Trans. Nonferrous Met. Soc. China 30(2020) 647-656

Influence of trace Ca addition on texture and stretch formability of AM50 magnesium alloy sheet

Hong-liang ZHAO¹, Yun-xiao HUA¹, Xiang-lei DONG¹, Hui XING², Yan-li LU³

1. College of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, China;

2. The MOE Key Laboratory of Material Physics and Chemistry under Extraordinary Conditions, Northwestern Polytechnical University, Xi’an 710129, China;

3. State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, China

Received 13 June 2019; accepted 17 December 2019

Abstract: The AM50, AM50-0.1Ca, AM50-0.3Ca and AM50-0.5Ca (wt.%) alloys were hot-rolled and their mechanical properties were determined for the purpose of investigating the effect of trace Ca addition on the texture and stretch formability of AM50 alloy. The results show that the addition of trace Ca can effectively modify the basal texture, which is characterized by the split of basal poles deviated from the normal direction (ND) after the hot rolling, while a broad spread of the basal planes toward the transverse direction (TD) after the annealing. Such change of the basal texture is related to the prior formation of massive compression twins and the decrease of the c/a ratio. Erichsen value increases from 2.25 to 4.21 mm with the increase of Ca content. The enhancement of stretch formability is ascribed to the weakened basal texture, which results in the increase of n-value and the decrease of r-value.

Key words: magnesium alloy; Ca addition; texture; stretch formability; mechanical properties

1 Introduction

Wrought magnesium alloy sheets have gained extensive attentions from the automotive industry owing to their better processing properties in comparison with those cast ones [1-3]. Compared to the traditional commercial wrought magnesium AZ31 alloys, the AM series alloy such as AM50 alloy has the unique advantages of higher strength and comparatively favorable ductility [4,5], which exhibits a promising potential for wrought magnesium materials. However, the poor stretch formability, due to the formation of strong basal texture during the deformation process has limited their further applications. Therefore, modifying the basal texture to enhance the stretch formability of the Mg alloy sheet presents great scientific and industrial significance.

Microalloying is an effective method to modify the basal texture, which can be easily achieved by adding trace amounts of alloying elements into the matrix alloys [6]. It has been reported that the trace addition of rare earth elements such as Y, Ce and Gd can effectively contribute to the weakened basal texture and thus result in the enhanced stretch formability [6-8]. But, these expensive rare earth elements substantially increase the costs of raw materials. It is, therefore, essential to find other cost-effective alternatives. Ca, as a low-cost alloying element, can not only refine the microstructure of magnesium alloy, but also has the similar effect on the texture weakening compared to rare earth elements [9-13]. However, the mechanism on the texture control is still not clear and the study on the stretch formability by adding trace Ca individually into Mg alloys still obtains few attentions. Therefore, it can be provided an opportunity for the development of the low-cost Mg alloy sheets with Ca addition based on the AM50 alloy.

In this study, the alloys with different Ca additions have been prepared for evaluating the effect of Ca on the texture and stretch formability of AM50 alloy sheet. The microstructure evolution during the hot rolling, and the basal textures of the as-rolled and annealed alloy sheets were investigated. The corresponding mechanical properties and stretch formability were determined. This work is expected to provide the theoretical and technological guidance for the development of wrought magnesium alloy sheets in automobile industry.

2 Experimental

Experiments were started on the preparation of the investigated alloys, with the chemical compositions given in Table 1. The melting process was carried out in a crucible resistance furnace under the protective atmosphere (SF₆+CO₂) at 740 °C. Mn and Ca were added in form of Mg-15wt.%Mn and Mg-30wt.%Ca master alloys. After being completely melted, the alloys were cast into a preheated steel mold with the sizes of 150 mm × 120 mm × 9 mm. The ingots for rolling were machined into the samples with the dimensions of 80 mm × 80 mm × 8 mm and then homogenized at 420 °C for 24 h. Rolling process was carried out at 400 °C and held at this temperature for 10 min between passes, and the final thickness of the sheets was 1 mm with a total cumulative reduction of 87% over 13 passes. Finally, the rolled sheets were annealed at 350 °C for 1 h.

Table 1 Chemical compositions of investigated alloys (wt.%)

The metallographic samples were sectioned, cold-mounted, polished and then etched in the picric acid solution (2.1 g picric acid, 10 mL acetic acid, 10 mL distilled water, and 70 mL ethanol) for the microstructure observation on optical microscope. Electron back scattered diffraction (EBSD) analysis were performed on a Zeiss/Auriga FIB SEM machine operating at 15 kV to reveal the textures of the as-rolled and annealed alloy samples. The Ultima IV X-ray diffractometer using Cu K_α radiation was employed to measure the lattice parameters of the powder alloy samples. The tensile specimens with 1 mm in gauge thickness, 20 mm in gauge length and 10 mm in gauge width were machined from the annealed sheets with the angles of 0°, 45° and 90° between the tensile direction and the rolling direction. The Shimadzu/ AG-X plus electronic universal testing machine was used for the measurement of the mechanical properties at an initial strain rate of 1×10^-3 s^-1 at room temperature. The strain hardening exponent values (n-values) were obtained by the data of tensile tests with a strain interval from 5% to 15% and the Lankford values (r-values) were measured on the specimens deformed at a plastic strain of 10%. The samples for Erichsen tests were machined from the annealed sheets with the dimensions of 60 mm × 60 mm × 1 mm. Erichsen tests were carried out at room temperature on the CB2-60D cup convex testing machine using a hemispherical punch with a diameter of 20 mm, the speed and the blank holder force of which were 5 mm/min and 10 kN, respectively. The Erichsen values (IE), which were the punch stroke at the fracture initiation, were directly obtained from the data acquisition system. The average values were given by averaging the IE from three repeated experiments.

3 Results and discussion

3.1 Microstructures during hot rolling

The microstructures of the four alloy sheets on the RD-ND plane during hot rolling are shown in Fig. 1, with the cumulative total deformation of 29%, 47%, 66% and 87%. Generally, there exist two types of twins during the hot deformation: the tensile twins with lenticular morphology and the compression twins with narrow banded appearance [14,15]. It is observed that the tensile twins are apparent and the recrystallization has occurred in some twin regions in the early deformation stages of the AM50 alloy, but there still exist a large proportion of original grains. By comparison, massive narrow banded compression twins appear and interlace with each other in the Ca-containing alloys. Original grains are divided into many long and block regions, and few recrystallized grains can be observed in these intersection areas. With the increase of thickness reduction, the recrystallization processes occur remarkably in these alloys, showing that the deformation twins disappear, while the number of recrystallized grains increases rapidly. It is obvious that the shear bands with black stripe structures appear at the reduction of 66% in the AM50 and AM50-0.1Ca alloys, whereas the higher proportion of shear bands form at a lower reduction of 47% in the AM50-0.3Ca and AM50-0.5Ca alloys. These deformed shear bands continue to broaden in the further reductions, which gives rise to the finer and more homogeneous microstructures after the last pass rolling.

Fig. 1 Microstructures of alloys on RD-ND plane during hot rolling

Generally, owing to the minimum critical resolved shear stress (CRSS), the tensile twins are easy to occur in the initial stages of the hot deformation [16]. But, in the present study, the compression twins preferentially appear during the hot rolling in Ca-containing alloys. This indicates that the addition of trace Ca has changed the twinning behavior of the AM50 alloy. The dominant mechanism of twinning is highly dependent on the stacking fault energy (SFE) [17], which is closely related to the difference of the atomic size. The greater difference between alloying elements and Mg atomic radius tends to obtain the lower SFE [18]. For example, the addition of Y element has promoted the formation of compression twins and - secondary twins, which was ascribed to the reduction of SFE [19]. In our study, the radii of Al and Mn atoms are both smaller than radius of Mg atom, while Ca has the atomic radius equivalent to Y, as well as the largest atomic radius difference compared with other elements. This suggests that the addition of Ca into AM50 alloy can significantly induce the local lattice distortion and the consequent decrease of the SFE. Therefore, the prior formation of compression twins after Ca addition can be attributed to the decrease of the SFE.

The shear bands are apparent with various proportions at different reductions, which is related to the formation of different types of deformation twins. Compared to the tensile twins, the massive compression twins in Ca-containing alloys are more favorable for basal slip [14]. Thus, the severe stress concentration could be produced, since the dislocation slips are usually blocked at twin boundaries. In order to coordinate the further deformation, the recrystallization is accelerated and the numerous new twins will be formed to provide more space for basal slip. The deformed shear bands are most likely to originate from these intensive new twins. Therefore, the prior formation of the shear bands in AM50-0.3Ca and AM50-0.5Ca alloys can be attributed to the higher proportion of the compression twins. These deformed shear bands consisting of massive deformation twins can effectively facilitate the recrystallization processes, which leads to the more homogeneously-distributed fine recrystallized grains.

3.2 Textures of rolled and annealed alloy sheets

Fig. 2 (0002) pole figures, texture intensity as function of tilt angle of (0002) plane deviations from ND, and misorientation angle distributions of as-rolled alloys at reduction of 87%

The (0002) pole figures, the texture intensity as a function of the tilt of (0002) plane deviates from ND, and the misorientation angle distributions of the as-rolled alloy sheets at the reduction of 87% are shown in Fig. 2. The strong basal texture with a typical characteristic, <0001>//ND is observed in AM50 alloy, whereas the weakened basal texture with the slight tilt of basal poles can be found in the case of the Ca-containing alloys. The texture intensity remarkably decreases from 21.79 to 12.11 with the increase of Ca content. It is worth noting that the double-peak texture is apparent in AM50-0.3Ca alloy with the peak intensity points lying at ~4° and ~12°. Similarity, the split of basal pole also appears in the AM50-0.5Ca alloy. It can be seen that the tilt angles of most grains are still within 10°, but there exists a strengthening point at ~26°. The fraction of low-angle grain boundary (LAGB) decreases while that of the high angle grain boundary (HAGB) increases with the increase of Ca addition, which results in the average misorientation angle increasing from 11.30° to 19.58°. The higher content of HAGB in Ca-containing alloys suggests the increased orientations among the grains, indicating the accelerated recrystallization during the deformation process. Therefore, it can be considered that the addition of Ca can effectively promote the recrystallization and thus lead to the basal texture weakening.

During the recrystallization process, the compression twins are the more effective nucleation sites than the tensile twins due to the higher internal storage energy. New grains nucleate at compression twins mainly toward the orientations of the sub-grains, which have been subjected to the complicated orientation rotations [14]. Therefore, the accelerated nucleation at massive compression twins in the initial stage of the deformation is believed to be responsible for the texture weakening. With the increase of the reduction, the recrystallization is intensified and the shear bands become the main nucleation sites of the recrystallized grains. Previous studies have been reported that the shear bands often have basal planes parallel to the shear plane, which usually leads to a wide range of recrystallization orientations [19-21]. This suggests that the split of basal pole in the AM50-0.3Ca and AM50-0.5Ca alloy can be related to the higher proportion of the shear bands.

In addition, the texture weakening is also relevant to the change of the lattice parameters. For the metals with close-packed hexagonal structure, the non-basal slips are more likely to be activated when the alloy has a smaller axial ratio (c/a). It has been studied that the addition of RE element resulted in the decrease of c/a, which contributed to the activation of non-basal slip and then weakened the basal texture [22]. As mentioned above, the radii of Al and Mn atoms are both smaller than the radius of Mg atom while Ca has the larger atomic size with respect to Mg. Therefore, it is suggested that the solid solution of Ca atoms can lead to the lattice distortion and the consequent decrease of c/a ratio. The XRD patterns of four investigated alloys are shown in Fig. 3. The corresponding calculated c/a values are summarized in Table 2. It is observed that compared to the pure Mg, the AM50 alloy exhibits the higher c/a despite the lower a and c values. With the increase of Ca content, the a value increases while the c value decreases, resulting in the c/a ratio decreasing from 1.6245 to 1.6198. This implies that the decrease of c/a also plays an important role in the texture weakening of Ca- containing alloys after the hot rolling.

Fig. 3 XRD patterns of alloys for calculation of c/a ratio

Table 2 Lattice parameters and axial ratios of alloys

The orientation maps and (0002) pole figures, as well as the distributions of grain size and misorientation angle of the annealed alloys are shown in Fig. 4. It can be observed that the average grain size and texture intensity decrease from 12.52 μm to 7.87 μm and from 17.98 to 7.29, respectively, but the average misorientation angle increases from 32.07° to 38.69° with the increase of Ca content. Compared to the alloys in as-rolled condition, the textures of four annealed alloy sheets are further weakened and the fraction of HAGB increases significantly. This suggests that the static recrystallization is promoted during the annealing. In addition, it can be seen that the typical rolling textures are still retained for the AM50 and AM50-0.1Ca alloys, whereas the split of basal poles in AM50-0.3Ca and AM50-0.5Ca alloys are replaced by the single-peak textures with the orientations spreading to TD. As mentioned above, the AM50-0.3Ca and AM50-0.5Ca alloys obtain a portion of deformed structures after the hot rolling. Therefore, the broad spread of the basal planes toward TD can be attributed to the full recrystallization in those fine-grained regions, where the enough storage energy for the nucleation and grain growth can be provided.

3.3 Mechanical properties and stretch formability

The nominal stress-strain curves of the annealed alloy sheets are shown in Fig. 5. Compared to the AM50 alloy, three Ca-containing alloy sheets exhibit the smaller in-plane anisotropy counterpart, which is supported by the closer curves in three tensile directions. The mechanical properties such as the 0.2% proof stress (σ_s), elongation to failure (δ), strain hardening exponent (n-value), and Lankford value (r-value) in the RD, 45° and TD tensile directions, as well as their average values (X^-=(X_RD+2X_45°+X_TD)/4) are listed in Table 3. It is known from the Hall-Petch formula that the material with smaller average grain size often obtains the higher yield strength. However, the average yield strength decreases from 201 to 180 MPa with the increase of Ca addition in spite of the average grain size decreasing from 12.52 to 7.87 μm. This indicates that the effect of texture weakening plays an important role in the variation of the yield strength. Moreover, the highest σ_s appears in the TD direction in AM50 alloy while the lowest σ_s in this direction is observed in AM50-0.3Ca and AM50-0.5Ca alloys. It can be ascribed to the spread of the (0002) orientation towards TD which is favorable for the basal slip during the tensile deformation.

Fig. 4 Orientation maps and (0002) pole figures, as well as distributions of grain size and misorientation angle of annealed alloys

Fig. 5 Nominal stress-strain curves in RD, 45° and TD tensile directions of annealed alloy sheets

Table 3 Tensile properties in RD, 45° and TD tensile directions of annealed alloy sheets

However, the change of the elongation appears to be relatively insensitive to the weakened basal texture. Specifically, the average elongation decreases from 23.9% to 21.7% with the increase of Ca content. As mentioned above, the addition of Ca has promoted the formation of the shear bands, where numerous deformation twins are concentrated. Currently, the micro-cracks tend to occur in the twin cells or the twin interlaced regions [23,24]. Therefore, the decreased elongation could be ascribed to premature formation of the cracks in extensive twins during the tensile deformation.

The n-value reflects the ability of the alloy sheet to resist the uniform plastic deformation. The large n-value is known to lead to a reduction in plastic instability, resulting in the enhanced stretch formability. In the present study, the n-values in the three tensile directions gradually increase and the average n-value increases from 0.17 to 0.26 with the increase of Ca addition. Besides, the largest n-values of AM50 and AM50-0.1Ca alloys appear in the RD direction, while the largest n-values of AM50-0.3Ca and AM50-0.5Ca alloys are observed in the TD direction. This results from the texture modification. The weakening of basal texture can restrict the softening behavior of the dynamic recovery, which is beneficial to the improvement of the work-hardening ability [25]. For the alloys with 0.3 wt.% and 0.5 wt.% Ca, the basal slip can be facilitated due to the lower texture intensities and the split of basal poles. This contributes to the enhancement of the plastic deformation capacity and therefore results in the larger n-value in the TD direction.

The r-value is used to evaluate the deep drawability of the alloy sheet. The small r-value is beneficial to the sheet thinning which is favorable for the enhancement of stretch formability [26,27]. As shown in Table 3. The r-value decreases in the RD and 45° directions whereas the r-value decreases firstly and then has a little increase in the TD direction with the increase of Ca content. The average r-value decreases from 3.10 to 2.46 and the AM50-0.3Ca and AM50-0.5Ca alloys have the equivalent values. During the in-plane tensile deformation, the width strain and the thickness strain are essentially required due to the fact that the sheet is in a biaxial tensile stress state. It has been reported that the prismatic slip always dominates in the width strain, while the thickness strain is usually governed by the pyramidal slip and twinning [28]. The deformation in the thickness direction of AM50 alloy sheet is difficult to occur due to the strong basal texture, which restricts the pyramidal slip and twinning, and thus results in the larger r-value. However, the TD-inclined basal texture leads to the more complicated grain orientations in the case of Ca- containing alloys. This contributes to the activation of the pyramidal slip [29]. Moreover, the tilt of basal plane can promote the c axis parallel to the tensile direction, which makes it easier to produce the tensile twins during the tensile deformation and therefore results in the decrease of r-value.