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Improving tensile properties of a room-temperature formable and heat-treatable Mg–6Zn-0.2Ca (wt.%) alloy sheet via micro-alloying of Al and Mn
Materials Science & Engineering A xxx (xxxx) xxx
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Improving tensile properties of a room-temperature formable and heat-treatable Mg–6Zn-0.2Ca (wt.%) alloy sheet via micro-alloying of Al and Mn T. Nakata a, *, C. Xu b, S. Kamado a a b
Nagaoka University of Technology, 1603-1, Kamitomioka, Nagaoka, 940-2188, Japan School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Magnesium Rolling Room-temperature formability Aging Tensile property
We have developed a strong Mg–6Zn-0.2Ca (wt.%) based alloy sheet with good room-temperature formability via micro-alloying of Al and Mn. The Al and Mn micro-alloyed Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%, ZXAM6000) alloy sheet forms fine grain structure with an average grain size of 15 μm and weakened basal texture, in which (0001) planes of most grains split to both rolling and transverse directions, leading to a good Index Erichsen value of 7.2 mm along with moderate 0.2% proof stresses of 155 MPa (rolling direction) and 120 MPa (transverse direction). After an aging treatment, the alloy sheet shows notable increase in the strengths; the peak-aged ZXAM6000 alloy sheet exhibits high 0.2% proof stresses of 264 MPa and 209 MPa along the rolling and transverse directions. The balance of the 0.2% proof stresses and the Index Erichsen value of the developed alloy sheet is comparable to that of Al–Mg–Si based 6xxx alloy sheets; hence, the developed alloy sheet may broaden the application of wrought Mg alloys in automotive industries.
1. Introduction Development of strong Mg alloy sheets with room-temperature (RT) formability is inevitable for the wider application of light-weight Mg alloys in automotive industries. Commercial Mg alloy sheets such as Mg–3Al–1Zn (wt.%, AZ31) alloy sheet shows similar strengths to Al–Mg–Si based 6xxx alloy sheets which are widely used as automotive body panels [1]. However, the AZ31 alloy sheet generally shows poor RT formability compared to the Al–Mg–Si based alloy sheets. For example, the Index Erichsen (I.E.) value of the AZ31 alloy sheet is only ~3 mm at RT while that of the Al–Mg–Si based alloy sheets are around 10 mm [2, 3], resulting that the production cost of the Mg alloy sheets become much higher than the Al–Mg–Si based alloy sheets. The poor RT form ability in Mg alloy sheets arises from strong basal texture, in which the (0001) planes of most grains align parallel to the sheet planes [1]. To weaken the basal texture and improve the RT formability in Mg alloy sheets, much effort has been done on alloy development. Mg–Zn–Ca based alloys are the promising materials as RT formable alloy sheets developed so far, because they show unusual texture feature. The (0001) planes of the most grains in the Mg–Zn–Ca based alloy sheets split to the sheet transverse direction (TD) [4–6]. As a result, the Mg–Zn–Ca based
alloy sheets exhibit large I.E. values of 7–9 mm that are much better than the AZ31 alloy sheet [4–6]. However, the strengths of the Mg–Zn–Ca based alloy sheets are not sufficient compared to age-hardenable Al–Mg–Si based alloy sheets [1]. Especially, the 0.2% proof stresses (PS) along the TD of the Mg–Zn–Ca alloy sheets are very poor due to easy activation of basal slips [7–9]. This makes difficult to use Mg alloy sheets as structural components in automotive industries. On one hand, it has been demonstrated that a Ca micro-alloyed Mg-6.2Zn-0.5Zr-0.2Ca (wt. %, ZKX600) alloy sheet has a large age-hardening response and exhibits a very high PS of 286 MPa along with moderate elongation to failure (EF) of 15% due to finely dispersed β1’ precipitates [10]. Although the excellent tensile properties can be realized in the ZKX600 alloy sheet, the RT formability of the alloy sheet is poor, the I.E. value is only 5.2 mm [10]. Also, the addition of Zr is inevitable to obtain fine grain structure, and the Zr addition has been achieved mainly by using expensive Mg–Zr master alloys, causing an increase of cost of the final products [11]. In this work, we have focused on synergetic additions of low-cost Al and Mn elements [12] as the alternative for the Zr element and inves tigated the effect of micro-alloying of Al and Mn on the microstructures, RT stretch formability, and tensile properties of a Mg–6Zn-0.2Ca (wt.%) based alloy sheet.
* Corresponding author. E-mail address: [email protected] (T. Nakata). https://doi.org/10.1016/j.msea.2019.138690 Received 6 October 2019; Received in revised form 7 November 2019; Accepted 15 November 2019 Available online 16 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: T. Nakata, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138690
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Table 1 Nomenclatures and chemical compositions of alloy ingots used in this work [wt. %]. Alloy
Zn
Ca
Al
Mn
ZX60 ZXAM6000
5.6 5.6
0.2 0.2
– 0.1
– 0.1
Fig. 1. Appearances of fractured (a) Mg–6Zn-0.2Ca and (b) Mg–6Zn-0.2Ca0.1Al-0.1Mn (wt.%) alloy sheets after Erichsen cupping test. Inset values are the average Index Erichsen (I.E.) values with standard deviations.
2. Experimental procedure Alloy ingots with nominal compositions of Mg-6.0Zn-0.2Ca and Mg6.0Zn-0.2Ca-0.1Al-0.1Mn (wt%, ZX60 and ZXAM6000) alloys were prepared by KOJUNDO CHEMICAL LABORATORY CO., LTD.. The no menclatures and chemical compositions of the ingots are given in Table 1. The ingots were homogenized at 320 � C for 4 h and 350 � C for 4 h in an electronic furnace under an Ar atmosphere. The homogenized samples were sectioned into plates with 70 mm in length, 120 mm in width, and 10 mm in thickness for the subsequent rolling process. Firstly, the homogenized plates with 10 mm in thickness were hot-rolled to obtain the sheets with ~2 mm in thickness by 7 passes of rolling. The roller temperature, thickness reduction per pass, and roller speed were 300 � C, 20%/pass, and 5 m/min, respectively. Prior to the rolling, the homogenized plates were pre-heated at 300 � C for 10min, and the hot rolling was done without reheating the samples. The hot-rolled sheets were further rolled to a thickness ~1 mm by 4 passes of rolling. The roller temperature, thickness reduction per pass, and roller speed were 100 � C, 15%/pass, and 5 m/min, respectively. The sheets were reheated at 350 � C for 5 min and air-cooled to 200 � C prior to the each rolling pass. The rolled sheets were solution-treated at 350 � C for 30 min fol lowed by water quenching, and age hardening responses of the solutiontreated samples were measured by a Vickers Hardness tester (Mitsutoyo, HM-221). The aging was done at 170 � C in an oil bath. RT formability of the solution-treated samples were evaluated by an Erichsen cupping test. The Erichsen cupping tests were carried out at RT on rectangular sam ples with 60 � 60 mm2 using a sheet metal testing machine (ERICHSEN, Model 100) with a 20 mm hemispherical punch. In accordance with the JIS Z 2247, the blank force and punch speed were 10 kN and ~6 mm/ min, respectively. Tensile tests were conducted at RT on the solutiontreated and peak-aged samples using an Autograph AG-50kNI (Shi madzu) at an initial strain rate of 10 3 s 1. The loading directions were parallel to the rolling and transverse directions (RD and TD) of the sheets. For the tensile tests, sub-size specimens having gauge length and width of 20 mm and 4 mm were used based on the JIS Z 2241. The Erichsen cupping tests were repeated 4 times for each alloy, and the tensile tests were repeated for 3 times for each condition for reproduc ibility. Microstructural characterization was done using a scanning electron microscope (SEM, JEOL IT-500) equipped with an energy dispersive X-ray spectroscopy. To evaluate the grain size and texture of the solution-treated samples, an electron backscattered diffraction (EBSD) was carried out on a SEM (JEOL JSM-7000F) with TSL EBSD apparatus. Transmission electron microscope (TEM, JEOL JEM-2100F) equipped with high-angle annular scanning transmission electron mi croscopy (HAADF-STEM) and energy dispersive X-ray spectroscopy (EDS) detectors was also used to evaluate the sizes and number densities of precipitates in the peak-aged samples. Thicknesses measurement of the thin foils in the TEM observation was carried out using the conver gent beam electron diffraction (CBED) method [13]. Samples for TEM observation were prepared by a twin jet electro-polishing (Struers, TenuPol-5) using a solution of 300 ml 2-butoxy ethanol, 15.9 g lithium chloride, 33.5 g of magnesium perchlorate in 1500 ml methanol at a temperature of 40–50 � C. After the electro-polishing, the samples were ion-polished by a Gatan Precision Ion Polishing System (PIPS).
Fig. 2. Variations in the Vickers hardness of solution-treated Mg–6Zn-0.2Ca and Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets as a function of aging time. S.T. in Fig. 2 stands for solution-treated condition.
Fig. 3. Nominal tensile stress-strain curves of solution-treated and peak-aged Mg–6Zn-0.2Ca and Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets stretched along the rolling and transverse directions.
ZXAM6000 samples after the Erichsen cupping test. Inset values are the average I.E. values with standard deviations. The ZX60 sample shows better I.E. value of 6.8 mm than previously reported high strength ZKX600 alloy sheet [10]. The good room-temperature formability is slightly increased by the micro-alloying of Al and Mn, the ZXAM6000 sample exhibits good I.E. value of 7.2 mm. Fig. 2 shows variations in the Vickers hardness of the solution-treated ZX60 and ZXAM6000 alloy samples as a function of the aging time. S.T. in Fig. 2 stands for the solution-treated condition. The hardness of the solution-treated ZX60 sample is 52 HV. After the additions of Al and Mn, the hardness increases to 60 HV. The aging at 170 � C gradually increases the hardness of both samples, and they reach their peak hardness after the aging for 64 h. The hardness increment of ZX60 sample is 18 HV, and the sample reaches its peak hardness of 70 HV. The ZXAM6000 sample shows almost the same hardness increment, and the ZXAM6000 sample reaches its peak hardness of 78 HV by the aging for 64 h. Fig. 3 shows nominal tensile stress-strain curves of the solutiontreated and peak-aged (a) ZX60 and (b) ZXAM6000 samples stretched along the RD and TD. Table 2 summarizes their 0.2% proof stress (PS),
3. Results Fig. 1 shows appearances of the fractured (a) ZX60 and (b) 2
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figures of the solution-treated (a, c) ZX60 and (b, d) ZXAM6000 sam ples. The (0001) pole figures were obtained by rotating the EBSD data so that we could observe from the RD-TD plane of the samples. The average grain sizes and the Schmid factors for (0001) [1120] slips (basal slips) of the ZX60 and ZXAM6000 samples are also summarized in Table 3. The ZX60 sample has coarse grain structure with its average grain size over 30 μm. The additions of Al and Mn lead to substantial grain refinement, and the ZXAM6000 sample has much finer average grain size of 15 μm. Interestingly, the ZX60 sample has unique texture feature, where their (0001) planes of most grains split to both RD and TD with low maximum intensity of 0001 poles of 3.4 MRD. The quadruple texture feature also appears in the ZXAM6000 sample, and the sample also shows low maximum intensity of 0001 poles of 3.2 MRD. In both samples, the Schmid factors for basal slips along the TD are higher than that along the RD. In addition, the Schmid factors for basal slips in both samples are same, suggesting that both samples have similar activity of basal slips, and basal slips occur easily if the loading direction is parallel to the TD. Fig. 5 shows backscattered electron (BSE) images of the solutiontreated (a, c) ZX60 and (b, d) ZXAM6000 samples. Upper images were taken at a low magnification, and those displayed on the bottom column were taken at higher magnification. Both samples contain sec ond phase particles, and their distributions are almost the same. The average diameter of the particles are ~7 μm, and their areal fractions are
Table 2 0.2% proof stress (PS), ultimate tensile strength (UTS), and elongation to failure (EF) of solution-treated (S.T.) and peak-aged (P.A.) ZX60 and ZXAM6000 alloy sheets stretched along rolling and transverse directions. Alloy
ZX60
ZXAM6000
State
Stretched along rolling direction (RD)
Stretched along transverse direction (TD)
PS [MPa]
UTS [MPa]
EF (%)
PS [MPa]
UTS [MPa]
EF (%)
S.T.
121 � 4
246 � 5
13.5 � 0.8
93 � 1
244 � 5
14.0 � 0.4
P.A.
239 � 7
282 � 8
5.5 � 2.5
171 � 3
264 � 8
8.8 � 2.8
S.T.
155 � 2
273 � 1
20.0 � 1.1
120 � 4
268 � 3
23.4 � 2.7
P.A.
264 � 4
298 � 2
9.3 � 0.9
209 � 6
286 � 3
11.9 � 2.6
ultimate tensile strength (UTS), and elongation to failure (EF). S.T. and P.A. in Fig. 3 and Table 2 represent solution-treated and peak-aged conditions, respectively. In the solution-treated condition, the Al and Mn additions are effective to improve the PS and UTS. The PS and UTS are increased about 30 MPa and 20 MPa, respectively. The additions of Al and Mn also result in the improvement of the ductility. The ZXAM6000 sample exhibits EF over 20% in both RD and TD, which are 1.5 times higher than the ZX60 sample. Upon the aging, the ZX60 sample shows substantial improvement in the PS, and the PS along the RD and TD increases 118 MPa and 78 MPa, respectively. The PS of the ZXAM6000 sample are also substantially enhanced by the aging, the improvements of the PS along the RD and TD in the ZXAM6000 are 109 MPa and 79 MPa, respectively. After the aging, the UTS also increases, while the EF deteriorates. However, the ZXAM6000 sample keeps moderate EF of ~10% even after the aging. Fig. 4 shows (a, b) inverse pole figure maps and (c, d) (0001) pole
Table 3 Average grain sizes and Schmid factors for (0001) [1120] slips (basal slips) of the ZX60 and ZXAM6000 alloy sheets. Alloy
ZX60 ZXAM6000
Average grain size [μm]
Schmid factors for basal slips Stretched along rolling direction
Stretched along transverse direction
32 15
0.25 0.25
0.31 0.31
Fig. 4. Inverse pole figure maps and (0001) pole figures of solution-treated (a, c) Mg–6Zn-0.2Ca and (b, d) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets. Note that the (0001) pole figures were obtained by rotating the EBSD data so that we could observe from the rolling plane of the sheets. 3
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Fig. 5. Backscattered electron (BSE) images of solution-treated (a, c) Mg–6Zn-0.2Ca and (b, d) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets. Upper images were taken at a low magnification, and those displayed on the bottom column were taken at higher magnification.
about 7% for both samples. The microstructure of the solution-treated ZXAM6000 sample was further analyzed by the EDS, a BSE image and EDS elemental maps for Zn, Ca, Al, and Mn elements are displayed in Fig. 6. Zn, Ca, and Al elements are mainly enriched in the second phase particles. They may be Ca2Mg6Zn3 ternary phases (trigonal, a ¼ 0.97 nm, c ¼ 1.0 nm [14]) containing Al element. As indicated by white arrow-heads, some Mn-enriched particles are found at the grain boundaries, suggesting that they may pin the grain boundaries. Fig. 7 shows bright-field TEM images and selected area diffraction patterns of the peak-aged (a, c, e) ZX60 and (b, d, f) ZXAM6000 sam ples. Fig. 7 (a, b), (c, d), and (e, f) were recorded from the [0001], [1010], and [1120] directions, respectively. From the bright-field TEM images, it can be found that majority of the strengthening phases are rod-like precipitates growing along the [0001] direction. The selected area diffraction patterns recorded from the [1010] and [1120] di rections show extra spots along the [1120] and [1010] directions. This indicates that the rod-like precipitates are β1’ phases having an orien tation relationship of {1120}Mg//(0001) β1’ and [0001]Mg//[110] β1’ [15]. The number densities, diameters, and lengths of the β1’ phases are summarized in Table 4. Both samples form almost the same number density of the β1’ phases, and the sizes of these phases are also similar. Except for the rod-shaped β1’ phases, the ZX60 sample forms precipitates lying on the (0001) planes as indicated by yellow arrow-heads. They may be β2’ phases with plate-shaped morphology lying on the (0001) planes [15]. In the ZXAM6000 sample, as indicated by white arrow-heads, spherical precipitates also exist, and these precipitates are in contact with the rod-shaped β1’ phases. Fig. 8 shows a HAADF-STEM image and EDS elemental maps for Zn, Ca, Al, and Mn elements of the peak-aged ZXAM6000 sample. Note that they were taken along the [1120] direction. As indicated by white arrow-heads, the spherical precipitates which are contact with rod-shaped precipitates contain Al
Fig. 6. A BSE image and EDS elemental maps for Zn, Ca, Al, and Mn elements of a Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet. 4
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Fig. 7. Bright-field TEM images and selected area diffraction patterns of the peak-aged (a, c, e) Mg–6Zn-0.2Ca and (b, d, f) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets recorded from the (a, b) [0001], (c, d) [1010], and (e, f) [1120] directions.
Table 4 Number densities, diameters, and lengths of β1’ phases in the peak-aged ZX60 and ZXAM6000 alloy sheets. Alloy
Number density [m 3]
Diameter [nm]
Length [nm]
ZX60 ZXAM6000
1.1 � 1020 1.3 � 1020
14 16
278 264
Fig. 8. A HAADF-STEM image and EDS elemental maps for Zn, Ca, Al, and Mn elements of the peak-aged Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet.
various Mg alloy sheets [4–6,9,10,18–36] and Al–Mg–Si based 6xxx Al alloy sheets [3,37]. Some Mg–Al and Mg–Zn–Ca/RE based alloy sheets show excellent I.E. values of ~9 mm. However, their PS are much lower than the aged (or paint-baked) Al–Mg–Si based alloy sheets, this means that these Mg alloy sheets have insufficient strengths for automotive applications. The Mg–6Zn-0.2Ca-0.1Al-0.2Mn (wt.%, ZXAM6000) alloy sheet developed in this work exhibits large I.E. value of 7.2 mm. In addition, after the aging, the PS along the RD is increased to 264 MPa. Such balance of PS and I.E. value is similar to those of recently devel oped Mg–Al–Ca based alloy sheets with good strengths and RT form ability and Al–Mg–Si based alloy sheets. Also, it is interesting to note that the developed ZXAM6000 alloy sheet shows the PS of 209 MPa even along the TD, this is much higher than that of recently developed RT formable Mg alloy sheets [5,7–9]. It is generally accepted that poor strengths along certain loading direction are major hindrance of in dustrial applications of wrought Mg alloys; therefore, the ZXAM6000 alloy may become a potential sheet alloy that can be used as structural components such as automotive body panels. Both ZX60 and ZXAM6000 alloy sheets show good I.E. values of 6.8 mm and 7.2 mm, respectively. This is because the ZX based alloy sheets form quadruple texture where the (0001) poles of the most grains incline to both RD and TD. This texture feature contributes to the activation of basal slips during the stretch forming at room temperature, resulting in much higher I.E. values than those of commercial Mg–Al alloy sheets with strong basal texture [2]. It is interesting to note that the quadruple texture feature has not been obtained in the previously reported Ca
and Mn elements, suggesting that they may be Al8Mn5 phases [16]. They seem to be preferential nucleation sites of the Zn-enriched rod-shaped β1’ phases as previously reported Mg–Zn based alloy sheets [16,17]. Fig. 9 shows secondary electron (SE) and BSE images of the fracture surfaces of the solution-treated and peak-aged (a, b, c, d) ZX60 and (e, f, g, h) ZXAM6000 samples after the tensile test along the RD. S.T. and P. A. in Fig. 9 represent the solution-treated and peak-aged samples. The solution-treated ZX60 sample shows large amount of cleavage regions. Ductile dimple fracture with second phase particles also exist in the solution-treated ZX60 sample; however, the area of such region is limited. This indicates that twinning deformation dominates the fracture in the solution-treated ZX60 sample [18]. After the peak-aging, the ZX60 sample also shows brittle fracture feature, most of regions are occupied by cleavage regions. On the other hand, the fracture surfaces of the solution-treated and peak-aged ZXAM6000 sample mainly consist of ductile dimple fracture feature, implying that the additions of Al and Mn suppress twinning deformation during the tensile test. 4. Discussion In this work, we have developed a strong Mg–Zn–Ca alloy sheet with good room-temperature stretch formability via micro-alloying of Al and Mn. Fig. 10 shows relationships of PS along the RD and I.E. value of 5
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Fig. 10. Relationships of 0.2% proof stress (PS) along the rolling direction and Index Erichsen (I.E.) value of various Mg alloy and Al–Mg–Si based 6xxx Al alloy sheets.
that the additions of Al and Mn do not deteriorate the age hardenability of the ZX60 alloy sheet unlike Zr addition [10]. The ZXAM6000 alloy sheet forms almost the same number density of β1’ phases after the aging at 170 � C for 64 h. This causes significant improvement of the PS in the ZXAM6000 alloy sheet, resulting in the high PS of 264 MPa and 209 MPa along the RD and TD. The planar yield anisotropy still exist in the peak-aged ZXAM6000 alloy sheet; however, the PS along the TD is much higher than that of other RT formable Mg alloy sheets [5,7–9]. As well as the high strengths, the ZXAM6000 alloy sheet keeps its moderate ductility of ~10% even after the aging. This is because the fine grain structure of the ZXAM6000 alloy sheet could suppress the twinning deformation (Fig. 9) [39]. It has been demonstrated that the refinement of precipitates improves both strengths and ductility in age-hardened Mg alloys [10,43,44], and Al–Mn precipitates in Mg–Zn based alloy sheets may act as preferential nucleation sites for β1’ phases [16,17] (Figs. 7 and 8). Therefore, further enhancement of the strengths and ductility in the Mg–6Zn-0.2Ca (wt.%) alloy sheet can be possible via optimization of Al and Mn contents.
Fig. 9. Secondary electron (SE) and Backscattered electron (BSE) images of fracture surfaces after a tensile test of solution-treated and peak-aged (a, b, c, d) Mg–6Zn-0.2Ca and (e, f, g, h) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets stretched along the rolling direction. S.T. and P.A. in Fig. 9 represent solution-treated and peak-aged conditions, respectively.
micro-alloyed Mg–6Zn (wt.%) based alloy sheets [10], further study will be required on the texture development of the Mg–Zn–Ca based alloy sheets to unveil the texture weakening mechanism obtained in this work. Grain refining also contributes to enhanced RT stretch formability in Mg alloy sheets [38,39]. The additions of Al and Mn refine the average grain size of the ZX60 alloy sheet from 32 μm to 15 μm without a change of texture feature and activity of basal slips (Table 3). Therefore, the slightly improved I.E. value in the ZXAM6000 alloy sheet is attributed to the finer grain structure due to the additions of Al and Mn. The fine grain size may be attributed to particle pinning of Al–Mn particles and/or precipitates [12,16,17,40], and it is reported that grain refining will also contribute to enhanced tensile properties [41,42]. As shown in Fig. 3, the PS of the ZX60 alloy sheet along the RD and TD are increased from 120 MPa to 155 MPa and 90 MPa to 120 MPa by the additions of Al and Mn. Also, the additions of Al and Mn improve ductility of the ZX60 alloy sheet, so the solution-treated ZXAM6000 alloy sheet exhibits large elongation to failure over 20% in both RD and TD, which are ~1.5 times higher than those of the ZX60 alloy sheet. As shown in Fig. 4 and Table 3, the texture feature and activity of basal slips in both ZX60 and ZXAM6000 alloy sheets are the same, the enhanced tensile properties in the ZXAM6000 alloy sheet are attributed to the fine grain structure. Although the both alloy sheets have planar yield anisotropy: the PS along the TD is lower than that for the RD due to ease activation of basal slips along the TD (Table 3), it is worth mentioning
5. Summary In this work, we have successfully improved room-temperature stretch formability and tensile properties of a Mg–6Zn-0.2Ca (wt.%) alloy sheet via micro-alloying of low-cost Al and Mn. A Mg–6Zn-0.2Ca0.1Al-0.1Mn (wt.%) alloy sheet forms fine grain structure with an average grain size of 15 μm and a unique texture feature with splitting of (0001) planes towards the rolling and transverse directions. Such microstructural feature contributes to a good Index Erichsen value of 7.2 mm, large elongation to failure over 20%, and moderate 0.2% proof stress of 155 MPa along the rolling direction. After an aging treatment, high number density of rod-shaped β1’precipitates are distributed within 6
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the Mg matrix. This leads to a significant improvement of the 0.2% proof stress of the Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet; the peakaged alloy sheet exhibits the high 0.2% proof stress of 264 MPa along the rolling direction. Also, the peak-aged Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet shows the 0.2% proof stress of 209 MPa even along the transverse direction. These results indicate that Al and Mn microalloyed Mg–Zn–Ca based alloy sheets will become promising candi dates in automotive industries where the room-temperature formability, strengths, and ductility are important properties.
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Nakata, C. Xu, Y. Uehara, T.T. Sasaki, S. Kamado, Origin of texture weakening in a rolled ZEX4101 alloy sheet and its effect on room temperature formability and tensile property, J. Alloy. Comp. 782 (2019) 304–314. [31] Y. Chino, M. Kado, M. Mabuchi, Enhancement of tensile ductility and stretch formability of magnesium by addition of 0.2 wt%(0.035 at%)Ce, Mater. Sci. Eng. A 494 (2008) 343–349. [32] D. Wu, R.S. Chen, E.H. Han, Excellent room-temperature ductility and formability of rolled Mg–Gd–Zn alloy sheets, J. Alloy. Comp. 509 (2011) 2856–2863. [33] M.Z. Bian, T.T. Sasaki, B.C. Suh, T. Nakata, S. Kamado, K. Hono, A heat-treatable Mg–Al–Ca–Mn–Zn sheet alloy with good room temperature formability, Scr. Mater. 138 (2017) 151–155. [34] M.Z. Bian, T.T. Sasaki, T. Nakata, S. Kamado, K. Hono, Effects of rolling conditions on the microstructure and mechanical properties in a Mg–Al–Ca–Mn–Zn alloy sheet, Mater. Sci. Eng. A 730 (2018) 147–154. [35] M.Z. Bian, T.T. Sasaki, T. Nakata, Y. Yoshida, N. Kawabe, S. Kamado, K. Hono, Bake-hardenable Mg–Al–Zn–Mn–Ca sheet alloy processed by twin-roll casting, Acta Mater. 158 (2018) 278–288. [36] T.T.T. Trang, J.H. Zhang, J.H. Kim, A. Zargaran, J.H. Hwang, B.C. Suh, N.J. Kim, Designing a magnesium alloy with high strength and high formability, Nat. Commun. 9 (2018), 2522. [37] F. Ozturk, E. Esener, S. Toros, C.R. Picu, Effects of aging parameters on formability of 6061-O alloy, Mater. Des. 31 (2010) 4847–4852. [38] H. Somekawa, A. Kinoshita, K. Washio, A. Kato, Enhancement of room temperature stretch formability via grain boundary sliding in magnesium alloy, Mater. Sci. Eng. A 676 (2016) 427–433. [39] X. Huang, Y. Chino, M. Mabuchi, M. Matsuda, Influences of grain size on mechanical properties and cold formability of Mg–3Al–1Zn alloy sheets with similar weak initial textures, Mater. Sci. Eng. A 611 (2014) 152–161. [40] S.A. Khan, Y. Miyashita, Y. Mutoh, Z.B. Sajuri, Influence of Mn content on mechanical properties and fatigue behavior of extruded Mg alloys, Mater. Sci. Eng. A 420 (2006) 315–321. [41] J. Koike, T. Kobayashi, T. Mukai, H. Watanabe, M. Suzuki, K. Maruyama, K. Higashi, The activity of non-basal slip systems and dynamic recovery at room temperature in fine-grained AZ31B magnesium alloys, Acta Mater. 51 (2003) 2055–2065. [42] R.W. Armstrong, The influence of polycrystal grain size on several mechanical properties of materials, Metall. Mater. Trans. A 1 (1970) 1169–1176. [43] J.F. Nie, Effects of precipitate shape and orientation on dispersion strengthening in magnesium alloys, Scr. Mater. 48 (2003) 1009–1015. [44] T.T. Sasaki, F.R. Elsayed, T. Nakata, T. Ohkubo, S. Kamado, K. Hono, Strong and ductile heat-treatable Mg–Sn–Zn–Al wrought alloys, Acta Mater. 99 (2015) 176–186.
Data availability All data are available from the corresponding author on reasonable request. 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. Acknowledgement This work was supported by JSPS KAKENHI Grant Numbers JP19K15321, JP18H03837, and Advanced Low Carbon Technology Research and Development Program(ALCA), 12102886. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2019.138690. References [1] J. Hirsh, T. Al-Samman, Superior light metals by texture engineering: optimized aluminum and magnesium alloys for automotive applications, Acta Mater. 61 (2013) 818–843. [2] K. Iwanaga, H. Tashiro, H. Okamoto, K. Shimizu, Improvement of formability from room temperature to warm temperature in AZ-31 magnesium alloy, J. Mater. Process. Technol. 155–156 (2004) 1313–1316. [3] S.M. Hirth, G.J. Marshall, S.A. Court, D.J. Lloyd, Effects of Si on the aging behaviour and formability of aluminium alloys based on AA6016, Mater. Sci. Eng. A 319–321 (2001) 452–456. [4] D.-W. Kim, B.-C. Suh, M.-S. Shim, J.H. Bae, D.H. Kim, N.J. Kim, Texture evolution in Mg-Zn-Ca alloy sheets, Metall. Mater. Trans. A 44 (2013) 2950–2961. [5] Y. Chino, T. Ueda, Y. Otomatsu, K. Sassa, X. Huang, K. Suzuki, M. Mabuchi, Effects of Ca on tensile properties and stretch formability at room temperature in Mg-Zn and Mg-Al alloys, Mater. Trans. 52 (2011) 1477–1482. [6] S.J. Park, H.C. Jung, K.S. Shin, Deformation behaviors of twin roll cast Mg-Zn-X-Ca alloys for enhanced room-temperature formability, Mater. Sci. Eng. A 679 (2017) 329–339. [7] J. Bohlen, M.R. Nürnberg, J.W. Senn, D. Letzig, S.R. Agnew, The texture and anisotropy of magnesium–zinc–rare earth alloy sheets, Acta Mater. 55 (2007) 2101–2112. [8] S. Yi, J. Bohlen, F. Heinemann, D. Letzig, Mechanical anisotropy and deep drawing behaviour of AZ31 and ZE10 magnesium alloy sheets, Acta Mater. 58 (2010) 592–605. [9] B.-C. Suh, J.H. Kim, J.H. Hwang, M.-S. Shim, N.J. Kim, Twinning-mediated formability in Mg alloys, Sci. Rep. 6 (2016) 22364. [10] T. Bhattacharjee, B.-C. Suh, T.T. Sasaki, T. Ohkubo, N.J. Kim, K. Hono, High strength and formable Mg–6.2Zn–0.5Zr–0.2Ca alloy sheet processed by twin roll casting, Mater. Sci. Eng. A 609 (2014) 154–160. [11] M. Qian, A. Das, Grain refinement of magnesium alloys by zirconium: formation of equiaxed grains, Scr. Mater. 54 (2006) 881–886. [12] T. Nakata, C. Xu, R. Ajima, Y. Matsumoto, K. Shimizu, T.T. Sasaki, K. Hono, S. Kamado, Improving mechanical properties and yield asymmetry in high-speed extrudable Mg-1.1Al-0.24Ca (wt%) alloy by high Mn addition, Mater. Sci. Eng. A 712 (2018) 12–19. [13] P.M. Kelly, A. Jostons, R.G. Blake, J.G. Napier, The determination of foil thickness by scanning transmission electron microscopy, Phys. Status Solidi 31 (1975) 771–780. [14] P.M. Jardim, G. Sol� orzano, J.B.V. Sande, Precipitate crystal structure determination in melt spun Mg-1.5wt%Ca-6wt%Zn alloy, Microsc. Microanal. 8 (2002) 487–496.
7
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Improving tensile properties of a room-temperature formable and heat-treatable Mg–6Zn-0.2Ca (wt.%) alloy sheet via micro-alloying of Al and Mn T. Nakata a, *, C. Xu b, S. Kamado a a b
Nagaoka University of Technology, 1603-1, Kamitomioka, Nagaoka, 940-2188, Japan School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Magnesium Rolling Room-temperature formability Aging Tensile property
We have developed a strong Mg–6Zn-0.2Ca (wt.%) based alloy sheet with good room-temperature formability via micro-alloying of Al and Mn. The Al and Mn micro-alloyed Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%, ZXAM6000) alloy sheet forms fine grain structure with an average grain size of 15 μm and weakened basal texture, in which (0001) planes of most grains split to both rolling and transverse directions, leading to a good Index Erichsen value of 7.2 mm along with moderate 0.2% proof stresses of 155 MPa (rolling direction) and 120 MPa (transverse direction). After an aging treatment, the alloy sheet shows notable increase in the strengths; the peak-aged ZXAM6000 alloy sheet exhibits high 0.2% proof stresses of 264 MPa and 209 MPa along the rolling and transverse directions. The balance of the 0.2% proof stresses and the Index Erichsen value of the developed alloy sheet is comparable to that of Al–Mg–Si based 6xxx alloy sheets; hence, the developed alloy sheet may broaden the application of wrought Mg alloys in automotive industries.
1. Introduction Development of strong Mg alloy sheets with room-temperature (RT) formability is inevitable for the wider application of light-weight Mg alloys in automotive industries. Commercial Mg alloy sheets such as Mg–3Al–1Zn (wt.%, AZ31) alloy sheet shows similar strengths to Al–Mg–Si based 6xxx alloy sheets which are widely used as automotive body panels [1]. However, the AZ31 alloy sheet generally shows poor RT formability compared to the Al–Mg–Si based alloy sheets. For example, the Index Erichsen (I.E.) value of the AZ31 alloy sheet is only ~3 mm at RT while that of the Al–Mg–Si based alloy sheets are around 10 mm [2, 3], resulting that the production cost of the Mg alloy sheets become much higher than the Al–Mg–Si based alloy sheets. The poor RT form ability in Mg alloy sheets arises from strong basal texture, in which the (0001) planes of most grains align parallel to the sheet planes [1]. To weaken the basal texture and improve the RT formability in Mg alloy sheets, much effort has been done on alloy development. Mg–Zn–Ca based alloys are the promising materials as RT formable alloy sheets developed so far, because they show unusual texture feature. The (0001) planes of the most grains in the Mg–Zn–Ca based alloy sheets split to the sheet transverse direction (TD) [4–6]. As a result, the Mg–Zn–Ca based
alloy sheets exhibit large I.E. values of 7–9 mm that are much better than the AZ31 alloy sheet [4–6]. However, the strengths of the Mg–Zn–Ca based alloy sheets are not sufficient compared to age-hardenable Al–Mg–Si based alloy sheets [1]. Especially, the 0.2% proof stresses (PS) along the TD of the Mg–Zn–Ca alloy sheets are very poor due to easy activation of basal slips [7–9]. This makes difficult to use Mg alloy sheets as structural components in automotive industries. On one hand, it has been demonstrated that a Ca micro-alloyed Mg-6.2Zn-0.5Zr-0.2Ca (wt. %, ZKX600) alloy sheet has a large age-hardening response and exhibits a very high PS of 286 MPa along with moderate elongation to failure (EF) of 15% due to finely dispersed β1’ precipitates [10]. Although the excellent tensile properties can be realized in the ZKX600 alloy sheet, the RT formability of the alloy sheet is poor, the I.E. value is only 5.2 mm [10]. Also, the addition of Zr is inevitable to obtain fine grain structure, and the Zr addition has been achieved mainly by using expensive Mg–Zr master alloys, causing an increase of cost of the final products [11]. In this work, we have focused on synergetic additions of low-cost Al and Mn elements [12] as the alternative for the Zr element and inves tigated the effect of micro-alloying of Al and Mn on the microstructures, RT stretch formability, and tensile properties of a Mg–6Zn-0.2Ca (wt.%) based alloy sheet.
* Corresponding author. E-mail address: [email protected] (T. Nakata). https://doi.org/10.1016/j.msea.2019.138690 Received 6 October 2019; Received in revised form 7 November 2019; Accepted 15 November 2019 Available online 16 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: T. Nakata, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138690
T. Nakata et al.
Materials Science & Engineering A xxx (xxxx) xxx
Table 1 Nomenclatures and chemical compositions of alloy ingots used in this work [wt. %]. Alloy
Zn
Ca
Al
Mn
ZX60 ZXAM6000
5.6 5.6
0.2 0.2
– 0.1
– 0.1
Fig. 1. Appearances of fractured (a) Mg–6Zn-0.2Ca and (b) Mg–6Zn-0.2Ca0.1Al-0.1Mn (wt.%) alloy sheets after Erichsen cupping test. Inset values are the average Index Erichsen (I.E.) values with standard deviations.
2. Experimental procedure Alloy ingots with nominal compositions of Mg-6.0Zn-0.2Ca and Mg6.0Zn-0.2Ca-0.1Al-0.1Mn (wt%, ZX60 and ZXAM6000) alloys were prepared by KOJUNDO CHEMICAL LABORATORY CO., LTD.. The no menclatures and chemical compositions of the ingots are given in Table 1. The ingots were homogenized at 320 � C for 4 h and 350 � C for 4 h in an electronic furnace under an Ar atmosphere. The homogenized samples were sectioned into plates with 70 mm in length, 120 mm in width, and 10 mm in thickness for the subsequent rolling process. Firstly, the homogenized plates with 10 mm in thickness were hot-rolled to obtain the sheets with ~2 mm in thickness by 7 passes of rolling. The roller temperature, thickness reduction per pass, and roller speed were 300 � C, 20%/pass, and 5 m/min, respectively. Prior to the rolling, the homogenized plates were pre-heated at 300 � C for 10min, and the hot rolling was done without reheating the samples. The hot-rolled sheets were further rolled to a thickness ~1 mm by 4 passes of rolling. The roller temperature, thickness reduction per pass, and roller speed were 100 � C, 15%/pass, and 5 m/min, respectively. The sheets were reheated at 350 � C for 5 min and air-cooled to 200 � C prior to the each rolling pass. The rolled sheets were solution-treated at 350 � C for 30 min fol lowed by water quenching, and age hardening responses of the solutiontreated samples were measured by a Vickers Hardness tester (Mitsutoyo, HM-221). The aging was done at 170 � C in an oil bath. RT formability of the solution-treated samples were evaluated by an Erichsen cupping test. The Erichsen cupping tests were carried out at RT on rectangular sam ples with 60 � 60 mm2 using a sheet metal testing machine (ERICHSEN, Model 100) with a 20 mm hemispherical punch. In accordance with the JIS Z 2247, the blank force and punch speed were 10 kN and ~6 mm/ min, respectively. Tensile tests were conducted at RT on the solutiontreated and peak-aged samples using an Autograph AG-50kNI (Shi madzu) at an initial strain rate of 10 3 s 1. The loading directions were parallel to the rolling and transverse directions (RD and TD) of the sheets. For the tensile tests, sub-size specimens having gauge length and width of 20 mm and 4 mm were used based on the JIS Z 2241. The Erichsen cupping tests were repeated 4 times for each alloy, and the tensile tests were repeated for 3 times for each condition for reproduc ibility. Microstructural characterization was done using a scanning electron microscope (SEM, JEOL IT-500) equipped with an energy dispersive X-ray spectroscopy. To evaluate the grain size and texture of the solution-treated samples, an electron backscattered diffraction (EBSD) was carried out on a SEM (JEOL JSM-7000F) with TSL EBSD apparatus. Transmission electron microscope (TEM, JEOL JEM-2100F) equipped with high-angle annular scanning transmission electron mi croscopy (HAADF-STEM) and energy dispersive X-ray spectroscopy (EDS) detectors was also used to evaluate the sizes and number densities of precipitates in the peak-aged samples. Thicknesses measurement of the thin foils in the TEM observation was carried out using the conver gent beam electron diffraction (CBED) method [13]. Samples for TEM observation were prepared by a twin jet electro-polishing (Struers, TenuPol-5) using a solution of 300 ml 2-butoxy ethanol, 15.9 g lithium chloride, 33.5 g of magnesium perchlorate in 1500 ml methanol at a temperature of 40–50 � C. After the electro-polishing, the samples were ion-polished by a Gatan Precision Ion Polishing System (PIPS).
Fig. 2. Variations in the Vickers hardness of solution-treated Mg–6Zn-0.2Ca and Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets as a function of aging time. S.T. in Fig. 2 stands for solution-treated condition.
Fig. 3. Nominal tensile stress-strain curves of solution-treated and peak-aged Mg–6Zn-0.2Ca and Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets stretched along the rolling and transverse directions.
ZXAM6000 samples after the Erichsen cupping test. Inset values are the average I.E. values with standard deviations. The ZX60 sample shows better I.E. value of 6.8 mm than previously reported high strength ZKX600 alloy sheet [10]. The good room-temperature formability is slightly increased by the micro-alloying of Al and Mn, the ZXAM6000 sample exhibits good I.E. value of 7.2 mm. Fig. 2 shows variations in the Vickers hardness of the solution-treated ZX60 and ZXAM6000 alloy samples as a function of the aging time. S.T. in Fig. 2 stands for the solution-treated condition. The hardness of the solution-treated ZX60 sample is 52 HV. After the additions of Al and Mn, the hardness increases to 60 HV. The aging at 170 � C gradually increases the hardness of both samples, and they reach their peak hardness after the aging for 64 h. The hardness increment of ZX60 sample is 18 HV, and the sample reaches its peak hardness of 70 HV. The ZXAM6000 sample shows almost the same hardness increment, and the ZXAM6000 sample reaches its peak hardness of 78 HV by the aging for 64 h. Fig. 3 shows nominal tensile stress-strain curves of the solutiontreated and peak-aged (a) ZX60 and (b) ZXAM6000 samples stretched along the RD and TD. Table 2 summarizes their 0.2% proof stress (PS),
3. Results Fig. 1 shows appearances of the fractured (a) ZX60 and (b) 2
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figures of the solution-treated (a, c) ZX60 and (b, d) ZXAM6000 sam ples. The (0001) pole figures were obtained by rotating the EBSD data so that we could observe from the RD-TD plane of the samples. The average grain sizes and the Schmid factors for (0001) [1120] slips (basal slips) of the ZX60 and ZXAM6000 samples are also summarized in Table 3. The ZX60 sample has coarse grain structure with its average grain size over 30 μm. The additions of Al and Mn lead to substantial grain refinement, and the ZXAM6000 sample has much finer average grain size of 15 μm. Interestingly, the ZX60 sample has unique texture feature, where their (0001) planes of most grains split to both RD and TD with low maximum intensity of 0001 poles of 3.4 MRD. The quadruple texture feature also appears in the ZXAM6000 sample, and the sample also shows low maximum intensity of 0001 poles of 3.2 MRD. In both samples, the Schmid factors for basal slips along the TD are higher than that along the RD. In addition, the Schmid factors for basal slips in both samples are same, suggesting that both samples have similar activity of basal slips, and basal slips occur easily if the loading direction is parallel to the TD. Fig. 5 shows backscattered electron (BSE) images of the solutiontreated (a, c) ZX60 and (b, d) ZXAM6000 samples. Upper images were taken at a low magnification, and those displayed on the bottom column were taken at higher magnification. Both samples contain sec ond phase particles, and their distributions are almost the same. The average diameter of the particles are ~7 μm, and their areal fractions are
Table 2 0.2% proof stress (PS), ultimate tensile strength (UTS), and elongation to failure (EF) of solution-treated (S.T.) and peak-aged (P.A.) ZX60 and ZXAM6000 alloy sheets stretched along rolling and transverse directions. Alloy
ZX60
ZXAM6000
State
Stretched along rolling direction (RD)
Stretched along transverse direction (TD)
PS [MPa]
UTS [MPa]
EF (%)
PS [MPa]
UTS [MPa]
EF (%)
S.T.
121 � 4
246 � 5
13.5 � 0.8
93 � 1
244 � 5
14.0 � 0.4
P.A.
239 � 7
282 � 8
5.5 � 2.5
171 � 3
264 � 8
8.8 � 2.8
S.T.
155 � 2
273 � 1
20.0 � 1.1
120 � 4
268 � 3
23.4 � 2.7
P.A.
264 � 4
298 � 2
9.3 � 0.9
209 � 6
286 � 3
11.9 � 2.6
ultimate tensile strength (UTS), and elongation to failure (EF). S.T. and P.A. in Fig. 3 and Table 2 represent solution-treated and peak-aged conditions, respectively. In the solution-treated condition, the Al and Mn additions are effective to improve the PS and UTS. The PS and UTS are increased about 30 MPa and 20 MPa, respectively. The additions of Al and Mn also result in the improvement of the ductility. The ZXAM6000 sample exhibits EF over 20% in both RD and TD, which are 1.5 times higher than the ZX60 sample. Upon the aging, the ZX60 sample shows substantial improvement in the PS, and the PS along the RD and TD increases 118 MPa and 78 MPa, respectively. The PS of the ZXAM6000 sample are also substantially enhanced by the aging, the improvements of the PS along the RD and TD in the ZXAM6000 are 109 MPa and 79 MPa, respectively. After the aging, the UTS also increases, while the EF deteriorates. However, the ZXAM6000 sample keeps moderate EF of ~10% even after the aging. Fig. 4 shows (a, b) inverse pole figure maps and (c, d) (0001) pole
Table 3 Average grain sizes and Schmid factors for (0001) [1120] slips (basal slips) of the ZX60 and ZXAM6000 alloy sheets. Alloy
ZX60 ZXAM6000
Average grain size [μm]
Schmid factors for basal slips Stretched along rolling direction
Stretched along transverse direction
32 15
0.25 0.25
0.31 0.31
Fig. 4. Inverse pole figure maps and (0001) pole figures of solution-treated (a, c) Mg–6Zn-0.2Ca and (b, d) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets. Note that the (0001) pole figures were obtained by rotating the EBSD data so that we could observe from the rolling plane of the sheets. 3
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Fig. 5. Backscattered electron (BSE) images of solution-treated (a, c) Mg–6Zn-0.2Ca and (b, d) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets. Upper images were taken at a low magnification, and those displayed on the bottom column were taken at higher magnification.
about 7% for both samples. The microstructure of the solution-treated ZXAM6000 sample was further analyzed by the EDS, a BSE image and EDS elemental maps for Zn, Ca, Al, and Mn elements are displayed in Fig. 6. Zn, Ca, and Al elements are mainly enriched in the second phase particles. They may be Ca2Mg6Zn3 ternary phases (trigonal, a ¼ 0.97 nm, c ¼ 1.0 nm [14]) containing Al element. As indicated by white arrow-heads, some Mn-enriched particles are found at the grain boundaries, suggesting that they may pin the grain boundaries. Fig. 7 shows bright-field TEM images and selected area diffraction patterns of the peak-aged (a, c, e) ZX60 and (b, d, f) ZXAM6000 sam ples. Fig. 7 (a, b), (c, d), and (e, f) were recorded from the [0001], [1010], and [1120] directions, respectively. From the bright-field TEM images, it can be found that majority of the strengthening phases are rod-like precipitates growing along the [0001] direction. The selected area diffraction patterns recorded from the [1010] and [1120] di rections show extra spots along the [1120] and [1010] directions. This indicates that the rod-like precipitates are β1’ phases having an orien tation relationship of {1120}Mg//(0001) β1’ and [0001]Mg//[110] β1’ [15]. The number densities, diameters, and lengths of the β1’ phases are summarized in Table 4. Both samples form almost the same number density of the β1’ phases, and the sizes of these phases are also similar. Except for the rod-shaped β1’ phases, the ZX60 sample forms precipitates lying on the (0001) planes as indicated by yellow arrow-heads. They may be β2’ phases with plate-shaped morphology lying on the (0001) planes [15]. In the ZXAM6000 sample, as indicated by white arrow-heads, spherical precipitates also exist, and these precipitates are in contact with the rod-shaped β1’ phases. Fig. 8 shows a HAADF-STEM image and EDS elemental maps for Zn, Ca, Al, and Mn elements of the peak-aged ZXAM6000 sample. Note that they were taken along the [1120] direction. As indicated by white arrow-heads, the spherical precipitates which are contact with rod-shaped precipitates contain Al
Fig. 6. A BSE image and EDS elemental maps for Zn, Ca, Al, and Mn elements of a Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet. 4
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Fig. 7. Bright-field TEM images and selected area diffraction patterns of the peak-aged (a, c, e) Mg–6Zn-0.2Ca and (b, d, f) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets recorded from the (a, b) [0001], (c, d) [1010], and (e, f) [1120] directions.
Table 4 Number densities, diameters, and lengths of β1’ phases in the peak-aged ZX60 and ZXAM6000 alloy sheets. Alloy
Number density [m 3]
Diameter [nm]
Length [nm]
ZX60 ZXAM6000
1.1 � 1020 1.3 � 1020
14 16
278 264
Fig. 8. A HAADF-STEM image and EDS elemental maps for Zn, Ca, Al, and Mn elements of the peak-aged Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet.
various Mg alloy sheets [4–6,9,10,18–36] and Al–Mg–Si based 6xxx Al alloy sheets [3,37]. Some Mg–Al and Mg–Zn–Ca/RE based alloy sheets show excellent I.E. values of ~9 mm. However, their PS are much lower than the aged (or paint-baked) Al–Mg–Si based alloy sheets, this means that these Mg alloy sheets have insufficient strengths for automotive applications. The Mg–6Zn-0.2Ca-0.1Al-0.2Mn (wt.%, ZXAM6000) alloy sheet developed in this work exhibits large I.E. value of 7.2 mm. In addition, after the aging, the PS along the RD is increased to 264 MPa. Such balance of PS and I.E. value is similar to those of recently devel oped Mg–Al–Ca based alloy sheets with good strengths and RT form ability and Al–Mg–Si based alloy sheets. Also, it is interesting to note that the developed ZXAM6000 alloy sheet shows the PS of 209 MPa even along the TD, this is much higher than that of recently developed RT formable Mg alloy sheets [5,7–9]. It is generally accepted that poor strengths along certain loading direction are major hindrance of in dustrial applications of wrought Mg alloys; therefore, the ZXAM6000 alloy may become a potential sheet alloy that can be used as structural components such as automotive body panels. Both ZX60 and ZXAM6000 alloy sheets show good I.E. values of 6.8 mm and 7.2 mm, respectively. This is because the ZX based alloy sheets form quadruple texture where the (0001) poles of the most grains incline to both RD and TD. This texture feature contributes to the activation of basal slips during the stretch forming at room temperature, resulting in much higher I.E. values than those of commercial Mg–Al alloy sheets with strong basal texture [2]. It is interesting to note that the quadruple texture feature has not been obtained in the previously reported Ca
and Mn elements, suggesting that they may be Al8Mn5 phases [16]. They seem to be preferential nucleation sites of the Zn-enriched rod-shaped β1’ phases as previously reported Mg–Zn based alloy sheets [16,17]. Fig. 9 shows secondary electron (SE) and BSE images of the fracture surfaces of the solution-treated and peak-aged (a, b, c, d) ZX60 and (e, f, g, h) ZXAM6000 samples after the tensile test along the RD. S.T. and P. A. in Fig. 9 represent the solution-treated and peak-aged samples. The solution-treated ZX60 sample shows large amount of cleavage regions. Ductile dimple fracture with second phase particles also exist in the solution-treated ZX60 sample; however, the area of such region is limited. This indicates that twinning deformation dominates the fracture in the solution-treated ZX60 sample [18]. After the peak-aging, the ZX60 sample also shows brittle fracture feature, most of regions are occupied by cleavage regions. On the other hand, the fracture surfaces of the solution-treated and peak-aged ZXAM6000 sample mainly consist of ductile dimple fracture feature, implying that the additions of Al and Mn suppress twinning deformation during the tensile test. 4. Discussion In this work, we have developed a strong Mg–Zn–Ca alloy sheet with good room-temperature stretch formability via micro-alloying of Al and Mn. Fig. 10 shows relationships of PS along the RD and I.E. value of 5
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Fig. 10. Relationships of 0.2% proof stress (PS) along the rolling direction and Index Erichsen (I.E.) value of various Mg alloy and Al–Mg–Si based 6xxx Al alloy sheets.
that the additions of Al and Mn do not deteriorate the age hardenability of the ZX60 alloy sheet unlike Zr addition [10]. The ZXAM6000 alloy sheet forms almost the same number density of β1’ phases after the aging at 170 � C for 64 h. This causes significant improvement of the PS in the ZXAM6000 alloy sheet, resulting in the high PS of 264 MPa and 209 MPa along the RD and TD. The planar yield anisotropy still exist in the peak-aged ZXAM6000 alloy sheet; however, the PS along the TD is much higher than that of other RT formable Mg alloy sheets [5,7–9]. As well as the high strengths, the ZXAM6000 alloy sheet keeps its moderate ductility of ~10% even after the aging. This is because the fine grain structure of the ZXAM6000 alloy sheet could suppress the twinning deformation (Fig. 9) [39]. It has been demonstrated that the refinement of precipitates improves both strengths and ductility in age-hardened Mg alloys [10,43,44], and Al–Mn precipitates in Mg–Zn based alloy sheets may act as preferential nucleation sites for β1’ phases [16,17] (Figs. 7 and 8). Therefore, further enhancement of the strengths and ductility in the Mg–6Zn-0.2Ca (wt.%) alloy sheet can be possible via optimization of Al and Mn contents.
Fig. 9. Secondary electron (SE) and Backscattered electron (BSE) images of fracture surfaces after a tensile test of solution-treated and peak-aged (a, b, c, d) Mg–6Zn-0.2Ca and (e, f, g, h) Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheets stretched along the rolling direction. S.T. and P.A. in Fig. 9 represent solution-treated and peak-aged conditions, respectively.
micro-alloyed Mg–6Zn (wt.%) based alloy sheets [10], further study will be required on the texture development of the Mg–Zn–Ca based alloy sheets to unveil the texture weakening mechanism obtained in this work. Grain refining also contributes to enhanced RT stretch formability in Mg alloy sheets [38,39]. The additions of Al and Mn refine the average grain size of the ZX60 alloy sheet from 32 μm to 15 μm without a change of texture feature and activity of basal slips (Table 3). Therefore, the slightly improved I.E. value in the ZXAM6000 alloy sheet is attributed to the finer grain structure due to the additions of Al and Mn. The fine grain size may be attributed to particle pinning of Al–Mn particles and/or precipitates [12,16,17,40], and it is reported that grain refining will also contribute to enhanced tensile properties [41,42]. As shown in Fig. 3, the PS of the ZX60 alloy sheet along the RD and TD are increased from 120 MPa to 155 MPa and 90 MPa to 120 MPa by the additions of Al and Mn. Also, the additions of Al and Mn improve ductility of the ZX60 alloy sheet, so the solution-treated ZXAM6000 alloy sheet exhibits large elongation to failure over 20% in both RD and TD, which are ~1.5 times higher than those of the ZX60 alloy sheet. As shown in Fig. 4 and Table 3, the texture feature and activity of basal slips in both ZX60 and ZXAM6000 alloy sheets are the same, the enhanced tensile properties in the ZXAM6000 alloy sheet are attributed to the fine grain structure. Although the both alloy sheets have planar yield anisotropy: the PS along the TD is lower than that for the RD due to ease activation of basal slips along the TD (Table 3), it is worth mentioning
5. Summary In this work, we have successfully improved room-temperature stretch formability and tensile properties of a Mg–6Zn-0.2Ca (wt.%) alloy sheet via micro-alloying of low-cost Al and Mn. A Mg–6Zn-0.2Ca0.1Al-0.1Mn (wt.%) alloy sheet forms fine grain structure with an average grain size of 15 μm and a unique texture feature with splitting of (0001) planes towards the rolling and transverse directions. Such microstructural feature contributes to a good Index Erichsen value of 7.2 mm, large elongation to failure over 20%, and moderate 0.2% proof stress of 155 MPa along the rolling direction. After an aging treatment, high number density of rod-shaped β1’precipitates are distributed within 6
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the Mg matrix. This leads to a significant improvement of the 0.2% proof stress of the Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet; the peakaged alloy sheet exhibits the high 0.2% proof stress of 264 MPa along the rolling direction. Also, the peak-aged Mg–6Zn-0.2Ca-0.1Al-0.1Mn (wt.%) alloy sheet shows the 0.2% proof stress of 209 MPa even along the transverse direction. These results indicate that Al and Mn microalloyed Mg–Zn–Ca based alloy sheets will become promising candi dates in automotive industries where the room-temperature formability, strengths, and ductility are important properties.
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Data availability All data are available from the corresponding author on reasonable request. 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. Acknowledgement This work was supported by JSPS KAKENHI Grant Numbers JP19K15321, JP18H03837, and Advanced Low Carbon Technology Research and Development Program(ALCA), 12102886. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2019.138690. References [1] J. Hirsh, T. Al-Samman, Superior light metals by texture engineering: optimized aluminum and magnesium alloys for automotive applications, Acta Mater. 61 (2013) 818–843. [2] K. Iwanaga, H. Tashiro, H. Okamoto, K. Shimizu, Improvement of formability from room temperature to warm temperature in AZ-31 magnesium alloy, J. Mater. Process. Technol. 155–156 (2004) 1313–1316. [3] S.M. Hirth, G.J. Marshall, S.A. Court, D.J. Lloyd, Effects of Si on the aging behaviour and formability of aluminium alloys based on AA6016, Mater. Sci. Eng. A 319–321 (2001) 452–456. [4] D.-W. Kim, B.-C. Suh, M.-S. Shim, J.H. Bae, D.H. Kim, N.J. Kim, Texture evolution in Mg-Zn-Ca alloy sheets, Metall. Mater. Trans. A 44 (2013) 2950–2961. [5] Y. Chino, T. Ueda, Y. Otomatsu, K. Sassa, X. Huang, K. Suzuki, M. Mabuchi, Effects of Ca on tensile properties and stretch formability at room temperature in Mg-Zn and Mg-Al alloys, Mater. Trans. 52 (2011) 1477–1482. [6] S.J. Park, H.C. Jung, K.S. Shin, Deformation behaviors of twin roll cast Mg-Zn-X-Ca alloys for enhanced room-temperature formability, Mater. Sci. Eng. A 679 (2017) 329–339. [7] J. Bohlen, M.R. Nürnberg, J.W. Senn, D. Letzig, S.R. Agnew, The texture and anisotropy of magnesium–zinc–rare earth alloy sheets, Acta Mater. 55 (2007) 2101–2112. [8] S. Yi, J. Bohlen, F. Heinemann, D. Letzig, Mechanical anisotropy and deep drawing behaviour of AZ31 and ZE10 magnesium alloy sheets, Acta Mater. 58 (2010) 592–605. [9] B.-C. Suh, J.H. Kim, J.H. Hwang, M.-S. Shim, N.J. Kim, Twinning-mediated formability in Mg alloys, Sci. Rep. 6 (2016) 22364. [10] T. Bhattacharjee, B.-C. Suh, T.T. Sasaki, T. Ohkubo, N.J. Kim, K. Hono, High strength and formable Mg–6.2Zn–0.5Zr–0.2Ca alloy sheet processed by twin roll casting, Mater. Sci. Eng. A 609 (2014) 154–160. [11] M. Qian, A. Das, Grain refinement of magnesium alloys by zirconium: formation of equiaxed grains, Scr. Mater. 54 (2006) 881–886. [12] T. Nakata, C. Xu, R. Ajima, Y. Matsumoto, K. Shimizu, T.T. Sasaki, K. Hono, S. Kamado, Improving mechanical properties and yield asymmetry in high-speed extrudable Mg-1.1Al-0.24Ca (wt%) alloy by high Mn addition, Mater. Sci. Eng. A 712 (2018) 12–19. [13] P.M. Kelly, A. Jostons, R.G. Blake, J.G. Napier, The determination of foil thickness by scanning transmission electron microscopy, Phys. Status Solidi 31 (1975) 771–780. [14] P.M. Jardim, G. Sol� orzano, J.B.V. Sande, Precipitate crystal structure determination in melt spun Mg-1.5wt%Ca-6wt%Zn alloy, Microsc. Microanal. 8 (2002) 487–496.
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