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Microstructure and texture evolution during hot rolling and subsequent annealing of Mg–1Gd alloy
Materials Science & Engineering A 582 (2013) 194–202
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Microstructure and texture evolution during hot rolling and subsequent annealing of Mg–1Gd alloy W.X. Wu a, L. Jin a,n, F.H. Wang a, J. Sun a, Z.Y. Zhang a, W.J. Ding a,b, J. Dong a a b
National Engineering Research Center of Light Alloy Net Forming, Shanghai Jiao Tong University, Shanghai 200240, PR China State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 11 March 2013 Received in revised form 14 May 2013 Accepted 17 May 2013 Available online 18 June 2013
The microstructure, texture and tensile ductility of Mg–1Gd alloy were investigated and compared to pure Mg following multipass rolling at 300 1C and isothermal annealing at 400 1C. The addition of Gd weakens the basal texture in both the as-rolled and annealed conditions, which is related to the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling and the suppressed grain boundaries migration due to Gd segregation at grain boundary during annealing. A large number of secondary twins and shear bands formed during hot rolling of Mg–1Gd sheet may serve as favorable nucleation sites for static recrystallization during annealing. The recrystallized grains at bands/twins in Mg–1Gd alloy display a wide spread of orientations, which is similar to that in conventional Mg alloys. Pure Mg sheet shows a strong {0002} 〈11−20〉 texture component due to the preferred growth of grains with the 〈11−20〉 component during annealing. However, Gd solute segregation at grain boundary could inhibit the preferred growth of 〈11−20〉 grains, leading to a weak basal texture and a single 〈10−10〉 texture component in Mg–1Gd alloy sheet after annealing. The room-temperature ductility is significantly improved by the addition of Gd, which is mainly attributed to the texture weakening and grain refinement. & 2013 Elsevier B.V. All rights reserved.
Keywords: Magnesium Texture Recrystallization Hot rolling Annealing
1. Introduction Strong texture with the c-axis of grains perpendicular to the extrusion direction or rolling plane generally develops in conventional wrought Mg alloys, resulting in limited formability, poor ductility and strong anisotropy in their mechanical properties. One of the most effective approaches to improve the ductility of these Mg alloys is weakening the recrystallization texture. Recently, it has been found that the recrystallization texture following hot processing and annealing can be significantly weakened by the addition of rare earth (RE) elements, such as Nd, Ce, Gd and Y [1–6], and therefore, leading to improved formability and ductility. However, the mechanisms by which RE elements weaken the texture remain unclear, although several mechanisms have been proposed, such as shear bands nucleation [7,8], particle-stimulated nucleation (PSN) [9], dynamic strain aging (DSA) [3,10] and solute segregation at grain boundaries [11]. In our previous studies, we found that DRX of Mg alloys can be inhibited during hot deformation by the addition of RE elements. Under this condition, static recrystallization (SRX) between passes during multipass rolling plays an important role in the microstructure
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0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.05.080
and texture evolution of Mg–RE alloys. Therefore, an understanding of microstructure and texture evolution during annealing treatment is essential in order to optimize the rolling and annealing parameters. The previous studies have already shown that the texture of Mg–RE alloys is significantly weakened after post-deformation annealing compared to the as-rolled materials [5,12]. However, the mechanisms for such weaker texture development during static recrystallization (SRX) are not clearly defined. Farzadfar et al. [13] found that upon isothermal annealing of a single-pass rolled Mg–2.9 wt% Y alloy, the texture is weakened, and ascribed it to SRX in basal parent grains. However, Wu et al. [14] found that the tensile twins observed in hot-rolled Mg–3Gd–Zn alloy might serve as favorable sites for SRX, and the recrystallized grains were oriented randomly, resulting in a weak texture. The aim of this study is to investigate the effect of 1 wt% Gd addition on the microstructure, texture evolution and mechanical properties of Mg sheet by multipass rolling and annealing treatment, and elucidate the mechanisms by which the texture is weakened. In order to examine the effect of SRX on the microstructure and texture, the Mg–1Gd alloy is rolled at 300 1C and annealed at 400 1C to suppress DRX and promote SRX. Another goal of this study is to explore whether Mg–RE alloys can be rolled at a low temperature and annealed at a relatively high temperature, which would provide useful guidance to design a rolling processing route for Mg–RE alloys. Gd is used for such a study
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because of its high solubility in order to avoid Mg–RE particle effects.
2. Experimental Mg–1 wt% Gd (referred as Mg–1Gd hereafter) alloy was prepared from high purity Mg and Mg–25 wt% Gd master alloy, and was produced by an electric furnace under a mixed protection gas atmosphere of SF6/CO2 and casting into a preheated steel mold. Then, the as-cast ingots were sectioned into slabs with a dimension of 110 60 10 mm3 for rolling experiments. Pure Mg was selected as a reference alloy for comparison. The rolling experiments were performed on a rolling mill with a roller diameter of 320 mm. The rollers were heated at 180 1C and operated at a rolling rate of 4 m/min. A total cumulative rolling reduction of 80% was achieved through a 6-pass rolling schedule (∼20% per pass). The plates were preheated to 400 1C and held for 30 min before the 1st pass rolling. Then, the plates were taken out of the annealing furnace. When the preheated plates were cooled in the air to about 300 1C, the rolling of the samples started. Thus, the temperature at which hot rolling was conducted is considered to be 300 1C. After rolling, the sheet was quenched into water immediately, and then, reheated to 400 1C and held for 15 min as intermediate annealing for further rolling. To study the static recrystallization behavior, a small sample was taken from the quenched sheet after the 1st pass, and annealed at 400 1C for 5 min. The final sheet samples were also annealed for 30 min at 400 1C in order to compare the texture and mechanical properties of Mg–1Gd alloy and pure Mg. The rolling temperature of 300 1C was chosen to inhibit DRX, and therefore, the role of SRX on the microstructure and texture evolution can be precisely examined. Optical microscopy was used to analyze the microstructure of the sheet samples. Standard metallographic sample preparation
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techniques were used, and all specimens were etched using a solution of 3 g picric acid, 50 ml alcohol, 20 ml acetic acid and 20 ml water. The micro-texture and the macro-texture of rolled and annealed samples were examined in the plane containing the RD (Rolling Direction) and TD (Transverse Direction) using electron backscattering diffraction (EBSD) on a field emission gun scanning electron microscope (HITACHI SU-70, operated at 20 kV, EDAX/TSL EBSD system) and XRD technique using a Bruker D8 diffractometer with a Cu Kα source. Tensile tests were carried out at room temperature with the tensile direction parallel to the rolling direction of the sheet samples using a conventional testing machine ZWICK. The fracture surfaces of the tensile samples were observed using the same SEM.
3. Results 3.1. Microstructure and texture evolution during hot rolling and subsequent annealing Fig. 1a and b shows the macrostructures of the as-cast samples of pure Mg and Mg–1Gd alloy, respectively. Both samples show very large columnar grains with some millimeters in length and width, which indicates that the 1 wt% addition of Gd does not result in a grain refinement of the cast microstructure. This is in agreement with a previous study by Stanford et al. [6]. The optical microstructures of pure Mg and Mg–1Gd alloy in the as-rolled condition with a total thickness reduction of 80% are shown in Fig. 1c and d, respectively. The as-rolled pure Mg and Mg–1Gd alloy sheets both contain a large amount of twins in the microstructure. Some small grains are also observed which nucleate and grow during annealing between passes. After annealing for 30 min at 400 1C, SRX starts, and a fully recrystallized microstructure is formed as shown in Fig. 1e and f for pure Mg and Mg–1Gd
Fig. 1. Macrostructures of as-cast (a) pure Mg and (b) Mg–1Gd alloy. Microstructures in as-rolled condition: (c) pure Mg, (d) Mg–1Gd, and after annealing for 30 min at 400 1C: (e) pure Mg and (f) Mg–1Gd.
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alloy, respectively. It can be readily found that the grain size of pure Mg sheet is larger compared with the Mg–1Gd alloy, which suggests that the microstructure is significantly refined during hot rolling and subsequent annealing by the addition of 1 wt% Gd. Fig. 2 shows the {0002} pole figures of the pure Mg and Mg– 1Gd samples after 6-pass hot rolling and annealing for 30 min at 400 1C. As expected, a typical strong basal texture is formed in as-rolled pure Mg sheet, where the orientation distribution of the most basal poles is parallel to the sheet normal direction ND. After annealing, the texture strength is enhanced from 9.16 to 14.63. In contrast, the qualitative character of the Mg–1Gd alloy sheet texture is obviously distinct from typical basal texture for pure Mg sheet. The texture has a clearly broader distribution of basal poles, as compared with its distribution in the pure Mg. In addition, the Mg–1Gd alloy in as-rolled condition shows a split of basal pole intensity peak from ND toward RD, which is typically found for Mg alloys containing RE elements [12]. This is consistent with previous experimental observations and theoretical simulation in AZ31 alloys that the split of the maximum basal pole intensity about ND is due to the increased activity of pyramidal 〈c+a〉-slip [15]. After annealing, a much weaker sheet texture is obtained in Mg–1Gd alloy sheet compared with the as-rolled one.
Fig. 2. Recalculated {0002} pole figures of pure Mg and Mg–1Gd alloy sheet in as-rolled and annealed conditions.
Fig. 3a shows the microstructure of the as-rolled Mg–1Gd sample in the form of inverse pole figure (IPF) map, and their corresponding axis/angle misorientation distributions for special angle pairs with high density values. The IPF map of the as-rolled Mg–1Gd alloy sheet revealed a non-negligible proportion of twins in the deformed microstructures. The misorientation distributions in Fig. 3a indicate that the special boundaries correspond to {10 −11}–{10−12} secondary twin (381 〈11−20〉) (S-twin), {10−11} contraction twin (561 〈11−20〉) (C-twin) and {10−12} extension twin (861 〈11−20〉) (T-twin), respectively. An example of extension twin, contraction and secondary twin is illustrated in Fig. 3b. In Mg–1Gd alloy, there are some grains aligned favorably for extension twinning due to the weak basal texture. The accompanying {0002} pole figures illustrate the orientation relationship between extension twin, contraction twin, secondary twin and matrix. The accompanying {11−20} pole figures in Fig. 3b reveal that the matrix, the primary extension twin, contraction twin and the secondary twin share the same 〈11−20〉 rotation axis. Fig. 4 shows an EBSD Kernel average misorientation (KAM) map of the as-rolled pure Mg sheet with 80% thickness reduction. The KAM map is constructed based on the average misorientation between a measuring point and all its neighbors. Therefore, the local misorientation and strain energy can be clearly reflected in the KAM map. As can be seen from Fig. 4, the microstructure is heterogeneous, and can be divided into highly deformed regions with high KAM values and less deformed regions with low KAM values. And also, the twin boundaries are marked with different colors in Fig. 4. There are many contraction twins and secondary twins in the deformed matrix, but few extension twins are observed due to the strong basal texture of the sample. Fig. 5a shows the IPF map of pure Mg sample after annealing for 30 min at 400 1C, and a strong basal texture and a coarse microstructure can be observed. In contrast, a weak texture and fine-grained microstructure are observed in the Mg–1Gd alloy (Fig. 5b). The corresponding inverse pole figures in the RD of pure Mg and Mg–1Gd alloy samples are shown in Fig. 5c and d, respectively. Interestingly, there is a distinct difference in the texture component between pure Mg and Mg–1Gd alloy. The 〈10 −10〉 and 〈11−20〉 fiber texture components are both observed in annealed pure Mg sheet, which is typically found in extruded AZ31 Mg alloys. In contrast, only 〈10−10〉 texture component is observed for Mg–1Gd alloy sheet. Although the inverse pole figures show a
Fig. 3. (a) IPF map of the as-rolled Mg–1Gd sample and their corresponding axis/angle misorientation distributions for special angle pairs and (b) examples of extension, contraction and secondary twins. The accompanying {0002} pole figures illustrate the crystallographic arrangement of the twins and the matrix. The accompanying {11−20} pole figures indicate a common rotation axis during both twinning deformation.
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Fig. 4. Kernel average misorientation map constructed using the orientation relationship with 3rd neighbors, of 80% hot rolled pure Mg sample. Twin boundaries are marked with different colors: red lines correspond to extension twin, green lines to contraction twin and blue lines to secondary twin. (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 5. IPF map of (a) pure Mg, (b) Mg–1Gd alloy sheet after annealing for 30 min at 400 1C; corresponding inverse pole figures in the RD of (c) pure Mg and (d) Mg–1Gd alloy sheet; (e) corresponding misorientation angle distribution and (f) grain size distribution. Note the higher magnification in (b).
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similar peak intensity in pure Mg and Mg–1Gd alloy, the basal texture is weakened significantly by Gd addition. Fig. 5e shows the number of fractions of misorientation angles for both the annealed sheets. In pure Mg, a distribution peak around 301 is observed, which is related to (0002) 〈10−10〉 and (0002) 〈11−20〉 texture component. The majority of misorientation angles are below 451. However, the misorientation distribution appears quite different in the Mg–1Gd alloy sheet, and large misorientation angles are more significant with a distribution peak around 401, which suggests a more uniform misorientation distribution than pure Mg. This is correlated to the weak texture with the addition of Gd. Fig. 5f shows the grain size distribution in the pure Mg and Mg– 1Gd alloy sheets, as determined by OIM analysis. As can be seen, in Mg–1Gd alloy, the grain size ranges between 6 μm and 29 μm, and the average grain size is about 18 μm. However, a large variety in grain size from 15 μm to 133 μm was observed in pure Mg. Therefore, a homogeneous and fine-grained microstructure with a weak texture is obtained by the addition of Gd. 3.2. Mechanical properties The typical stress–strain curves of the hot rolled pure Mg and Mg–1Gd alloy after annealing for 30 min at 400 1C along the RD are shown in Fig. 6. The ultimate tensile strength (UTS), yield strength (YS) and elongation to failure of the annealed pure Mg sheet are 167.6 MPa, 69.5 MPa and 5.4%, respectively. In contrast, the room-temperature UTS and YS of annealed Mg–1Gd alloy sheet are 193.5 MPa and 129.9 MPa respectively, with the elongation dramatically increasing to 15.6%, indicating an enhanced combination of strength and ductility. Stanford et al. [6] have reported that Gd had effective solid solution strengthening effect on Mg sheet and the UTS and YS both increased with Gd concentration, but were basically unchanged above 1 wt% Gd. In addition, Mg–1Gd alloy shows a much finer grain size than pure Mg (Fig. 5f), which contributes to UTS and YS. Therefore, the increase in UTS and YS by the addition of Gd is mainly ascribed to the solid solution strengthening effect and the grain refinement. The tensile fractograph of the hot rolled pure Mg and Mg–1Gd sheet samples after annealing is shown in Fig. 7a and b, respectively. The fracture surface of pure Mg shows cleavage facets, which contain tear ridges, and a small number of dimples. The fracture mode of pure Mg can therefore be categorized as a quasicleavage fracture. In the Mg–1Gd alloy, however, a ductile fracture surface characterized by the presence of elongated dimples is observed.
Fig. 6. Typical stress–strain curves of the hot rolled pure Mg and Mg–1Gd alloy after annealing for 30 min at 400 1C in the tensile directions of RD.
Fig. 7. Typical SEM images showing the fractograph of the tensile tested pure Mg (a) and Mg–1Gd alloy (b) along the RD.
4. Discussion 4.1. Texture weakening mechanism by Gd addition In the present work, DRX is inhibited during hot rolling at a low temperature of 300 1C in Mg–1Gd alloy. Thus, SRX is the most important mechanism during annealing at 400 1C. Therefore, it seems reasonable to suggest that the texture weakening is related to the SRX. The recrystallization behavior of pure Mg, which is similar to conventional Mg–Al based alloys such as AZ31, has already been discussed in the literature. During the annealing process of the as-rolled pure Mg sample, nucleation occurs at preferred sites such as contraction and secondary twins with a high stored energy. Moreover, rapid recovery may occur in the less deformed matrix grains with low KAM values, which contributes to the persistence of the strong (0002) basal texture [16]. Here, more emphasis is placed on the Mg–1Gd alloy and particularly on the recrystallization texture. Fig. 8a shows the optical microstructure of Mg–1Gd sample in as-rolled condition with 20% thickness reduction. The as-rolled microstructure contains a large amount of twins, and no sign of DRX is found in the sample. Deformation accumulates in the contraction, secondary twins, and shear bands evolved from them [17,18]. Upon annealing of the rolled Mg–1Gd alloy at 400 1C for 5 min, such regions of high dislocation density are favorable sites for nucleation, as shown in Fig. 8b. In order to investigate the micro-texture of recrystallized grains, the recrystallized region is presented in more detail in
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Fig. 8. Optical microstructure of Mg–1Gd sample in (a) as-rolled condition with 20% thickness reduction and (b) after annealing for 5 min at 400 1C.
Fig. 9. (a) IPF map showing onset of SRX at bands and twins after annealing of Mg–1Gd at 400 1C: for 5 min; (b) corresponding KAM map marked with the twin boundaries; (c) (0002) pole figure of unrecrystallized grains illustrating the crystallographic arrangement of the different twins and the matrix grain and (d) (0002) pole figure of recrystallized grains.
Fig. 9a in term of an IPF map. It can be found that the distribution of recrystallization grains was very heterogeneous due to the heterogeneity of the as-deformed microstructure. These observations show that SRX takes place at twins and shear bands. Fig. 9b shows the corresponding KAM map. Newly recrystallized grains are easily detected by their low KAM values. Different types of
twin boundaries are marked in different colors, and a large amount of primary {10−12} extension twins can be seen in Fig. 9b. In addition, a few {10−11} contraction twins and some {10−11}–{10−12} secondary twins are also presented in the as-rolled sample after annealing for 5 min at 400 1C. Fig. 9c displays (0002) pole figure of unrecrystallized grains illustrating
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Fig. 10. (0001) 〈11−20〉 basal slip Schmid factor distribution maps before the tensile test: (a) pure Mg; (b) Mg–1Gd; and corresponding fraction (c) before and (d) after tensile test. (For interpretation of references to color in this figure, the reader is referred to the web version of this article.)
the crystallographic arrangement of the different twins and the matrix grain. T-twin A corresponds to region A, and T-twin B to region B (Fig. 9a). Thus, regions A and B are two different twin variants. According to the work by Nave and Barnett [19], three different misorientation relationships are possible between different twin variants because of the six {10−12} planes available for twinning, i.e. 7.41〈1−210〉, 60.01〈10−10〉, 60.41〈8−1−70〉. In the current work, the misorientation relationship between the two variants is a 551〈10−10〉 relationship (Fig. 9a), which is close to 60.01〈10−10〉 relationship. The difference in rotation angle may result from lattice rotation due to deformation. In the early stage of rolling, primary {10−12} extension twinning can be easily activated. Under further deformation, most of the matrix is consumed and reoriented by extension twinning, which contributes to the formation of basal texture. Moreover, contraction and secondary twins develop in the primary extension twinning region in order to accommodate deformation. This phenomenon was also reported by Jiang et al. [20], who found that a high density of secondary {10−12} twins form on the primary {10−12} extension twins at a high compression strain level in AZ31 alloy. Fig. 9d shows the orientation of the recrystallized grains in terms of (0002) pole figure. Apparently, the new grains nucleated at bands and secondary twins show a much wider spread of orientations than that of deformed parent grains and twins, leading to the weakening of the basal texture. As can be seen from Fig. 1c and d, there is more twinning in the as-rolled Mg–1Gd alloy compared with pure Mg. This result is consistent with the work by Hantzsche et al. [12], who found that the as-rolled microstructure of Mg alloys with higher RE contents contains a higher amount of twins, especially contraction and secondary
twins, than those with lower RE content. In addition, Muzyk et al. [21] calculated the generalized stacking fault energies of Mg alloys using density functional theory, and demonstrated that alloying with RE increases the possibility of twin formation during deformation of Mg alloys. Therefore, it is reasonable to suggest that alloying with Gd would increase the tendency for contraction and secondary twinning, which contributes to the texture weakening after annealing. Solute atoms with a larger atomic size than solvent atoms will tend to segregate to the grain boundaries in order to reduce the elastic strain energy [22,23]. Gd is significantly larger in atomic size than Mg, and it is confirmed by atom probe tomography (APT) that Gd solute atoms segregate strongly to grain boundaries [11]. During the growth of recrystallized grains, Gd solute atoms will move along with the boundary via slow diffusion, and thus reduce the grain boundary mobility. On the other hand, there are indications in the literature that with increasing segregation, the energy of special grain boundaries tends towards that of random boundaries [24]. This could eventually reduce the driving force for the preferred migration of special grain boundary, such as 301 〈0001〉 boundary in the conventional Mg alloys. Consequently, it is responsible for the texture weakening after annealing. Actually, the microstructure and texture evolution during static recrystallization of conventional Mg alloys without RE have been investigated extensively. Yi [16] and Li [18] have found that in the early stage of annealing, the orientation of recrystallized grains formed at secondary and contraction twins is also evenly distributed in pure Mg and AZ31 alloy. Farzadfar et al. [25] found that in Mg–2.9Zn alloy, new grains recrystallized in bands exhibit orientations close to the ones associated with twins, and contribute
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thus to texture weakening. With the increase in grain size during coarsening, the maximum intensity of basal pole figures rises linearly. They ascribed this texture strengthening to the consumption of small recrystallized grains by larger ones. As can be seen from Fig. 5c, besides {0002} 〈10−10〉 texture component, a strong {0002} 〈11−20〉 component is observed in as-rolled pure Mg after annealing for 30 min at 400 1C. This is because grain growth is controlled mainly by the preferred growth of grains with the 〈11 −20〉 component during annealing [26], and it could be speculated that, upon further annealing, the {0002} 〈11−20〉 component strengthens gradually and becomes the main texture component. However, in Mg–1Gd alloy, only 〈10−10〉 texture component is observed, which indicates that the growth of grains with 〈11−20〉 component is restricted by solute segregation of Gd. From the above analysis, it can be concluded that the recrystallization texture modification by the addition of Gd mainly takes place during the growth step. The growth of recrystallized grains of Mg–1Gd alloy sample is inhibited by solute segregation of Gd, which is a major contributor to the texture weakening. However, the role of deformation behavior in the recrystallization texture evolution cannot be neglected entirely. As mentioned in Section 3.1, the split of basal pole intensity peak from ND toward RD is a typical indication of pyramidal 〈c+a〉-slip. It has been reported that the addition of Gd to Mg alters the bond energy between Gd atoms and Mg atoms [27], increasing the probability of non-basal slip. Sandlöbes et al. [17,28] proposed that an addition of Y to the Mg alloy can influence the stacking fault energy (SFE) on either basal or pyramidal planes and reduce the differences of critical resolved shear stresses (CRSS) for the corresponding deformation mechanisms, and therefore promote the activity of pyramidal 〈c+a〉-slip. Therefore, in the current work, the activity of pyramidal 〈c+a〉-slip of Mg sheet during rolling could be enhanced by Gd additions, which results in the split of the maximum basal pole intensity. Thus, the texture strength of the as-rolled Mg–1Gd alloy sheet is decreased significantly (see Fig. 2). After annealing for 30 min at 400 1C, the recrystallized grains show a much wider spread of orientations than the basal axis orientations of deformed matrix and the texture strength is decreased further. Therefore, the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling by Gd addition is also a non-negligible contributor to the texture weakening. 4.2. Effect of Gd on the ductility of Mg sheet As can be seen from Fig. 6, the room-temperature ductility of Mg sheet is significantly improved by the addition of Gd. In order to fully understand the effect of Gd addition on the ductility of hot rolled Mg sheet, the distribution of the (0001) 〈11−20〉 basal slip Schmid factor is investigated by EBSD analysis. The results in Fig. 10a and b show that the blue grains usually have their basal planes parallel to the tensile direction, as shown by hexagonal unit cell. The blue grains have a relatively lower (0001) 〈11−20〉 basal slip Schmid factor and is unfavorable for basal slip. However, the red grains usually have their basal planes inclined about 451 to the tensile planes as shown by hexagonal unit cell. Thus, the red grains have a higher (0001) 〈11−20〉 basal slip Schmid factor, which is favorable for basal slip. Compared to pure Mg, a large number of red grains are observed in Mg–1Gd alloy, which is correlated to the weak texture by Gd addition. Fig. 10c and d shows the fraction distribution of the Schmid factor for (0001) 〈11−20〉 basal slip along the RD of both sheets before and after tensile test, respectively. It can be seen that about 35% grains in pure Mg show low Schmid factor values less than 0.1. The fraction of grains gradually decreases with the increasing Schmid factor. The highest Schmid factor ( 40.4) corresponds to the lowest fraction of 0.04. In contrast, there is a more uniform
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Fig. 11. Variation of ln s against ln ε for pure Mg and Mg–1Gd alloy.
distribution of the Schmid factor for (0001) 〈11−20〉 basal slip along the RD of Mg–1Gd alloy sheet. The fraction of grains with a high Schmid factor (40.4) is greater than 0.18. In Mg–1Gd alloy, the proportion of deforming grains by basal slip is markedly higher than in pure Mg, due to the weaker basal texture component. After the tensile test, the fraction of grains with a high Schmid factor (40.4) for basal slip is decreased for both materials (Fig. 10d), which means that basal slip is the dominant deformation mechanism during tensile test. Therefore, we propose here that the weak texture is a major contributor to the improved ductility. This result is consistent with the work by Stanford et al. [6], who found that a significant increase in ductility is achieved due to texture weakening resulting from Gd addition. Moreover, from Fig. 5f, the annealed microstructure is significantly refined by the addition of Gd due to the solute segregation effect at grain boundaries. Therefore, grain refinement would be another important contributor to the enhanced ductility. In addition, strain hardening behavior is another metallurgical factor that may contribute to ductility, which is characterized by a power law relationship, s ¼Kεn, where s and ε are the true stress and strain, K is the strength coefficient and n is the strain hardening exponent [15]. A high value of n would inhibit inhomogeneous deformation, and thus, improve formability and ductility. Fig. 11 shows the variation of ln s against ln ε of tensile data for the pure Mg and Mg–1Gd alloy along the RD. The obtained n values during tension along the RD are 0.326 and 0.172 for pure Mg and Mg–1Gd, respectively. Apparently, the strain hardening exponent n value is decreased by the addition of Gd, which indicates that this parameter does not contribute to the increased ductility obtained in Mg–1Gd alloys.
5. Conclusions The microstructure, texture and tensile ductility of Mg–1Gd alloy and pure Mg sheet were investigated following rolling and subsequent isothermal annealing. The addition of Gd weakens the basal texture in both the as-hot rolled and annealed conditions. This is related to the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling by Gd addition and the suppressed grain boundaries migration due to solute segregation of Gd at grain boundary during annealing. Pure Mg sheet shows a strong {0002} 〈11−20〉 texture component due to the preferred growth of grains with the 〈11−20〉 component. However, Gd solute segregation at grain boundary could inhibit the preferred growth of grains with the 〈11−20〉 component, which leads to a weak basal texture and a
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single 〈10−10〉 texture component in Mg–1Gd alloy sheet. A large amount of secondary twins and deformation bands form during hot rolling of Mg–1Gd sheet, and may serve as favorable nucleation sites for SRX during annealing. The recrystallized grains at bands/twins in Mg–1Gd alloy display a relatively wide spread of orientations. The room temperature ductility is significantly improved by the addition of Gd, which is mainly attributed to the texture weakening and grain refinement. Acknowledgments The authors gratefully acknowledge the support of the National High-Tech R&D Program of China (Grant no. 2011BAE22B06), the National Natural Sciences Foundation of China (Grant nos. 50901044 and 51271118) and the Shanghai Rising-Star Program (B type) (Grant no. 12QB1403300). References [1] [2] [3] [4]
N. Stanford, M. Barnett, Scr. Mater. 58 (2008) 179–182. N. Stanford, Mater. Sci. Eng. A 527 (2010) 2669–2677. L. Jin, R.K. Mishra, A.K. Sachdev, Metall. Mater. Trans. A 43 (2012) 2148–2157. J.P. Hadorn, K. Hantzsche, S.B. Yi, J. Bohlen, D. Letzig, S.R. Agnew, Metall. Mater. Trans. A 43 (2012) 1363–1375. [5] J.P. Hadorn, K. Hantzsche, S. Yi, J. Bohlen, D. Letzig, J.A. Wollmershauser, S.R. Agnew, Metall. Mater. Trans. A 43 (2012) 1347–1362. [6] N. Stanford, D. Atwell, M.R. Barnett, Acta Mater. 58 (2010) 6773–6783.
[7] N. Stanford, M.R. Barnett, Mater. Sci. Eng. A 496 (2008) 399–408. [8] H. Yan, S.W. Xu, R.S. Chen, S. Kamado, T. Honma, E.H. Han, Scr. Mater. 64 (2011) 141–144. [9] E.A. Ball, P.B. Prangnell, Scr. Metall. Mater. 31 (1994) 111–116. [10] L. Jiang, J.J. Jonas, R. Mishra, Mater. Sci. Eng. A 528 (2011) 6596–6605. [11] N. Stanford, G. Sha, J.H. Xia, S.P. Ringer, M.R. Barnett, Scr. Mater. 65 (2011) 919–921. [12] K. Hantzsche, J. Bohlen, J. Wendt, K.U. Kainer, S.B. Yi, D. Letzig, Scr. Mater. 63 (2010) 725–730. [13] S.A. Farzadfar, E. Martin, M. Sanjari, E. Essadiqi, M.A. Wells, S. Yue, Mater. Sci. Eng. A 534 (2012) 209–219. [14] D. Wu, R.S. Chen, E.H. Han, in: W.J. Poole, K.U. Kainer (Eds.), Proceedings of the 9th International Conference on Magnesium Alloys and their Applications, Canada, 2012, pp. 531–537. [15] S.R. Agnew, Ö. Duygulu, Int. J. Plasticity 21 (2005) 1161–1193. [16] S.B. Yi, I. Schestakow, S. Zaefferer, Mater. Sci. Eng. A 516 (2009) 58–64. [17] S. Sandlöbes, S. Zaefferer, I. Schestakow, S. Yi, R. Gonzalez-Martinez, Acta Mater. 59 (2011) 429–439. [18] X. Li, P. Yang, L.N. Wang, L. Meng, F. Cui, Mater. Sci. Eng. A 517 (2009) 160–169. [19] M.D. Nave, M.R. Barnett, Scr. Mater. 51 (2004) 881–885. [20] L. Jiang, J.J. Jonas, A.A. Luo, A.K. Sachdev, S. Godet, Mater. Sci. Eng. A 445–446 (2007) 302–309. [21] M. Muzyk, Z. Pakiela, K.J. Kurzydlowski, Scr. Mater. 66 (2012) 219–222. [22] J. Friedel, Adv. Phys 3 (1954) 446–507. [23] P. Wynblatt, R.C. Ku, Surf. Sci. 65 (1977) 511–531. [24] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon Press, Oxford,UK, 2004. [25] S.A. Farzadfar, E. Martin, M. Sanjari, E. Essadiqi, S. Yue, J. Mater. Sci. 47 (2012) 5488–5500. [26] S. Yi, H.G. Brokmeier, D. Letzig, J. Alloys Compd. 506 (2010) 364–371. [27] L. Gao, J. Zhou, Z.M. Sun, R.S. Chen, E.H. Han, Chin. Sci. Bull. 56 (2011) 1038–1042. [28] S. Sandlöbes, M. Friák, S. Zaefferer, A. Dick, S. Yi, D. Letzig, Z. Pei, L.F. Zhu, J. Neugebauer, D. Raabe, Acta Mater. 60 (2012) 3011–3021.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and texture evolution during hot rolling and subsequent annealing of Mg–1Gd alloy W.X. Wu a, L. Jin a,n, F.H. Wang a, J. Sun a, Z.Y. Zhang a, W.J. Ding a,b, J. Dong a a b
National Engineering Research Center of Light Alloy Net Forming, Shanghai Jiao Tong University, Shanghai 200240, PR China State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 11 March 2013 Received in revised form 14 May 2013 Accepted 17 May 2013 Available online 18 June 2013
The microstructure, texture and tensile ductility of Mg–1Gd alloy were investigated and compared to pure Mg following multipass rolling at 300 1C and isothermal annealing at 400 1C. The addition of Gd weakens the basal texture in both the as-rolled and annealed conditions, which is related to the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling and the suppressed grain boundaries migration due to Gd segregation at grain boundary during annealing. A large number of secondary twins and shear bands formed during hot rolling of Mg–1Gd sheet may serve as favorable nucleation sites for static recrystallization during annealing. The recrystallized grains at bands/twins in Mg–1Gd alloy display a wide spread of orientations, which is similar to that in conventional Mg alloys. Pure Mg sheet shows a strong {0002} 〈11−20〉 texture component due to the preferred growth of grains with the 〈11−20〉 component during annealing. However, Gd solute segregation at grain boundary could inhibit the preferred growth of 〈11−20〉 grains, leading to a weak basal texture and a single 〈10−10〉 texture component in Mg–1Gd alloy sheet after annealing. The room-temperature ductility is significantly improved by the addition of Gd, which is mainly attributed to the texture weakening and grain refinement. & 2013 Elsevier B.V. All rights reserved.
Keywords: Magnesium Texture Recrystallization Hot rolling Annealing
1. Introduction Strong texture with the c-axis of grains perpendicular to the extrusion direction or rolling plane generally develops in conventional wrought Mg alloys, resulting in limited formability, poor ductility and strong anisotropy in their mechanical properties. One of the most effective approaches to improve the ductility of these Mg alloys is weakening the recrystallization texture. Recently, it has been found that the recrystallization texture following hot processing and annealing can be significantly weakened by the addition of rare earth (RE) elements, such as Nd, Ce, Gd and Y [1–6], and therefore, leading to improved formability and ductility. However, the mechanisms by which RE elements weaken the texture remain unclear, although several mechanisms have been proposed, such as shear bands nucleation [7,8], particle-stimulated nucleation (PSN) [9], dynamic strain aging (DSA) [3,10] and solute segregation at grain boundaries [11]. In our previous studies, we found that DRX of Mg alloys can be inhibited during hot deformation by the addition of RE elements. Under this condition, static recrystallization (SRX) between passes during multipass rolling plays an important role in the microstructure
n
Corresponding author. Tel.: +86 21 34203052; fax: +86 21 34202794. E-mail address: [email protected] (L. Jin).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.05.080
and texture evolution of Mg–RE alloys. Therefore, an understanding of microstructure and texture evolution during annealing treatment is essential in order to optimize the rolling and annealing parameters. The previous studies have already shown that the texture of Mg–RE alloys is significantly weakened after post-deformation annealing compared to the as-rolled materials [5,12]. However, the mechanisms for such weaker texture development during static recrystallization (SRX) are not clearly defined. Farzadfar et al. [13] found that upon isothermal annealing of a single-pass rolled Mg–2.9 wt% Y alloy, the texture is weakened, and ascribed it to SRX in basal parent grains. However, Wu et al. [14] found that the tensile twins observed in hot-rolled Mg–3Gd–Zn alloy might serve as favorable sites for SRX, and the recrystallized grains were oriented randomly, resulting in a weak texture. The aim of this study is to investigate the effect of 1 wt% Gd addition on the microstructure, texture evolution and mechanical properties of Mg sheet by multipass rolling and annealing treatment, and elucidate the mechanisms by which the texture is weakened. In order to examine the effect of SRX on the microstructure and texture, the Mg–1Gd alloy is rolled at 300 1C and annealed at 400 1C to suppress DRX and promote SRX. Another goal of this study is to explore whether Mg–RE alloys can be rolled at a low temperature and annealed at a relatively high temperature, which would provide useful guidance to design a rolling processing route for Mg–RE alloys. Gd is used for such a study
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because of its high solubility in order to avoid Mg–RE particle effects.
2. Experimental Mg–1 wt% Gd (referred as Mg–1Gd hereafter) alloy was prepared from high purity Mg and Mg–25 wt% Gd master alloy, and was produced by an electric furnace under a mixed protection gas atmosphere of SF6/CO2 and casting into a preheated steel mold. Then, the as-cast ingots were sectioned into slabs with a dimension of 110 60 10 mm3 for rolling experiments. Pure Mg was selected as a reference alloy for comparison. The rolling experiments were performed on a rolling mill with a roller diameter of 320 mm. The rollers were heated at 180 1C and operated at a rolling rate of 4 m/min. A total cumulative rolling reduction of 80% was achieved through a 6-pass rolling schedule (∼20% per pass). The plates were preheated to 400 1C and held for 30 min before the 1st pass rolling. Then, the plates were taken out of the annealing furnace. When the preheated plates were cooled in the air to about 300 1C, the rolling of the samples started. Thus, the temperature at which hot rolling was conducted is considered to be 300 1C. After rolling, the sheet was quenched into water immediately, and then, reheated to 400 1C and held for 15 min as intermediate annealing for further rolling. To study the static recrystallization behavior, a small sample was taken from the quenched sheet after the 1st pass, and annealed at 400 1C for 5 min. The final sheet samples were also annealed for 30 min at 400 1C in order to compare the texture and mechanical properties of Mg–1Gd alloy and pure Mg. The rolling temperature of 300 1C was chosen to inhibit DRX, and therefore, the role of SRX on the microstructure and texture evolution can be precisely examined. Optical microscopy was used to analyze the microstructure of the sheet samples. Standard metallographic sample preparation
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techniques were used, and all specimens were etched using a solution of 3 g picric acid, 50 ml alcohol, 20 ml acetic acid and 20 ml water. The micro-texture and the macro-texture of rolled and annealed samples were examined in the plane containing the RD (Rolling Direction) and TD (Transverse Direction) using electron backscattering diffraction (EBSD) on a field emission gun scanning electron microscope (HITACHI SU-70, operated at 20 kV, EDAX/TSL EBSD system) and XRD technique using a Bruker D8 diffractometer with a Cu Kα source. Tensile tests were carried out at room temperature with the tensile direction parallel to the rolling direction of the sheet samples using a conventional testing machine ZWICK. The fracture surfaces of the tensile samples were observed using the same SEM.
3. Results 3.1. Microstructure and texture evolution during hot rolling and subsequent annealing Fig. 1a and b shows the macrostructures of the as-cast samples of pure Mg and Mg–1Gd alloy, respectively. Both samples show very large columnar grains with some millimeters in length and width, which indicates that the 1 wt% addition of Gd does not result in a grain refinement of the cast microstructure. This is in agreement with a previous study by Stanford et al. [6]. The optical microstructures of pure Mg and Mg–1Gd alloy in the as-rolled condition with a total thickness reduction of 80% are shown in Fig. 1c and d, respectively. The as-rolled pure Mg and Mg–1Gd alloy sheets both contain a large amount of twins in the microstructure. Some small grains are also observed which nucleate and grow during annealing between passes. After annealing for 30 min at 400 1C, SRX starts, and a fully recrystallized microstructure is formed as shown in Fig. 1e and f for pure Mg and Mg–1Gd
Fig. 1. Macrostructures of as-cast (a) pure Mg and (b) Mg–1Gd alloy. Microstructures in as-rolled condition: (c) pure Mg, (d) Mg–1Gd, and after annealing for 30 min at 400 1C: (e) pure Mg and (f) Mg–1Gd.
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alloy, respectively. It can be readily found that the grain size of pure Mg sheet is larger compared with the Mg–1Gd alloy, which suggests that the microstructure is significantly refined during hot rolling and subsequent annealing by the addition of 1 wt% Gd. Fig. 2 shows the {0002} pole figures of the pure Mg and Mg– 1Gd samples after 6-pass hot rolling and annealing for 30 min at 400 1C. As expected, a typical strong basal texture is formed in as-rolled pure Mg sheet, where the orientation distribution of the most basal poles is parallel to the sheet normal direction ND. After annealing, the texture strength is enhanced from 9.16 to 14.63. In contrast, the qualitative character of the Mg–1Gd alloy sheet texture is obviously distinct from typical basal texture for pure Mg sheet. The texture has a clearly broader distribution of basal poles, as compared with its distribution in the pure Mg. In addition, the Mg–1Gd alloy in as-rolled condition shows a split of basal pole intensity peak from ND toward RD, which is typically found for Mg alloys containing RE elements [12]. This is consistent with previous experimental observations and theoretical simulation in AZ31 alloys that the split of the maximum basal pole intensity about ND is due to the increased activity of pyramidal 〈c+a〉-slip [15]. After annealing, a much weaker sheet texture is obtained in Mg–1Gd alloy sheet compared with the as-rolled one.
Fig. 2. Recalculated {0002} pole figures of pure Mg and Mg–1Gd alloy sheet in as-rolled and annealed conditions.
Fig. 3a shows the microstructure of the as-rolled Mg–1Gd sample in the form of inverse pole figure (IPF) map, and their corresponding axis/angle misorientation distributions for special angle pairs with high density values. The IPF map of the as-rolled Mg–1Gd alloy sheet revealed a non-negligible proportion of twins in the deformed microstructures. The misorientation distributions in Fig. 3a indicate that the special boundaries correspond to {10 −11}–{10−12} secondary twin (381 〈11−20〉) (S-twin), {10−11} contraction twin (561 〈11−20〉) (C-twin) and {10−12} extension twin (861 〈11−20〉) (T-twin), respectively. An example of extension twin, contraction and secondary twin is illustrated in Fig. 3b. In Mg–1Gd alloy, there are some grains aligned favorably for extension twinning due to the weak basal texture. The accompanying {0002} pole figures illustrate the orientation relationship between extension twin, contraction twin, secondary twin and matrix. The accompanying {11−20} pole figures in Fig. 3b reveal that the matrix, the primary extension twin, contraction twin and the secondary twin share the same 〈11−20〉 rotation axis. Fig. 4 shows an EBSD Kernel average misorientation (KAM) map of the as-rolled pure Mg sheet with 80% thickness reduction. The KAM map is constructed based on the average misorientation between a measuring point and all its neighbors. Therefore, the local misorientation and strain energy can be clearly reflected in the KAM map. As can be seen from Fig. 4, the microstructure is heterogeneous, and can be divided into highly deformed regions with high KAM values and less deformed regions with low KAM values. And also, the twin boundaries are marked with different colors in Fig. 4. There are many contraction twins and secondary twins in the deformed matrix, but few extension twins are observed due to the strong basal texture of the sample. Fig. 5a shows the IPF map of pure Mg sample after annealing for 30 min at 400 1C, and a strong basal texture and a coarse microstructure can be observed. In contrast, a weak texture and fine-grained microstructure are observed in the Mg–1Gd alloy (Fig. 5b). The corresponding inverse pole figures in the RD of pure Mg and Mg–1Gd alloy samples are shown in Fig. 5c and d, respectively. Interestingly, there is a distinct difference in the texture component between pure Mg and Mg–1Gd alloy. The 〈10 −10〉 and 〈11−20〉 fiber texture components are both observed in annealed pure Mg sheet, which is typically found in extruded AZ31 Mg alloys. In contrast, only 〈10−10〉 texture component is observed for Mg–1Gd alloy sheet. Although the inverse pole figures show a
Fig. 3. (a) IPF map of the as-rolled Mg–1Gd sample and their corresponding axis/angle misorientation distributions for special angle pairs and (b) examples of extension, contraction and secondary twins. The accompanying {0002} pole figures illustrate the crystallographic arrangement of the twins and the matrix. The accompanying {11−20} pole figures indicate a common rotation axis during both twinning deformation.
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Fig. 4. Kernel average misorientation map constructed using the orientation relationship with 3rd neighbors, of 80% hot rolled pure Mg sample. Twin boundaries are marked with different colors: red lines correspond to extension twin, green lines to contraction twin and blue lines to secondary twin. (For interpretation of references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 5. IPF map of (a) pure Mg, (b) Mg–1Gd alloy sheet after annealing for 30 min at 400 1C; corresponding inverse pole figures in the RD of (c) pure Mg and (d) Mg–1Gd alloy sheet; (e) corresponding misorientation angle distribution and (f) grain size distribution. Note the higher magnification in (b).
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similar peak intensity in pure Mg and Mg–1Gd alloy, the basal texture is weakened significantly by Gd addition. Fig. 5e shows the number of fractions of misorientation angles for both the annealed sheets. In pure Mg, a distribution peak around 301 is observed, which is related to (0002) 〈10−10〉 and (0002) 〈11−20〉 texture component. The majority of misorientation angles are below 451. However, the misorientation distribution appears quite different in the Mg–1Gd alloy sheet, and large misorientation angles are more significant with a distribution peak around 401, which suggests a more uniform misorientation distribution than pure Mg. This is correlated to the weak texture with the addition of Gd. Fig. 5f shows the grain size distribution in the pure Mg and Mg– 1Gd alloy sheets, as determined by OIM analysis. As can be seen, in Mg–1Gd alloy, the grain size ranges between 6 μm and 29 μm, and the average grain size is about 18 μm. However, a large variety in grain size from 15 μm to 133 μm was observed in pure Mg. Therefore, a homogeneous and fine-grained microstructure with a weak texture is obtained by the addition of Gd. 3.2. Mechanical properties The typical stress–strain curves of the hot rolled pure Mg and Mg–1Gd alloy after annealing for 30 min at 400 1C along the RD are shown in Fig. 6. The ultimate tensile strength (UTS), yield strength (YS) and elongation to failure of the annealed pure Mg sheet are 167.6 MPa, 69.5 MPa and 5.4%, respectively. In contrast, the room-temperature UTS and YS of annealed Mg–1Gd alloy sheet are 193.5 MPa and 129.9 MPa respectively, with the elongation dramatically increasing to 15.6%, indicating an enhanced combination of strength and ductility. Stanford et al. [6] have reported that Gd had effective solid solution strengthening effect on Mg sheet and the UTS and YS both increased with Gd concentration, but were basically unchanged above 1 wt% Gd. In addition, Mg–1Gd alloy shows a much finer grain size than pure Mg (Fig. 5f), which contributes to UTS and YS. Therefore, the increase in UTS and YS by the addition of Gd is mainly ascribed to the solid solution strengthening effect and the grain refinement. The tensile fractograph of the hot rolled pure Mg and Mg–1Gd sheet samples after annealing is shown in Fig. 7a and b, respectively. The fracture surface of pure Mg shows cleavage facets, which contain tear ridges, and a small number of dimples. The fracture mode of pure Mg can therefore be categorized as a quasicleavage fracture. In the Mg–1Gd alloy, however, a ductile fracture surface characterized by the presence of elongated dimples is observed.
Fig. 6. Typical stress–strain curves of the hot rolled pure Mg and Mg–1Gd alloy after annealing for 30 min at 400 1C in the tensile directions of RD.
Fig. 7. Typical SEM images showing the fractograph of the tensile tested pure Mg (a) and Mg–1Gd alloy (b) along the RD.
4. Discussion 4.1. Texture weakening mechanism by Gd addition In the present work, DRX is inhibited during hot rolling at a low temperature of 300 1C in Mg–1Gd alloy. Thus, SRX is the most important mechanism during annealing at 400 1C. Therefore, it seems reasonable to suggest that the texture weakening is related to the SRX. The recrystallization behavior of pure Mg, which is similar to conventional Mg–Al based alloys such as AZ31, has already been discussed in the literature. During the annealing process of the as-rolled pure Mg sample, nucleation occurs at preferred sites such as contraction and secondary twins with a high stored energy. Moreover, rapid recovery may occur in the less deformed matrix grains with low KAM values, which contributes to the persistence of the strong (0002) basal texture [16]. Here, more emphasis is placed on the Mg–1Gd alloy and particularly on the recrystallization texture. Fig. 8a shows the optical microstructure of Mg–1Gd sample in as-rolled condition with 20% thickness reduction. The as-rolled microstructure contains a large amount of twins, and no sign of DRX is found in the sample. Deformation accumulates in the contraction, secondary twins, and shear bands evolved from them [17,18]. Upon annealing of the rolled Mg–1Gd alloy at 400 1C for 5 min, such regions of high dislocation density are favorable sites for nucleation, as shown in Fig. 8b. In order to investigate the micro-texture of recrystallized grains, the recrystallized region is presented in more detail in
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Fig. 8. Optical microstructure of Mg–1Gd sample in (a) as-rolled condition with 20% thickness reduction and (b) after annealing for 5 min at 400 1C.
Fig. 9. (a) IPF map showing onset of SRX at bands and twins after annealing of Mg–1Gd at 400 1C: for 5 min; (b) corresponding KAM map marked with the twin boundaries; (c) (0002) pole figure of unrecrystallized grains illustrating the crystallographic arrangement of the different twins and the matrix grain and (d) (0002) pole figure of recrystallized grains.
Fig. 9a in term of an IPF map. It can be found that the distribution of recrystallization grains was very heterogeneous due to the heterogeneity of the as-deformed microstructure. These observations show that SRX takes place at twins and shear bands. Fig. 9b shows the corresponding KAM map. Newly recrystallized grains are easily detected by their low KAM values. Different types of
twin boundaries are marked in different colors, and a large amount of primary {10−12} extension twins can be seen in Fig. 9b. In addition, a few {10−11} contraction twins and some {10−11}–{10−12} secondary twins are also presented in the as-rolled sample after annealing for 5 min at 400 1C. Fig. 9c displays (0002) pole figure of unrecrystallized grains illustrating
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Fig. 10. (0001) 〈11−20〉 basal slip Schmid factor distribution maps before the tensile test: (a) pure Mg; (b) Mg–1Gd; and corresponding fraction (c) before and (d) after tensile test. (For interpretation of references to color in this figure, the reader is referred to the web version of this article.)
the crystallographic arrangement of the different twins and the matrix grain. T-twin A corresponds to region A, and T-twin B to region B (Fig. 9a). Thus, regions A and B are two different twin variants. According to the work by Nave and Barnett [19], three different misorientation relationships are possible between different twin variants because of the six {10−12} planes available for twinning, i.e. 7.41〈1−210〉, 60.01〈10−10〉, 60.41〈8−1−70〉. In the current work, the misorientation relationship between the two variants is a 551〈10−10〉 relationship (Fig. 9a), which is close to 60.01〈10−10〉 relationship. The difference in rotation angle may result from lattice rotation due to deformation. In the early stage of rolling, primary {10−12} extension twinning can be easily activated. Under further deformation, most of the matrix is consumed and reoriented by extension twinning, which contributes to the formation of basal texture. Moreover, contraction and secondary twins develop in the primary extension twinning region in order to accommodate deformation. This phenomenon was also reported by Jiang et al. [20], who found that a high density of secondary {10−12} twins form on the primary {10−12} extension twins at a high compression strain level in AZ31 alloy. Fig. 9d shows the orientation of the recrystallized grains in terms of (0002) pole figure. Apparently, the new grains nucleated at bands and secondary twins show a much wider spread of orientations than that of deformed parent grains and twins, leading to the weakening of the basal texture. As can be seen from Fig. 1c and d, there is more twinning in the as-rolled Mg–1Gd alloy compared with pure Mg. This result is consistent with the work by Hantzsche et al. [12], who found that the as-rolled microstructure of Mg alloys with higher RE contents contains a higher amount of twins, especially contraction and secondary
twins, than those with lower RE content. In addition, Muzyk et al. [21] calculated the generalized stacking fault energies of Mg alloys using density functional theory, and demonstrated that alloying with RE increases the possibility of twin formation during deformation of Mg alloys. Therefore, it is reasonable to suggest that alloying with Gd would increase the tendency for contraction and secondary twinning, which contributes to the texture weakening after annealing. Solute atoms with a larger atomic size than solvent atoms will tend to segregate to the grain boundaries in order to reduce the elastic strain energy [22,23]. Gd is significantly larger in atomic size than Mg, and it is confirmed by atom probe tomography (APT) that Gd solute atoms segregate strongly to grain boundaries [11]. During the growth of recrystallized grains, Gd solute atoms will move along with the boundary via slow diffusion, and thus reduce the grain boundary mobility. On the other hand, there are indications in the literature that with increasing segregation, the energy of special grain boundaries tends towards that of random boundaries [24]. This could eventually reduce the driving force for the preferred migration of special grain boundary, such as 301 〈0001〉 boundary in the conventional Mg alloys. Consequently, it is responsible for the texture weakening after annealing. Actually, the microstructure and texture evolution during static recrystallization of conventional Mg alloys without RE have been investigated extensively. Yi [16] and Li [18] have found that in the early stage of annealing, the orientation of recrystallized grains formed at secondary and contraction twins is also evenly distributed in pure Mg and AZ31 alloy. Farzadfar et al. [25] found that in Mg–2.9Zn alloy, new grains recrystallized in bands exhibit orientations close to the ones associated with twins, and contribute
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thus to texture weakening. With the increase in grain size during coarsening, the maximum intensity of basal pole figures rises linearly. They ascribed this texture strengthening to the consumption of small recrystallized grains by larger ones. As can be seen from Fig. 5c, besides {0002} 〈10−10〉 texture component, a strong {0002} 〈11−20〉 component is observed in as-rolled pure Mg after annealing for 30 min at 400 1C. This is because grain growth is controlled mainly by the preferred growth of grains with the 〈11 −20〉 component during annealing [26], and it could be speculated that, upon further annealing, the {0002} 〈11−20〉 component strengthens gradually and becomes the main texture component. However, in Mg–1Gd alloy, only 〈10−10〉 texture component is observed, which indicates that the growth of grains with 〈11−20〉 component is restricted by solute segregation of Gd. From the above analysis, it can be concluded that the recrystallization texture modification by the addition of Gd mainly takes place during the growth step. The growth of recrystallized grains of Mg–1Gd alloy sample is inhibited by solute segregation of Gd, which is a major contributor to the texture weakening. However, the role of deformation behavior in the recrystallization texture evolution cannot be neglected entirely. As mentioned in Section 3.1, the split of basal pole intensity peak from ND toward RD is a typical indication of pyramidal 〈c+a〉-slip. It has been reported that the addition of Gd to Mg alters the bond energy between Gd atoms and Mg atoms [27], increasing the probability of non-basal slip. Sandlöbes et al. [17,28] proposed that an addition of Y to the Mg alloy can influence the stacking fault energy (SFE) on either basal or pyramidal planes and reduce the differences of critical resolved shear stresses (CRSS) for the corresponding deformation mechanisms, and therefore promote the activity of pyramidal 〈c+a〉-slip. Therefore, in the current work, the activity of pyramidal 〈c+a〉-slip of Mg sheet during rolling could be enhanced by Gd additions, which results in the split of the maximum basal pole intensity. Thus, the texture strength of the as-rolled Mg–1Gd alloy sheet is decreased significantly (see Fig. 2). After annealing for 30 min at 400 1C, the recrystallized grains show a much wider spread of orientations than the basal axis orientations of deformed matrix and the texture strength is decreased further. Therefore, the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling by Gd addition is also a non-negligible contributor to the texture weakening. 4.2. Effect of Gd on the ductility of Mg sheet As can be seen from Fig. 6, the room-temperature ductility of Mg sheet is significantly improved by the addition of Gd. In order to fully understand the effect of Gd addition on the ductility of hot rolled Mg sheet, the distribution of the (0001) 〈11−20〉 basal slip Schmid factor is investigated by EBSD analysis. The results in Fig. 10a and b show that the blue grains usually have their basal planes parallel to the tensile direction, as shown by hexagonal unit cell. The blue grains have a relatively lower (0001) 〈11−20〉 basal slip Schmid factor and is unfavorable for basal slip. However, the red grains usually have their basal planes inclined about 451 to the tensile planes as shown by hexagonal unit cell. Thus, the red grains have a higher (0001) 〈11−20〉 basal slip Schmid factor, which is favorable for basal slip. Compared to pure Mg, a large number of red grains are observed in Mg–1Gd alloy, which is correlated to the weak texture by Gd addition. Fig. 10c and d shows the fraction distribution of the Schmid factor for (0001) 〈11−20〉 basal slip along the RD of both sheets before and after tensile test, respectively. It can be seen that about 35% grains in pure Mg show low Schmid factor values less than 0.1. The fraction of grains gradually decreases with the increasing Schmid factor. The highest Schmid factor ( 40.4) corresponds to the lowest fraction of 0.04. In contrast, there is a more uniform
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Fig. 11. Variation of ln s against ln ε for pure Mg and Mg–1Gd alloy.
distribution of the Schmid factor for (0001) 〈11−20〉 basal slip along the RD of Mg–1Gd alloy sheet. The fraction of grains with a high Schmid factor (40.4) is greater than 0.18. In Mg–1Gd alloy, the proportion of deforming grains by basal slip is markedly higher than in pure Mg, due to the weaker basal texture component. After the tensile test, the fraction of grains with a high Schmid factor (40.4) for basal slip is decreased for both materials (Fig. 10d), which means that basal slip is the dominant deformation mechanism during tensile test. Therefore, we propose here that the weak texture is a major contributor to the improved ductility. This result is consistent with the work by Stanford et al. [6], who found that a significant increase in ductility is achieved due to texture weakening resulting from Gd addition. Moreover, from Fig. 5f, the annealed microstructure is significantly refined by the addition of Gd due to the solute segregation effect at grain boundaries. Therefore, grain refinement would be another important contributor to the enhanced ductility. In addition, strain hardening behavior is another metallurgical factor that may contribute to ductility, which is characterized by a power law relationship, s ¼Kεn, where s and ε are the true stress and strain, K is the strength coefficient and n is the strain hardening exponent [15]. A high value of n would inhibit inhomogeneous deformation, and thus, improve formability and ductility. Fig. 11 shows the variation of ln s against ln ε of tensile data for the pure Mg and Mg–1Gd alloy along the RD. The obtained n values during tension along the RD are 0.326 and 0.172 for pure Mg and Mg–1Gd, respectively. Apparently, the strain hardening exponent n value is decreased by the addition of Gd, which indicates that this parameter does not contribute to the increased ductility obtained in Mg–1Gd alloys.
5. Conclusions The microstructure, texture and tensile ductility of Mg–1Gd alloy and pure Mg sheet were investigated following rolling and subsequent isothermal annealing. The addition of Gd weakens the basal texture in both the as-hot rolled and annealed conditions. This is related to the enhanced activity of pyramidal 〈c+a〉-slip during hot rolling by Gd addition and the suppressed grain boundaries migration due to solute segregation of Gd at grain boundary during annealing. Pure Mg sheet shows a strong {0002} 〈11−20〉 texture component due to the preferred growth of grains with the 〈11−20〉 component. However, Gd solute segregation at grain boundary could inhibit the preferred growth of grains with the 〈11−20〉 component, which leads to a weak basal texture and a
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single 〈10−10〉 texture component in Mg–1Gd alloy sheet. A large amount of secondary twins and deformation bands form during hot rolling of Mg–1Gd sheet, and may serve as favorable nucleation sites for SRX during annealing. The recrystallized grains at bands/twins in Mg–1Gd alloy display a relatively wide spread of orientations. The room temperature ductility is significantly improved by the addition of Gd, which is mainly attributed to the texture weakening and grain refinement. Acknowledgments The authors gratefully acknowledge the support of the National High-Tech R&D Program of China (Grant no. 2011BAE22B06), the National Natural Sciences Foundation of China (Grant nos. 50901044 and 51271118) and the Shanghai Rising-Star Program (B type) (Grant no. 12QB1403300). References [1] [2] [3] [4]
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