J. Cent. South Univ. (2014) 21: 2984−2990 DOI: 10.1007/s11771-014-2266-z
Microstructure and mechanical properties of AZ31 alloy ingot fabricated by semi-continuous casting LI Jin-zhu(李金柱)1, 2, NONG Deng(农登)2, ZHENG Kai-hong(郑开宏)2, LI Xiao-hui(黎小辉)2, ZHAO Ming-chun(赵明纯)1 1. School of Materials Science and Engineering, Central South University, Changsha 410083, China; 2. Institute of Materials Processing, Guangzhou General Research Institute for Industrial Technology (Guangzhou Research Institute of Non-ferrous Metals), Guangzhou 510650, China © Central South University Press and Springer-Verlag Berlin Heidelberg 2014 Abstract: The AZ31 alloy ingot with diameter of 110 mm and length of 3500 mm was fabricated successfully. The compositions and microstructure morphologies of the ingot at different locations were performed, which indicated that the chemical composition distributed homogeneously through the whole alloy ingot and the average grain size increased from the surface to the center. The results of the EDS and element face-scanning illustrated that the eutectic compounds mainly consisted of β-Mg17Al12 and a small amount of β-Mg17(AlZn)12. Furthermore, slight improvements of the strength and ductility were observed from the center to the surface along the axial direction of the alloy ingot, while both the strength and elongation to failure of the samples along the radial direction are higher than that along the axial direction. The fine grain strengthening was the main contributors to the strength of the as-casted AZ31 alloy. Key words: semi-continuous casting method; microstructure; mechanical properties; fine grain strengthening
1 Introduction With the dramatically increased emphasis on mass reduction of the next generation automobiles, aerospace and 3C industries, magnesium and its alloys are receiving much more attention. Low density and excellent castability, high specific stiffness and good damping capacity of magnesium were increasingly used in various applications [1−5]. What’s more, because of its low cost, high corrosion resistance and other perfect characteristics, AZ magnesium alloys have become one of the most popular materials to automobile manufacturers for development and research [6−8]. However, the Mg-Al magnesium alloys have a wide range of solid−liquid dual phase region relatively and close-packed hexagonal crystal structure, which result in the inhomogeneous diffusion of the alloying elements in the magnesium matrix, and even forming dendritic segregation and non-equilibrium phase terribly [9−10]. More specifically, most of the processes for producing magnesium parts is die casting technology. In the future, the die casting process may not meet a new demand for lighter, thinner, stronger, and more complex shape
applications [11]. That is why more and more researchers investigate various methods to obtain high-quality and large scale magnesium alloy billets, which is the premise of high-quality products. Semi-continuous casting technology may be a novel and promising approach to produce magnesium alloy products, exhibiting strong continuity of the production process, reducing chemical components segregation, casting shrinkage cavities and porosities [12]. HAN et al [13] pointed that the grain refinement of AC52 alloy ingot achieved the best conditions with a high cooling rate about 20−65 °C/s. MOSTAFA [14] studied the effects of continuous casting speed on the secondary dendrite arms, which demonstrated that the secondary dendrite arms dropped from 46.7 μm to 40.5 μm, when the casting speed increased from 50 mm/min to 90 mm/min. HUA et al [15] showed that the low pouring temperature would cause elements enrichment (e.g., Ce, Mn, et al), thus reducing the melt flow properties. However, only a few have researched on the casting process of magnesium alloy [16]. In order to enhance the application of magnesium alloy, it is necessary to be familiar with its casting property, especially microstructures and mechanical prosperities of magnesium alloy ingot. In the present
Foundation item: Project(2010A090200078) supported by the Special Foundation Project of Industry, University and Research Institute Collaboration of Guangdong Provincial Government and the Ministry of Education, China; Project(2010B090500010) supported by the Special Commissioners’ Workstation Construction Project of Guangdong Provincial Government, China Received date: 2013−04−07; Accepted date: 2013−11−06 Corresponding author: LI Jin-zhu; Tel: +86−13416199302; Fax: +86−20−37238039; E-mail:
[email protected]
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study, semi-continuous casting method was employed to produce the AZ31 magnesium alloy large ingot with the diameter of 110 mm and length of 3500 mm. The elements distribution and the phases in the different parts of the ingot as well as the relationship between the microstructure and mechanical properties were discussed, which will benefit to gating system design and improve the quality of magnesium alloy castings.
2 Experimental The chemical compositions (mass fraction, %) of raw materials AZ31 magnesium alloy used in this work, which is a commercially available wrought magnesium alloy, are listed in Table 1. They were remelted in an electric resistance furnace with the capacity of approximately 500 kg and protected by a shielding mixed atmosphere of CO2 and SF6 with the ratio of 100:1 at 670 °C, then standing for 40 min. The melt was transferred to the semi-continuous casting machine around 630 °C, the rate of the water in the cooling mould was 3000 L/min and the casting machine worked with the speed of 90 mm/min, respectively. The large AZ31 alloy ingot with the diameter of 110 mm and length of 3500 mm was successfully prepared. Table 1 Chemical composition of raw materials AZ31 magnesium alloy (mass fraction, %) Al
Zn
Si
Fe
Mg
2.71
0.78
0.020
0.004
Balance
The chemical compositions (mass fraction, %) of the alloy ingot at different positions (surface or center) were carried out by ICP analyzer (JY Ultima2); at least 10 points were measured to reduce the error. The samples for microscopy observation were mechanically polished and then etched with 2% aqueous oxalic acid solution plus 2 mL nitric acid, which was observed by Olympus optical microscopy (OM) and JOEL JSM-820 scanning electron microscope. The average grain size was measured by the average linear intercept method. What’s more, the quantitative analysis of the phases chemical compositions in the alloy ingots were performed with
energy dispersive X-ray spectrometer (EDS) equipped on JOEL JSM-820. And the wave dispersion/energy dispersion combined micro analyzers were used to make element area-scanning. The positions and the abridged general view of the tensile samples as well as the tension directions are illustrated in Fig. 1, respectively. And the tensile tests based on the International Standard of Tensile Test ISO6892-84 was carried out at ambient temperature using a DNS 200 testing machine with a cross-head speed of 2 mm/min. The tensile samples along axial directions for the test were cut from the center and near the surface of the alloy ingot by electrical discharge machining. And because of the limit of size, the tensile samples along radial directions were just only cut from the center of the alloy ingot. For each sample, three specimens were tested to calculate the average value as the final results.
3 Results and discussion 3.1 Microstructure The examined chemical compositions of the alloy at different locations of the ingot are listed in Table 2, which indicates that there was no macro segregation of the solute elements and the homogeneous distribution of chemical composition was obtained. Compared with the chemical composition of raw materials as listed in Table 1, the contents of the main elements (e.g., Al, Zn, etc) were reduced to a certain degree. This is because of the burning losses in the remelting process. But that of the Fe element increased slightly. The reason is that the Fe in the iron crucible fused into the melt at high temperature, which was higher than the burning losses in the remelting process. What’s more, the elements contents near the surface of the alloy ingot are slightly higher than that of the center, especially Al and Zn. This should be attributed primarily to the rapid cooling rate to the alloy surface during the semi-continuous casting, which caused the elements have no enough time to diffusion, thus resulted in the super-saturated solid solution. Figure 2 shows the microstructures of the asreceived ingot during the semi-continuous casting process observed at different positions. As can be seen
Fig. 1 Tensile samples abridged general view (a) and positions and tension directions (b) (Unit: mm)
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Mass fraction/% Center
Surface
Al
2.64
2.69
Zn
0.76
0.769
Si
0.019
0.018
Fe
0.007
0.009
Mg
Balance
Balance
Fig. 2 Microstructures of as-received ingot at different observation positions: (a) Center; (b) Surface
from these pictures, the whole alloy ingot was composed of α-Mg matrix, which is composed of coarse equiaxed grains, and a network of the discontinuous eutectic compounds on grain boundaries. Among other things, the dispersed distribution of needle-like precipitates is a typical casting divorced eutectic structure. This should be attributed primarily to the temperature gradient in the alloy ingot during the solidification process. The gradient was small combined with the adsorption of the atom clusters at high temperature that made the grain grew. Furthermore, the segregated Al and Zn elements at the grain boundaries formed the eutectic compounds. Generally, according to the Mg-Al diagram, the maximum equilibrium solid solubility of Al in Mg matrix is relatively 12.17% at 708.65 K (eutectic temperature) and decreases exponentially with decreasing temperature (2%) at 373.15 K. The content of Al in AZ31 alloy which has a wide range of solid−liquid two-phase region relatively is about 3%, and the α-Mg is close-packed
hexagonal crystal structure, thus the solid solution of solute atoms in the matrix was reduced, and the divorced eutectic was formed during the non-equilibrium solidification process. What’s more, β-Mg17Al12 being body-centered cubic structure presented in the eutectic organization and distributed in the α-Mg grain boundaries of the earlier precipitation. The average grain sizes at the center and near the surface were 256.3 μm and 182.6 μm, respectively. The average grain size decreased from the center to the surface of the ingot, while the volume fraction of the eutectic compounds increased obviously. Because of the rapid cooling rate at the alloy ingot surface, the solute segregation was so high that the average grain size was small and the volume fraction of the eutectic compounds was high. What can be seen from the SEM micrograph of the as-received alloy ingot (Fig. 3(a)) is that the eutectic compounds distributed homogeneously and discontinuously in the observed region. As shown in Fig. 3(b) which was the magnified image of the region marked B in Fig. 3(a), the eutectic compounds presented rod-shaped or small blocky-shaped and a small amount of that exhibited dot shaped. Obviously, it is the grain boundaries that the compounds surrounded in. Figure 3(c) revealed the further magnified SEM image of the region marked C in Fig. 3(b). Some of the eutectic compounds extended into the interior of the α-Mg matrix. What’s more, the energy dispersive X-ray spectrometer (EDS) was utilized to exam the compositions of the rod-shaped precipitation of the region marked D in the Fig. 3(c), and the results were shown in Fig. 3(d). It can be seen from the EDS analysis that the white block precipitates is the compounds of the β-Mg17Al12 or β-Mg17(AlZn)12 and a small amount of AlMn phase formed at the grain boundaries, which is the black spots on the block precipitates. The elements area-scanning results shown in Fig. 4 indicate that the Mg, Al and Zn elements were enriched in the region of the white block precipitates in Fig. 4(a). In addition, there were a small amount of Mn elements distributed in the block phase. Obviously, the results of the elements area-scanning test are consistent with EDS results. This is attributed to the fact that the initial β-Mg17 (AlZn)12 was formed by the eutectic reaction of the α and β phases during the solidification process. Furthermore, many investigations showed that the addition of Zn element is not only beneficial to the decreases of the solidification temperature, but also conducive to promote the precipitation of β-Mg17Al12 on the grain boundaries [17−19]. As the solidification process proceeds, the replacement of Al by Zn descended. At the end of solidification process, the Zn element of β phase would be replaced by Al element in the subsequent cooling process. Thus, the stable β- Mg17Al12 was formed.
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Fig. 3 SEM micrographs of the as-received alloy ingot: (a) Alloy ingot; (b) Magnified image of region marked B in (a); (c) Magnified image of region marked C in (b); (d) EDS results of phases marked D in (c)
3.2 Mechanical properties and fracture analysis Representative tensile stress−strain curves of the as-casted alloys at different locations along axial and radial directions at ambient temperature are shown in Fig. 5(a). The tensile mechanical properties including ultimate tensile strength (UTS), tensile yield strength (TYS) and elongation to failure (ε) of the casted AZ31 alloy at room temperature are listed in Table 3. It can be seen that, in terms of the axial direction of the as-casted AZ31 alloy, compared with the center, slight improvements of UTS and TYS as well as the elongation to failure are observed in the surface of the AZ31 alloy ingot, by 20 MPa, 7 MPa and 1.6%, respectively. There are several reasons to this phenomenon. Firstly, because of the rapid cooling rate during the semi-continuous casting process, the average grain size of the surface was smaller than that of the center of the as-casted alloy ingot as shown in Fig. 2, which is beneficial to the fine tension performance of the surface of the alloy ingot based on the Hall-Petch relationship [20−21]. At the mean time, it is well known that the eutectic compounds which dispersed at the grain boundaries could impede the dislocations slip. The cooling rate exponentially increased from the center to the surface, which led the volume fraction of the eutectic compounds of the surface would be higher than the center. Therefore, the strengthening effect of the surface was more significant
than the center. Secondly, according to Hume-Rothery theory [22], if the concentration of the solute atoms was higher, and the difference between the solute atoms and matrix atoms was larger, thus the strengthening effect would be more obvious. As far as we know, the diameter of the magnesium atom is larger than that of the zinc atom, so the supersaturated zinc atoms could improve the strength to a certain extent. What’s more, it is the improvement of ductility from the center to the surface of the as-casted alloy ingot that has prompted the work presented here. This characteristic can also be explained by the finer grain size of the surface. Anything else, the smaller measurement of the eutectic compounds of the surface was also conducive to the increase of the elongation. As shown in Fig. 5(a) and Table 3, both the strength and elongation to failure of the samples along the radial direction are higher than those along the axial direction, UTS of 199 MPa, TYS of 69 MPa and the elongation to failure of 11.2%. On the one hand, this characteristic can be explained by the anisotropy produced by the semi-continuous casting. On other hand, the different position from the alloy ingot may also cause this result. It is necessary to gain a better understanding about the reason that the mechanical properties of the samples along the radial direction are higher than that in the axial direction.
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Fig. 4 Part of magnified SEM image of Fig. 3(b) (a) and elements areascanning mapping of Mg (b), Al (c), Mn (d) and Zn (e)
Fig. 5 Tensile stress−strain curves and SEM micrographs of fracture surface after tension tests of as-casted alloys at different locations along axial and radial directions at ambient temperature: (a) Stress−strain curves; (b) SEM images along radial direction; (c) SEM images near center of alloy ingot in axial direction; (d) SEM image near surface of alloy ingot in axial direction
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Table 3 Tension properties of as-casted alloys at different locations along axial and radial directions Position
UTS/MPa
TYS/MPa
Elongation/%
Axial center
162
59
8.7
Axial surface
182
66
10.3
Radial
199
69
11.2
SEM micrographs of the fracture surface after tension tests are shown in Figs. 5(b), (c) and (d), revealing that there are a large of volume fractions of the cleavage planes and platforms both along axial and radial directions, which confirmed that the fracture of the specimens of the alloy ingot was brittle fracture. However, compared to micrographs (Figs. 5(c) and (d)) along the axial direction, much more dimples can be seen in the micrographs of the radial direction (Fig. 5(b)). This is why the ductility of the alloy ingot along the radial direction is better than that along the axial. In regard to the axial direction of the alloy ingot, the fraction of the dimples near the surface is higher than that of the center as shown in Figs. 5(c) and (d), which shows that the surface position exhibits a relatively better elongation. This should be attributed primarily to the fine grains near the surface of the alloy ingot. Firstly, the smaller the grain, the more the total amount of grains with the same area, thus the average stress and strain acting on each grain is relatively small, and the deformation is more uniform [23]. Secondly, grain boundaries would impede dislocations slip and shorten the distance of dislocations slip; when dislocation slipped to the grain boundaries, the dislocation group was formed by the jammed dislocation, which also induces some extrusion pressure on the grain boundaries. With the force increasing to a certain value，the pileup group will thrust the grain boundary [24−25]. In other words, the deformation of crystal proceeds. Therefore, the alloy ingot with fine grains has relatively high mechanical properties. Thirdly, the fine grains are able to release the stress concentration at the grain boundaries to avoid generating micro-cracks. For the reason that the crazes and the breakages occur relatively late, the material can withstand a large amount of deformation before breaking [22]. All of these analytical evidences demonstrate that the fracture of the samples is caused by the dislocations piled up and stress concentration occurs seriously at the center of the alloy ingot with coarse grains.
4 Conclusions 1) The microstructure of as-received AZ31 alloy consists of coarse dendrite grains and discontinuous networks of β-Mg17AL12 or β-Mg17(AlZn)12 eutectic compounds at the grain boundaries and the fine
precipitates dispersed in the crystal. Furthermore, the average grain size decreases from the center to the surface of the ingot, while the volume fraction of the eutectic compounds increases obviously. 2) Different regions of the alloy ingot exhibit different mechanical properties. The UTS of 199 MPa, TYS of 69 MPa and the elongation to failure of 11.2% of the alloy ingot along radial direction at the room temperature are obtained. What’s more, slight improvements of UTS and TYS as well as the elongation to failure are observed in the surface of the AZ31 alloy ingot, by 20 MPa, 7 MPa and 1.6%, respectively. 3) The fine grain strengthening is the largest contributor to the strength of the as-casted AZ31 alloy.
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