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New Mg–Al based alloy sheet with good room-temperature stretch formability and tensile properties
Scripta Materialia 180 (2020) 16–22
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
New Mg–Al based alloy sheet with good room-temperature stretch formability and tensile properties T. Nakata a,∗, C. Xu b, H. Ohashi a, Y. Yoshida c, K. Yoshida c, S. Kamado a a
Nagaoka University of Technology, 1603-1, Kamitomioka, Nagaoka 940-2188, Japan School of Materials Science and Engineering Harbin Institute of Technology, Harbin 150001, China c Sumitomo Electric Industries, Ltd., 1-1-1 Koyakita, Itami 664-0016, Japan b
a r t i c l e
i n f o
Article history: Received 5 October 2019 Revised 26 December 2019 Accepted 13 January 2020
Keyword: Magnesium Rolling Room temperature formability Tensile property Texture
a b s t r a c t We have successfully developed a Mg–3Al–0.4Mn (wt%) alloy sheet with good room-temperature formability. The alloy sheet shows a large Index Erichsen value of 8.2 mm due to its unique texture feature where the most of the grains have circular distribution of basal poles away from the normal direction of the sheet. The fine grain structure also leads to moderate strengths and large ductility with in-plane isotropic tensile property. These mechanical properties could be realized via industry viable rolling processing, thus the alloy sheet is highly possible to broaden the applications of Mg alloys in automotive industries. © 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
Mg–3Al–1Zn (AZ31, wt%) alloy is the most common wrought Mg alloy because of its moderate strengths, weldability, and corrosion resistance [1]. However, the rolled AZ31 alloy sheets usually show poor room-temperature (RT) formability [2–4], hindering the wide applications in automobiles. The limited formability at RT arises from a strong basal texture with basal poles parallel to the normal direction (ND) of the sheets, which is formed during the rolling processing [4–6]. Numerous attempts have been carried out to control the basal texture and improve the RT formability of commercial AZ31 alloy sheets [7–10]. High-temperature rolling is the most effective way to weaken the basal texture of the rolled sheets, and Huang et al. have successfully increased the Index Erichsen (I.E.) value of a commerial AZ31 alloy sheet from 3.7 mm to 9.5 mm [7]. However, this method will be undesirable in large-scale manufacturing process due to limited capacity of production facilities or high production cost. So texture weakening needs to be achieved via industry viable rolling processing. Mg-3Al-0.4Mn (AM30, wt%) alloy is a new wrought Mg alloy with a good balance of strengths and ductility [11]. Also, the AM30 alloy has slightly better extrudability and deformability at elevated temperatures than AZ31 alloy [11,12], implying that AM30 alloy has good rolling ability. If AM30 alloy also exhibits moderate RT formability, the AM30 alloy may become a candidate as a sheet alloy. In this work, we firstly report a rolled AM30 alloy sheet ∗
Corresponding author. E-mail address: [email protected] (T. Nakata).
with good combination of RT formability and tensile properties via industry viable rolling processing consisting of homogenization, conventional warm rolling, and annealing. Mg–3Al–0.8Zn–0.4Mn and Mg–3Al–0.4Mn (wt%, AZ31 and AM30) alloy ingots were prepared by a direct chill casting. The ingots were homogenized at 415 °C for 2 h and 500 °C for 12 h in an electronic furnace under an Ar atmosphere, followed by water quenching. The homogenized ingots were machined to make slabs with 10 mm in thickness and 120 mm in width. They were soaked at 220 °C for 30 min in an electronic furnace and continuously rolled with a rolling reduction of 20%/pass. The rolling was repeated 10 times without reheating to obtain sheets with 1 mm in thickness. During the rolling, the temperature and speed of the rollers were fixed at 220 °C and 5 m/min, respectively. The time between each rolling pass was within 10 s. The rolled sheets were annealed at 250 °C for 4 h, followed by water quenching. Stretch formability of the annealed sheets were evaluated by an Erichsen cupping test at RT using square shape specimens with 60 × 60 mm2 . The punch diameter and speed were 20 mm and 6 mm/min, respectively. Tensile tests were also conducted using the annealed specimens with gauge length, width, and thickness of 20, 4, and 1 mm at an initial strain rate of 10−3 s–1 and RT along the rolling and transverse directions (RD and TD). The Erichsen cupping test was repeated five times and tensile tests were repeated three times. A scanning electron microscope (JEOL, JSM-70 0 0F) equipped with TSL electron backscattered diffraction (EBSD) apparatus was used to characterize the second phases,
https://doi.org/10.1016/j.scriptamat.2020.01.015 1359-6462/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22
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Table 1 0.2% proof stress, ultimate tensile strength, and elongation to failure along rolling and transverse directions (RD and TD) of AZ31 and AM30 alloy sheets annealed for 250 °C for 4 h. Alloy
AZ31 AM30
Stretched along RD
Stretched along TD
0.2% proof stress [MPa]
Ultimate tensile strength [MPa]
Elongation to failure (%)
0.2% proof stress [MPa]
Ultimate tensile strength [MPa]
155±1 153±3
256±2 248±5
28±1 28±2
173±3 171±3
265±1 259±4
Elongation to failure (%) 25±1 25±2
Table 2 (0 0 01) [11-2¯ 0] Schmid factors (Schmid factors for basal slips) along rolling and transverse directions (RD and TD) of the AZ31 and AM30 alloy sheets annealed for 250 °C for 4 h. Alloy
AZ31 AM30
Schmid factor for basal slips Stretched along RD
Stretched along TD
0.27 0.28
0.22 0.26
grain sizes, and textures of the annealed samples. Microstructural development during the stretch forming in both AZ31 and AM30 samples was also characterized by the EBSD. The Erichsen cupping test was stopped after the punch stroke reaches to 2 mm and 4 mm. The deformed specimens were cut at the center along the RD, and the surface was polished by a Cross Section Polisher (JEOL, SM-09020CP). The EBSD scan was conducted on the top and bottom surfaces of the deformed specimens. The measured area of this analysis was 0.2 × 0.2 mm2 . To identify active dislocations, in-grain misorientation axes (IGMA) distribution was investigated using material point pairs having misorientation angles from 2° to 5° [13]. A quasi-in-situ EBSD analysis was also carried out to understand the microstructural development during the annealing. In the quasi-in-situ EBSD analysis, the EBSD scan was conducted firstly on the region of interest (ROI) of the as-rolled samples. The scanned samples were subjected to an annealing at 250 °C for 10, 30, 120 and 300 s in an infrared vacuum furnace (THERMO RIKO CO., LTD., IVF198W), and the EBSD scan was conducted again on the same ROI. Grains with grain orientation spread (GOS) less than 1° were regarded as recrystallized grains. The measured area for the quasi-in-situ EBSD analysis was about 0.2 mm2 . All microstructural characterization was done on the RD-ND plane, and the EBSD data was analyzed from the RD-TD plane. Fig. 1 shows tensile stress strain curves stretched along the RD and TD, appearances of the fractured specimens after the Erichsen cupping test, inverse pole figure maps, (0 0 01) pole figures, profile of the (0 0 01) pole density measured along the RD, and secondary electron images of the (a) AZ31 and (b) AM30 samples. Their average 0.2% proof stresses (PS), ultimate tensile strengths (UTS), and elongation to failure (EF) are summarized in Table 1 with the standard deviations, and Table 2 summarizes their (0 0 01) [11-2¯ 0] Schmid factors (Schmid factors for basal slips). Inset values in the snapshots of the fractured Erichsen cupping samples are average I.E. values with standard deviations. The PS and EF of both samples are almost the same. The AZ31 sample exhibits slightly higher UTS than those in the AM30 sample, which may be due to the solid-solution of Zn [14]. Though the AM30 sample stretched along the RD shows lower strengths and higher EF than those of the TD specimens, the difference is very small. The anisotropy is also small in the AZ31 sample; however, the AZ31 sample shows poor I.E. value of 4.7 mm. The AM30 sample shows significantly large I.E. value of 8.2 mm. This is similar to that of newly designed Mg–Al–Ca–Zn–Mn alloy sheets [15,16]. Both samples show fully recrystallized microstructure with average grain sizes of ~8 μm. The AZ31 sample forms RD-split texture as previously reported AZ31
Fig. 1. Tensile stress strain curves stretched along rolling and transverse directions (RD and TD), appearances of fractured specimens after Erichsen cupping test, inverse pole figure maps, (0 0 01) pole figures, profile of (0 0 01) pole density measured along the RD, and secondary electron images of (a) AZ31 and (b) AM30 alloy sheets. Inset I.E. values represent average Index Erichsen values with standard deviations. Dash lines in the profile of (0 0 01) pole density indicate the region of strong alignment of (0 0 01) poles to the ND.
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Fig. 2. Inverse pole figure maps, (0 0 01) pole figures of, and distributions in-grain misorientation axes (IGMA) of (a, c, e, f) AZ31 and (b, d, g, h) AM30 alloy sheets after interrupted Erichsen cupping test. Note that these images were taken from top surfaces of deformed specimens, and the inset values in the pole figures represent the maximum intensities of the poles.
alloy sheets [17,18], and the AM30 sample has similar preferential tilt of the basal poles to the RD. However, in the AM30 sample, most of the grains have circular distribution of basal poles away from the ND. Schmid factors for basal slips in the AM30 sample are higher than those in the AZ31 sample, and the AM30 sample has similar Schmid factors for basal slips along both RD and TD. These textural features indicate that, in the AM30 sample, large number of grains are favorably oriented for basal slips during RT stretch forming. Also, they are beneficial to reduce in-plane yield anisotropy. Both samples contain second phases. These phases may be Al8 Mn5 phases [11], and their sizes and fractions in both samples are almost the same. Fig. 2 shows inverse pole figure maps, (0 0 01) pole figures of the (a, b) AZ31 and (c, d) AM30 samples after the interrupted
Erichsen cupping test. Distributions of the IGMA in the (e, f) AZ31 and (g, h) AM30 samples are also shown in Fig. 2, and Table 3 summarizes fractions of low angle grain boundaries (LAGBs) having misorientation axis of [0 0 01], [101¯ 0], or [112¯ 0]. Note that misorientation angles of 2°−5° are considered. In both samples, the angular distribution of the basal poles gets narrower toward the ND, and the maximum intensities of the basal poles gradually increase with the progress of the deformation. Many IGMA are lying along [uv.0], and Table 3 shows that majority of LAGBs have misorientation axis of [10-1¯ 0] or [11-2¯ 0], suggesting that basal slips are the most dominant deformation mechanism during the RT stretch forming in both samples [13,19,20]. Some IGMA are lying along [0 0 01], and the fractions of LAGBs with misorientation axis of [0 0 01] are similar in both samples. This means that
T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22
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Fig. 3. Inverse pole figure maps, (0 0 01) pole figures, and misorientation angle distributions of (a, c, e) AZ31 and (b, d, f) AM30 alloy sheets after interrupted Erichsen cupping test taken from bottom surfaces of deformed specimens. The inset values in the pole figures represent the maximum intensities of the poles.
prismatic slips are also active and the activity is almost the same in both samples [13,20]. It is also interesting to note that intense shear bands are formed in the AZ31 sample as indicated by white dot lines, while such shear bands have not been observed in the AM30 sample. Fig. 3 shows inverse pole figure maps, (0 0 01) pole figures, and misorientation angle distributions of the (a, c, e) AZ31 and (b, d, f) AM30 samples after the interrupted Erichsen cupping test. Note that they were obtained from bottom surfaces of the deformed specimens. The inset values in the pole figures represent the maximum intensities of the poles. After the punch stroke reaches to 2 mm, distinct color changes occur within the grains,
and large fraction of the misorientation angles near 86° has been appeared, indicating that the formation of tensile twinning [21]. The AM30 sample has higher fraction of the misorientation angles near 86°, and the (0 0 01) pole figure shows clear change. These results suggest that tensile twinning is easily occurred in the AM30 sample. After the punch stroke reaches to 4 mm, the fraction of the misorientation angles near 86° decreases significantly, implying the occurrences of de-twinning [21]. In this work, we have found that AM30 alloy sheet exhibits excellent RT stretch formability and good tensile properties via industry viable rolling processing. The AM30 alloy sheet exhibits
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T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22
Fig. 4. Inverse pole figures and (0 0 01) pole figures of as-rolled and annealed (a) AZ31 and (b) AM30 alloy sheets. The inverse pole figure maps and the pole figures extracted from recrystallized grains (GOS < 1°) are also displayed in the right hand side, and the inset values in the pole figures represent maximum intensity of the (0 0 01) poles. Average sizes of recrystallized grains, drec , and percentages of the recrystallized grains with strong basal texture component, f0°−10° , (recrystallized grains having their basal poles tilting 0°−10° from the normal direction of the sheets) in the AZ31 and AM30 alloy sheets are also summarized in the (c) and (d). RD and TD in the figure represent rolling and transverse directions of the sheet.
large I.E. value of 8.2 mm, and the sheet shows PS and EF of 153 MPa and 28% along the RD. Also, the in-plane yield anisotropy is small, the sheet shows PS of 171 MPa along the TD with good EF of 25%. The balance of these properties is higher than the
recently developed RT formable Mg alloy sheets [7,22–24]. The PS of the AM30 alloy sheet is still lower than Mg–Al–Ca–Zn–Mn alloy sheets [15,16]; however, tensile properties can be improved by means of twin-roll casting [16]. Because the twin-roll casting is
T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22 Table 3 Fractions of low angle grain boundaries having misorientation axis of [0 0 01], [101¯ 0], or [11-2¯ 0] in AZ31 and AM30 alloy sheets after interrupted Erichsen cupping test with punch stroke of 2 mm and 4 mm. Note that misorientation angles of 2°−5° are considered for the calculation. Alloy
AZ31 AM30
Punch stroke [mm]
2 4 2 4
Misorientation axis [0 0 01]
[10-1¯ 0]
[11-2¯ 0]
3.6% 2.9% 2.8% 4.7%
37% 33% 38% 31%
35% 33% 37% 31%
becoming commercial processing for producing low-cost Mg alloy sheets [25], the AM30 alloy may become a promising sheet alloy in automotive industries. The moderate strengths and large ductility could be realized by the fine grain structure [26–28]. The fine grain structure along with unique texture feature, where the most of the grains have circular distribution of basal poles away from the ND, also contributes to small in-plane yield anisotropy due to similar activity of basal slips along the RD and TD (Table 2). The AZ31 alloy sheet also shows small in-plane yield anisotropy despite the difference of the Schmid factors along the RD and TD. This may be attributed to the higher solute strengthening of Zn against basal slips than prismatic slips [14]. Although both AZ31 and AM30 alloy sheets show similar tensile properties, the AM30 alloy sheet exhibits superior RT stretch formability with large I.E. value of 8.2 mm. In the AM30 alloy sheet, large number of grains are favorably oriented for basal slips during RT stretch forming due to unique texture feature, and this is important to accommodate thickness strain of the sheets [15]. As shown in Fig. 3, basal slips are dominant in both alloy sheets; therefore, it can be inferred that large I.E. value of the AM30 alloy sheet is mainly attributed to its unique texture feature. It is worth noting that tensile twinning is more actively occurred and the formation of shear bands is suppressed in the AM30 alloy sheet. It is reported that activation of twinning and de-twinning will contribute to the enhanced RT stretch formability to some extent [26,27], and formation of shear bands decreases RT stretch formability [20]. They suggest that enhanced activities of basal slips and twinning/de-twinning, and suppression of shear bands formation contribute to the excellent RT stretch formability in the AM30 alloy sheet. In-depth understanding of the effect of alloying element, microstructures, and texture on deformation behavior will be critical for further improvement of RT formability. To understand the formation of unique texture feature in the AM30 sample, quisi-in-situ EBSD was done. Fig. 4 shows inverse pole figures and (0 0 01) pole figures of the as-rolled and annealed (a) AZ31 and (b) AM30 samples. The inverse pole figure maps and the pole figures extracted from recrystallized grains are also displayed. Average recrystallized grain sizes, drec , and percentages of the recrystallized grains with strong basal texture component, f0°−10°, (recrystallized grains having their basal poles tilting 0°−10° from the ND) in both samples are also summarized in Fig. 4(c) and (d). In the as-rolled condition, both samples consist of deformed grains, shear bands, and small fraction of recrystallized grains. The areal fraction of the recrystallized grains and the drec in the AZ31 sample are 8% and 1.7 μm, while those in the AM30 sample are 3% and 0.9 μm, which means that the dynamic recrystallization and/or grain growth during the rolling is strongly retarded in the AM30 sample. The suppression of the dynamic recrystallization and/or grain growth may come from high melting point due to the absence of Zn [28]. The texture features of the as-rolled AZ31 and AM30 samples are almost the same; the f0°−10° in both samples are very similar. After the short-time annealing, recrystal-
21
lized grains with RD-split texture feature mainly form at the shear bands. Further annealing mainly causes grain growth and decreases the fraction of the recrystallized grains with strong basal texture component. During the annealing, the AZ31 sample keeps the RD-split texture feature of the statically recrystallized grains and strong basal texture component of the dynamically recrystallized grains. The AM30 sample also shows similar texture development; however, the recrystallized grains start to show circular distribution of basal poles, and the f0°−10° decreases significantly by the progress of the static recrystallization. Previous work has shown that basal poles of dynamically recrystallized grains gradually align to ND during a hot rolling. During the subsequent heat-treatment, the dynamically recrystallized grains grow and thus strong basal texture develops [29]. This result implies that suppression of dynamic recrystallization during rolling process is one of the important factors to produce Mg sheets with weakened basal texture. As shown in Fig. 4(a) and (b), the fraction of the recrystallized grains in the AM30 sample is smaller than the AZ31 sample. This should be one reason why the AM30 sample forms different texture feature compared to the AZ31 sample. Detailed observation of active deformation modes during the rolling needs to be done to fully understand the unique texture feature in the AM30 alloy sheet. In summary, we have successfully developed a Mg–3Al–0.4Mn (AM30) alloy sheet with excellent RT formability, moderate strengths, and large ductility along with isotropic tensile properties. The newly developed AM30 alloy sheet exhibits good balance of RT formability and tensile properties due to its fine grain structure and unique texture feature, where the basal poles of most grains have circular distribution away from the ND. Such unique texture feature and good properties could be obtained using only commercial rolling processing such as homogenization, rolling at warm temperature, and annealing; therefore, the AM30 alloy will become a promising sheet alloy in automotive industries. Understanding the unique texture formation will be an important work to optimize processing conditions and alloy compositions for the RT formable Mg alloy sheets. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgment This work was supported by JSPS KAKENHI Grant Numbers JP19K15321, JP18H03837, and Advanced Low Carbon Technology Research and Development Program (ALCA), 12102886. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2020.01. 015. References [1] E.F. Emley, Principle of Mangesium Technology, Pergamon Press, Oxford, 1966. [2] Y. Chino, H. Iwasaki, M. Mabuchi, H.T. Jeong, Mater. Sci. Eng.: A 466 (2007) 90–95. [3] B.-C. Suh, J.H. Kim, J.H. Bae, J.H. Hwang, M.-S. Shim, N.J. Kim, Acta Mater. 124 (2017) 268–279. [4] D.H. Kang, D.-W. Kim, S. Kim, G.T. Bae, K.H. Kim, N.J. Kim, Scr. Mater. 61 (2009) 768–771. [5] A. Styczynski, Ch. Hartig, J. Bohlen, D. Letzig, Scr. Mater. 50 (2004) 943–947. [6] Q. Jin, S.-Y. Shim, S.-G. Lim, Scr. Mater. 55 (2006) 843–846. [7] X. Huang, K. Suzuki, Y. Chino, Scr. Mater. 61 (2009) 445–448. [8] M. Eddahbi, J.A. del Valle, M.T. Pérez-Prado, O.A. Ruano, Mater. Sci. Eng.: A 410-411 (2005) 308–311.
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Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
New Mg–Al based alloy sheet with good room-temperature stretch formability and tensile properties T. Nakata a,∗, C. Xu b, H. Ohashi a, Y. Yoshida c, K. Yoshida c, S. Kamado a a
Nagaoka University of Technology, 1603-1, Kamitomioka, Nagaoka 940-2188, Japan School of Materials Science and Engineering Harbin Institute of Technology, Harbin 150001, China c Sumitomo Electric Industries, Ltd., 1-1-1 Koyakita, Itami 664-0016, Japan b
a r t i c l e
i n f o
Article history: Received 5 October 2019 Revised 26 December 2019 Accepted 13 January 2020
Keyword: Magnesium Rolling Room temperature formability Tensile property Texture
a b s t r a c t We have successfully developed a Mg–3Al–0.4Mn (wt%) alloy sheet with good room-temperature formability. The alloy sheet shows a large Index Erichsen value of 8.2 mm due to its unique texture feature where the most of the grains have circular distribution of basal poles away from the normal direction of the sheet. The fine grain structure also leads to moderate strengths and large ductility with in-plane isotropic tensile property. These mechanical properties could be realized via industry viable rolling processing, thus the alloy sheet is highly possible to broaden the applications of Mg alloys in automotive industries. © 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
Mg–3Al–1Zn (AZ31, wt%) alloy is the most common wrought Mg alloy because of its moderate strengths, weldability, and corrosion resistance [1]. However, the rolled AZ31 alloy sheets usually show poor room-temperature (RT) formability [2–4], hindering the wide applications in automobiles. The limited formability at RT arises from a strong basal texture with basal poles parallel to the normal direction (ND) of the sheets, which is formed during the rolling processing [4–6]. Numerous attempts have been carried out to control the basal texture and improve the RT formability of commercial AZ31 alloy sheets [7–10]. High-temperature rolling is the most effective way to weaken the basal texture of the rolled sheets, and Huang et al. have successfully increased the Index Erichsen (I.E.) value of a commerial AZ31 alloy sheet from 3.7 mm to 9.5 mm [7]. However, this method will be undesirable in large-scale manufacturing process due to limited capacity of production facilities or high production cost. So texture weakening needs to be achieved via industry viable rolling processing. Mg-3Al-0.4Mn (AM30, wt%) alloy is a new wrought Mg alloy with a good balance of strengths and ductility [11]. Also, the AM30 alloy has slightly better extrudability and deformability at elevated temperatures than AZ31 alloy [11,12], implying that AM30 alloy has good rolling ability. If AM30 alloy also exhibits moderate RT formability, the AM30 alloy may become a candidate as a sheet alloy. In this work, we firstly report a rolled AM30 alloy sheet ∗
Corresponding author. E-mail address: [email protected] (T. Nakata).
with good combination of RT formability and tensile properties via industry viable rolling processing consisting of homogenization, conventional warm rolling, and annealing. Mg–3Al–0.8Zn–0.4Mn and Mg–3Al–0.4Mn (wt%, AZ31 and AM30) alloy ingots were prepared by a direct chill casting. The ingots were homogenized at 415 °C for 2 h and 500 °C for 12 h in an electronic furnace under an Ar atmosphere, followed by water quenching. The homogenized ingots were machined to make slabs with 10 mm in thickness and 120 mm in width. They were soaked at 220 °C for 30 min in an electronic furnace and continuously rolled with a rolling reduction of 20%/pass. The rolling was repeated 10 times without reheating to obtain sheets with 1 mm in thickness. During the rolling, the temperature and speed of the rollers were fixed at 220 °C and 5 m/min, respectively. The time between each rolling pass was within 10 s. The rolled sheets were annealed at 250 °C for 4 h, followed by water quenching. Stretch formability of the annealed sheets were evaluated by an Erichsen cupping test at RT using square shape specimens with 60 × 60 mm2 . The punch diameter and speed were 20 mm and 6 mm/min, respectively. Tensile tests were also conducted using the annealed specimens with gauge length, width, and thickness of 20, 4, and 1 mm at an initial strain rate of 10−3 s–1 and RT along the rolling and transverse directions (RD and TD). The Erichsen cupping test was repeated five times and tensile tests were repeated three times. A scanning electron microscope (JEOL, JSM-70 0 0F) equipped with TSL electron backscattered diffraction (EBSD) apparatus was used to characterize the second phases,
https://doi.org/10.1016/j.scriptamat.2020.01.015 1359-6462/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22
17
Table 1 0.2% proof stress, ultimate tensile strength, and elongation to failure along rolling and transverse directions (RD and TD) of AZ31 and AM30 alloy sheets annealed for 250 °C for 4 h. Alloy
AZ31 AM30
Stretched along RD
Stretched along TD
0.2% proof stress [MPa]
Ultimate tensile strength [MPa]
Elongation to failure (%)
0.2% proof stress [MPa]
Ultimate tensile strength [MPa]
155±1 153±3
256±2 248±5
28±1 28±2
173±3 171±3
265±1 259±4
Elongation to failure (%) 25±1 25±2
Table 2 (0 0 01) [11-2¯ 0] Schmid factors (Schmid factors for basal slips) along rolling and transverse directions (RD and TD) of the AZ31 and AM30 alloy sheets annealed for 250 °C for 4 h. Alloy
AZ31 AM30
Schmid factor for basal slips Stretched along RD
Stretched along TD
0.27 0.28
0.22 0.26
grain sizes, and textures of the annealed samples. Microstructural development during the stretch forming in both AZ31 and AM30 samples was also characterized by the EBSD. The Erichsen cupping test was stopped after the punch stroke reaches to 2 mm and 4 mm. The deformed specimens were cut at the center along the RD, and the surface was polished by a Cross Section Polisher (JEOL, SM-09020CP). The EBSD scan was conducted on the top and bottom surfaces of the deformed specimens. The measured area of this analysis was 0.2 × 0.2 mm2 . To identify active dislocations, in-grain misorientation axes (IGMA) distribution was investigated using material point pairs having misorientation angles from 2° to 5° [13]. A quasi-in-situ EBSD analysis was also carried out to understand the microstructural development during the annealing. In the quasi-in-situ EBSD analysis, the EBSD scan was conducted firstly on the region of interest (ROI) of the as-rolled samples. The scanned samples were subjected to an annealing at 250 °C for 10, 30, 120 and 300 s in an infrared vacuum furnace (THERMO RIKO CO., LTD., IVF198W), and the EBSD scan was conducted again on the same ROI. Grains with grain orientation spread (GOS) less than 1° were regarded as recrystallized grains. The measured area for the quasi-in-situ EBSD analysis was about 0.2 mm2 . All microstructural characterization was done on the RD-ND plane, and the EBSD data was analyzed from the RD-TD plane. Fig. 1 shows tensile stress strain curves stretched along the RD and TD, appearances of the fractured specimens after the Erichsen cupping test, inverse pole figure maps, (0 0 01) pole figures, profile of the (0 0 01) pole density measured along the RD, and secondary electron images of the (a) AZ31 and (b) AM30 samples. Their average 0.2% proof stresses (PS), ultimate tensile strengths (UTS), and elongation to failure (EF) are summarized in Table 1 with the standard deviations, and Table 2 summarizes their (0 0 01) [11-2¯ 0] Schmid factors (Schmid factors for basal slips). Inset values in the snapshots of the fractured Erichsen cupping samples are average I.E. values with standard deviations. The PS and EF of both samples are almost the same. The AZ31 sample exhibits slightly higher UTS than those in the AM30 sample, which may be due to the solid-solution of Zn [14]. Though the AM30 sample stretched along the RD shows lower strengths and higher EF than those of the TD specimens, the difference is very small. The anisotropy is also small in the AZ31 sample; however, the AZ31 sample shows poor I.E. value of 4.7 mm. The AM30 sample shows significantly large I.E. value of 8.2 mm. This is similar to that of newly designed Mg–Al–Ca–Zn–Mn alloy sheets [15,16]. Both samples show fully recrystallized microstructure with average grain sizes of ~8 μm. The AZ31 sample forms RD-split texture as previously reported AZ31
Fig. 1. Tensile stress strain curves stretched along rolling and transverse directions (RD and TD), appearances of fractured specimens after Erichsen cupping test, inverse pole figure maps, (0 0 01) pole figures, profile of (0 0 01) pole density measured along the RD, and secondary electron images of (a) AZ31 and (b) AM30 alloy sheets. Inset I.E. values represent average Index Erichsen values with standard deviations. Dash lines in the profile of (0 0 01) pole density indicate the region of strong alignment of (0 0 01) poles to the ND.
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Fig. 2. Inverse pole figure maps, (0 0 01) pole figures of, and distributions in-grain misorientation axes (IGMA) of (a, c, e, f) AZ31 and (b, d, g, h) AM30 alloy sheets after interrupted Erichsen cupping test. Note that these images were taken from top surfaces of deformed specimens, and the inset values in the pole figures represent the maximum intensities of the poles.
alloy sheets [17,18], and the AM30 sample has similar preferential tilt of the basal poles to the RD. However, in the AM30 sample, most of the grains have circular distribution of basal poles away from the ND. Schmid factors for basal slips in the AM30 sample are higher than those in the AZ31 sample, and the AM30 sample has similar Schmid factors for basal slips along both RD and TD. These textural features indicate that, in the AM30 sample, large number of grains are favorably oriented for basal slips during RT stretch forming. Also, they are beneficial to reduce in-plane yield anisotropy. Both samples contain second phases. These phases may be Al8 Mn5 phases [11], and their sizes and fractions in both samples are almost the same. Fig. 2 shows inverse pole figure maps, (0 0 01) pole figures of the (a, b) AZ31 and (c, d) AM30 samples after the interrupted
Erichsen cupping test. Distributions of the IGMA in the (e, f) AZ31 and (g, h) AM30 samples are also shown in Fig. 2, and Table 3 summarizes fractions of low angle grain boundaries (LAGBs) having misorientation axis of [0 0 01], [101¯ 0], or [112¯ 0]. Note that misorientation angles of 2°−5° are considered. In both samples, the angular distribution of the basal poles gets narrower toward the ND, and the maximum intensities of the basal poles gradually increase with the progress of the deformation. Many IGMA are lying along [uv.0], and Table 3 shows that majority of LAGBs have misorientation axis of [10-1¯ 0] or [11-2¯ 0], suggesting that basal slips are the most dominant deformation mechanism during the RT stretch forming in both samples [13,19,20]. Some IGMA are lying along [0 0 01], and the fractions of LAGBs with misorientation axis of [0 0 01] are similar in both samples. This means that
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Fig. 3. Inverse pole figure maps, (0 0 01) pole figures, and misorientation angle distributions of (a, c, e) AZ31 and (b, d, f) AM30 alloy sheets after interrupted Erichsen cupping test taken from bottom surfaces of deformed specimens. The inset values in the pole figures represent the maximum intensities of the poles.
prismatic slips are also active and the activity is almost the same in both samples [13,20]. It is also interesting to note that intense shear bands are formed in the AZ31 sample as indicated by white dot lines, while such shear bands have not been observed in the AM30 sample. Fig. 3 shows inverse pole figure maps, (0 0 01) pole figures, and misorientation angle distributions of the (a, c, e) AZ31 and (b, d, f) AM30 samples after the interrupted Erichsen cupping test. Note that they were obtained from bottom surfaces of the deformed specimens. The inset values in the pole figures represent the maximum intensities of the poles. After the punch stroke reaches to 2 mm, distinct color changes occur within the grains,
and large fraction of the misorientation angles near 86° has been appeared, indicating that the formation of tensile twinning [21]. The AM30 sample has higher fraction of the misorientation angles near 86°, and the (0 0 01) pole figure shows clear change. These results suggest that tensile twinning is easily occurred in the AM30 sample. After the punch stroke reaches to 4 mm, the fraction of the misorientation angles near 86° decreases significantly, implying the occurrences of de-twinning [21]. In this work, we have found that AM30 alloy sheet exhibits excellent RT stretch formability and good tensile properties via industry viable rolling processing. The AM30 alloy sheet exhibits
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Fig. 4. Inverse pole figures and (0 0 01) pole figures of as-rolled and annealed (a) AZ31 and (b) AM30 alloy sheets. The inverse pole figure maps and the pole figures extracted from recrystallized grains (GOS < 1°) are also displayed in the right hand side, and the inset values in the pole figures represent maximum intensity of the (0 0 01) poles. Average sizes of recrystallized grains, drec , and percentages of the recrystallized grains with strong basal texture component, f0°−10° , (recrystallized grains having their basal poles tilting 0°−10° from the normal direction of the sheets) in the AZ31 and AM30 alloy sheets are also summarized in the (c) and (d). RD and TD in the figure represent rolling and transverse directions of the sheet.
large I.E. value of 8.2 mm, and the sheet shows PS and EF of 153 MPa and 28% along the RD. Also, the in-plane yield anisotropy is small, the sheet shows PS of 171 MPa along the TD with good EF of 25%. The balance of these properties is higher than the
recently developed RT formable Mg alloy sheets [7,22–24]. The PS of the AM30 alloy sheet is still lower than Mg–Al–Ca–Zn–Mn alloy sheets [15,16]; however, tensile properties can be improved by means of twin-roll casting [16]. Because the twin-roll casting is
T. Nakata, C. Xu and H. Ohashi et al. / Scripta Materialia 180 (2020) 16–22 Table 3 Fractions of low angle grain boundaries having misorientation axis of [0 0 01], [101¯ 0], or [11-2¯ 0] in AZ31 and AM30 alloy sheets after interrupted Erichsen cupping test with punch stroke of 2 mm and 4 mm. Note that misorientation angles of 2°−5° are considered for the calculation. Alloy
AZ31 AM30
Punch stroke [mm]
2 4 2 4
Misorientation axis [0 0 01]
[10-1¯ 0]
[11-2¯ 0]
3.6% 2.9% 2.8% 4.7%
37% 33% 38% 31%
35% 33% 37% 31%
becoming commercial processing for producing low-cost Mg alloy sheets [25], the AM30 alloy may become a promising sheet alloy in automotive industries. The moderate strengths and large ductility could be realized by the fine grain structure [26–28]. The fine grain structure along with unique texture feature, where the most of the grains have circular distribution of basal poles away from the ND, also contributes to small in-plane yield anisotropy due to similar activity of basal slips along the RD and TD (Table 2). The AZ31 alloy sheet also shows small in-plane yield anisotropy despite the difference of the Schmid factors along the RD and TD. This may be attributed to the higher solute strengthening of Zn against basal slips than prismatic slips [14]. Although both AZ31 and AM30 alloy sheets show similar tensile properties, the AM30 alloy sheet exhibits superior RT stretch formability with large I.E. value of 8.2 mm. In the AM30 alloy sheet, large number of grains are favorably oriented for basal slips during RT stretch forming due to unique texture feature, and this is important to accommodate thickness strain of the sheets [15]. As shown in Fig. 3, basal slips are dominant in both alloy sheets; therefore, it can be inferred that large I.E. value of the AM30 alloy sheet is mainly attributed to its unique texture feature. It is worth noting that tensile twinning is more actively occurred and the formation of shear bands is suppressed in the AM30 alloy sheet. It is reported that activation of twinning and de-twinning will contribute to the enhanced RT stretch formability to some extent [26,27], and formation of shear bands decreases RT stretch formability [20]. They suggest that enhanced activities of basal slips and twinning/de-twinning, and suppression of shear bands formation contribute to the excellent RT stretch formability in the AM30 alloy sheet. In-depth understanding of the effect of alloying element, microstructures, and texture on deformation behavior will be critical for further improvement of RT formability. To understand the formation of unique texture feature in the AM30 sample, quisi-in-situ EBSD was done. Fig. 4 shows inverse pole figures and (0 0 01) pole figures of the as-rolled and annealed (a) AZ31 and (b) AM30 samples. The inverse pole figure maps and the pole figures extracted from recrystallized grains are also displayed. Average recrystallized grain sizes, drec , and percentages of the recrystallized grains with strong basal texture component, f0°−10°, (recrystallized grains having their basal poles tilting 0°−10° from the ND) in both samples are also summarized in Fig. 4(c) and (d). In the as-rolled condition, both samples consist of deformed grains, shear bands, and small fraction of recrystallized grains. The areal fraction of the recrystallized grains and the drec in the AZ31 sample are 8% and 1.7 μm, while those in the AM30 sample are 3% and 0.9 μm, which means that the dynamic recrystallization and/or grain growth during the rolling is strongly retarded in the AM30 sample. The suppression of the dynamic recrystallization and/or grain growth may come from high melting point due to the absence of Zn [28]. The texture features of the as-rolled AZ31 and AM30 samples are almost the same; the f0°−10° in both samples are very similar. After the short-time annealing, recrystal-
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lized grains with RD-split texture feature mainly form at the shear bands. Further annealing mainly causes grain growth and decreases the fraction of the recrystallized grains with strong basal texture component. During the annealing, the AZ31 sample keeps the RD-split texture feature of the statically recrystallized grains and strong basal texture component of the dynamically recrystallized grains. The AM30 sample also shows similar texture development; however, the recrystallized grains start to show circular distribution of basal poles, and the f0°−10° decreases significantly by the progress of the static recrystallization. Previous work has shown that basal poles of dynamically recrystallized grains gradually align to ND during a hot rolling. During the subsequent heat-treatment, the dynamically recrystallized grains grow and thus strong basal texture develops [29]. This result implies that suppression of dynamic recrystallization during rolling process is one of the important factors to produce Mg sheets with weakened basal texture. As shown in Fig. 4(a) and (b), the fraction of the recrystallized grains in the AM30 sample is smaller than the AZ31 sample. This should be one reason why the AM30 sample forms different texture feature compared to the AZ31 sample. Detailed observation of active deformation modes during the rolling needs to be done to fully understand the unique texture feature in the AM30 alloy sheet. In summary, we have successfully developed a Mg–3Al–0.4Mn (AM30) alloy sheet with excellent RT formability, moderate strengths, and large ductility along with isotropic tensile properties. The newly developed AM30 alloy sheet exhibits good balance of RT formability and tensile properties due to its fine grain structure and unique texture feature, where the basal poles of most grains have circular distribution away from the ND. Such unique texture feature and good properties could be obtained using only commercial rolling processing such as homogenization, rolling at warm temperature, and annealing; therefore, the AM30 alloy will become a promising sheet alloy in automotive industries. Understanding the unique texture formation will be an important work to optimize processing conditions and alloy compositions for the RT formable Mg alloy sheets. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgment This work was supported by JSPS KAKENHI Grant Numbers JP19K15321, JP18H03837, and Advanced Low Carbon Technology Research and Development Program (ALCA), 12102886. Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2020.01. 015. References [1] E.F. Emley, Principle of Mangesium Technology, Pergamon Press, Oxford, 1966. [2] Y. Chino, H. Iwasaki, M. Mabuchi, H.T. Jeong, Mater. Sci. Eng.: A 466 (2007) 90–95. [3] B.-C. Suh, J.H. Kim, J.H. Bae, J.H. Hwang, M.-S. Shim, N.J. Kim, Acta Mater. 124 (2017) 268–279. [4] D.H. Kang, D.-W. Kim, S. Kim, G.T. Bae, K.H. Kim, N.J. Kim, Scr. Mater. 61 (2009) 768–771. [5] A. Styczynski, Ch. Hartig, J. Bohlen, D. Letzig, Scr. Mater. 50 (2004) 943–947. [6] Q. Jin, S.-Y. Shim, S.-G. Lim, Scr. Mater. 55 (2006) 843–846. [7] X. Huang, K. Suzuki, Y. Chino, Scr. Mater. 61 (2009) 445–448. [8] M. Eddahbi, J.A. del Valle, M.T. Pérez-Prado, O.A. Ruano, Mater. Sci. Eng.: A 410-411 (2005) 308–311.
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