alloys, in some grains f1011g compression twins may be formed as primary twins by the compression strain induced
by rolling deformation, as shown in Fig. 4a. Because the critical resolved shear stress (CRSS) value for the f1011g compression twin is much larger than that for the dislocation slip on the basal plane in magnesium alloys [5,26], a large number of basal plane dislocations will be simultaneously activated in the matrix, resulting in an accumulation of bB dislocations at the twin boundary. As a result, the lattice of the f1011g primary twin will deviate from the perfect f1011g twin relationship. In the process of deformation, the more the dislocations accumulated, the more the lattice deviated (see the SAED pattern in Fig. 1d). Meanwhile, the secondary twin of f1012g was likely to nucleate and grow in the f1011g primary twin due to the higher internal strain for primary twins and the low CRSS value for f1012g twins, as shown in Fig. 4b. Additionally, in the presence of a large number of basal dislocations in the primary twin (Fig. 1a), the lattice of the secondary twin will deflect from the f1011g-f1012g double twin relationship under the action of dislocations and will eventually be similar to high index twins with a specific orientation relationship, as shown in Fig. 4c. Obviously, the secondary twinning process will absorb a considerable amount of the stored strain energy that was inside the primary twins. Therefore, high index twins are formed by double twinning combined with a dislocation reaction, a process which absorbs a large number of dislocations. It should be noted that the twin boundaries for double twins and high index twins do not have a coherent structure, so the twin boundaries will behave as ordinary large angle boundaries. Therefore, these twin bands with large angle boundaries and nanometer-size thicknesses will easily result in a
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Fig. 3. TEM micrographs of rolled samples and corresponding SAED patterns with different deformation degrees. (a, b) 5%; (c, d) 8%; (e, f) 22%. (g)e(i) Simulated SAED patterns for various high index twins f1016g, f3032g, and f1015g, respectively.
Fig. 4. (a)e(c) Schematic illustration of the f1015g high index twin formation mechanism.
significant concentration of strain, eventually leading to the failure of the material. This mechanism provides a new pathway for explaining the formation of particular high index twins during the deformation process and may also be very helpful for understanding the fact that large numbers of double twins have been discovered in the fracture of magnesium alloys during plastic deformation [14e17]. Finding that the high index twin formation process results from a significant concentration of strain that eventually leads to the failure of the material is of great significance for research into deformation twinning in magnesium alloys.
EBSD technology is often used to analyze the evolution of deformation structures and is especially useful for identifying twinning relationships. Fig. 5a and b are the EBSD inverse pole figure (IPF) maps showing the microstructures of AZ31 magnesium alloys with 10% and 22% strain deformation, respectively. It is obvious that some zonal grains formed in the original large grains, but EBSD could only identify a very small proportion of these zonal grains as deformation twinning relationships, as shown in Fig. 5a and b indicated by the colored twin boundaries of the f1011g compression twin (white), the f1012g tension twin (black), and the
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Fig. 5. Microstructure evolution of AZ31 alloys upon rolling deformation. (a, b) The EBSD IPF maps with the conventional EBSD identification, (c, d) the EBSD IPF maps with the modified EBSD identification, and (e, f) the corresponding point-to-origin misorientation distribution profiles of the samples with rolling strains of 10% and 22%, respectively. In (a)e(d), the white lines indicate the f1011g compression twin boundaries, the black lines indicate the f1012g tension twin boundaries, and the red and the blue lines indicate the f1011g-f1012g double twin boundaries of variant SA and variant SD, respectively. In (e)e(f), the dashed gray lines indicate standard misorientation angle/axis pairs of different deformation twins. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
f1011g-f1012g double twin variants (red for SA and blue for SD), respectively. Comparing Fig. 5a and b, with an increase in deformation, more zonal grains were observed, but fewer twinning relationships were identified. Among these very few identifiable twins, the number of f1011g-f1012g double twins increased relative to the total number of twins. It is well known that twinning deformation plays a significant role in mediating strain along the caxis in magnesium alloys. However, these zonal grains, which likely result from twinning deformation, commonly cannot be recognized as deformation twins by using conventional EBSD identification (Fig. 5a and b). Based on the above SAED and HRTEM results, many f1011g f1012g double twins deviated from their standard misorientation angle/axis pairs of four variants as a result of a combination of dislocations. The dislocation-modified f1011g f1012g double twins could present as any of a variety
of high index twins with a specific geometrical orientation, such as the misorientation of the angle/axis pair from the standard 37.5 to 41 < 1210 > or more. Also, the lattices of these high index twins were confirmed as dislocation-modified f1011g f1012g double twins. Therefore, some deformation twins were unidentified in Fig. 5a and b, probably because they deviated from the double twin parameters. With an increase in deformation, more dislocations will be activated and react with the twins, changing the deformation twinning relationship even more. Obviously, it is not appropriate to identify the deformation twins and double twins using the traditional preset parameters, as the traditional method causes the indexing rate to be very low. The deviation from the double twin relationship can be checked for each angle/axis pair only by measuring the relative point-to-origin misorientation (taking the first point as a reference origin) across the twin boundaries. In
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Fig. 5e and f, it can be seen that the orientation differences of these unidentified parts differed slightly from their standard angles (the dashed gray lines). Fig. 5e shows the relative point-to-origin misorientation profiles measured along four straight lines in Fig. 5c, drawn across the zonal grains and one inner boundary. It is clear that T1 has about a 53 misorientation with the matrix and can be identified as a f1011g twin, while T*2 has a misorientation of ~90 with T1 and corresponds to a f1012g twin relationship. Moreover, T2 and T*2 respectively have 41 and 70 misorientations with the matrix, which indicates that they are modified f1011g-f1012g noncoaxial variants, namely variants SA and SD. Variant SA requires the least accommodation strain, and variant SD is favored by the highest CRSS, which obeys Schmid's law [19]. When taking the dislocation-modified double twin relationship, such as a f1015g twin with an angle/axis pair of 41 < 1210 >, as the preset parameter for the EBSD identification, the zonal grains can almost completely be identified. For example, by using the new parameter, the indexing rate increased from 23% in Fig. 5a to 98% in Fig. 5c. Similarly, Fig. 5d shows the EBSD IPF map of the 22% deformed sample, in which the indexed twin boundaries were identified by using the modified preset parameters (f1015g high index twin); then, the indexing rate increased obviously compared with that in Fig. 5b. The dislocation-modified double twins of T4 and T*4 are recognized based on their corresponding misorientation profiles in Fig. 5d. Note that the dislocation-modified double twins came to dominate as a result of the deformation twinning behavior at greater rolling deformation strain. Clearly, the high index twins observed by TEM are highly consistent with the unidentified twin structures in traditional EBSD maps, and using the modified preset parameters to identify the double twinning modified by dislocation greatly improves the identification rates obtained using EBSD maps. This modified double twinning model not only shed light on the formation mechanism of so-called high index twins with unfavorable energy but also significantly improved the identification in EBSD maps of various types of double twins, such as f1013g f1012g double twins [14] and f1012g f1012g double twins [27,28], which have formed at severe plastic deformations. 4. Conclusions In summary, the present results demonstrate the formation mechanism associated with high index twins in magnesium alloys. By using HRTEM to identify the structural characteristics and interfacial atomic arrangement, we found that high index twins do not have a coherent structure but rather some deformation structures of f1011g f1012g double twins with specific geometric relationships. With an increase in deformation, the misorientation between the matrix and a deformation twin becomes greater, resulting in the evolution of double twins into different structures. The process of dislocation-modified double twinning promotes the formation of a variety of high index twins. These results provide a new way to identify twins in EBSD data. Moreover, these results greatly aid the understanding of the role of double twinning in fractures resulting from deformations and, thus, can play a significant role in improving the properties of magnesium alloys. Acknowledgements The authors are grateful for financial support from the National Natural Science Foundation of China (Grants Nos. 11374028, U1330112 and 51621003) and the Scientific Research Key Program
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