Magnesium alloys are the lightest constructional materials owing to their low density, high specific strength, and hardness [1, 2]. Recently, there has been significant increase in the usage of magnesium alloy for aerospace components, automobile, 3C products, etc. [3, 4]. However, commercial applications of magnesium alloys to important structural parts are still limited due to some undesirable properties. The strength and ductility of magnesium alloys are generally poorer than those of competitive aluminium alloys. In order to meet the requirements of more applications, magnesium alloys with high strength and excellent room-temperature properties must be developed. Grain refinement is a promising way to enhance yield strength (YS) based on the Hall–Petch relationship [5, 6] and alloying is an effective grain refinement method for magnesium alloys. Some researches show that improvement in microstructure and mechanical properties of the alloys can be achieved by the addition of minor alloying element [7, 8]. Magnesium is usually alloyed with Al, Zn, Mn, Si, Sn and rare earth (RE) element. In particular, RE elements are important alloy elements to magnesium alloys, which can enhance the binding force of magnesium atoms, and reduce the diffusion velocity of atoms. Meanwhile, RE elements can form the high-melting point compounds with other elements and refine the grains, and thus they can effectively improve the heat resistance of magnesium alloy. There are some literatures that discuss the effect of Y additions on the microstructure and mechanical properties of magnesium alloy. N. Kashefi et al. studied the microstructure and impression creep behaviour of cast AZ80 magnesium alloy with Y additions . Tong et al. reported the effect of RE additions on microstructure and mechanical properties of AZ91 magnesium alloys . Wang et al. studied the microstructure and mechanical properties of AZ91 alloys by addition of Y . Luo et al. reported the effects of Y on microstructure and property of heat-treated AZ91D magnesium alloy prepared by lost foam-casting process .
Recently, our research shows a minor addition of Sn can contribute to the improvement of the mechanical properties of AZ80 magnesium alloys. However, there are few reports on the effects of Y addition on the microstructure and mechanical properties of as-cast Sn-containing AZ80 magnesium alloys. In this paper, microstructure of as-cast AZ80–2Sn alloy by addition of Y has been discussed; the mechanical properties at room temperature were investigated.
High-purity Mg (>99.7%), Al (>99.7%), Zn (>99.6%), Sn (>99.8%), the Mg-5% Mn master alloy and the Mg-30% Y master alloy were used as raw materials. The additions of Y were 0, 0.50, 0.95 and 1.93 mass%. The chemical compositions of the studied alloys are shown in Table 1.
The experimental procedures are described as follows. The alloys were melted in a semicontinuous vacuum induction melting furnace. The molten metal was held for about 20 min at 720℃, then poured into a columniform iron mould with internal diameter of 85 mm and cooled rapidly by water. The melt was protected by a mixture of SF6 and CO2 with the ratio of 1:100 during the melting and pouring stage. Microstructure and morphology were investigated by an optical microscopy (OM, NEISS NEOPHOT-30) and a scanning electron microscopy (SEM, TESCAN VEGA II). The specimens for OM and SEM were prepared by standard technique of grinding with SiC abrasive and polishing with a diamond paste (2.5 μm). The specimens were etched by 4% oxalic acid, for about 10 s before observed. The Oxford X-ray energy-dispersive spectroscopy (EDS) was employed to determine the phase’s chemical constitution of the specimens. The overall phase constitutions of these alloys were analysed by a Rigaku D/max-3C X-ray diffraction (XRD) instrument with Cu Kα and a scanning rate of 0.01°/s.
Tensile tests at room temperature were performed on WE-10 materials test machine at a constant speed of 2 mm/min. The ultimate tensile strength (UTS), YS and elongation to rupture (Er) were the average values of three specimens. Fractographic observation was conducted on the SEM to clarify the fracture process.
Effects of yttrium content on the microstructure
XRD analysis shows that the microstructure of as-cast AZ80–2Sn alloy is mainly consisted of α-Mg phase, β-Mg17Al12 phase and Mg2Sn phase; additional diffraction peaks of Al2Y phase emerge as a result of Y addition (Figure 1). The OM and SEM images of as-cast AZ80–2Sn alloys with different Y contents are shown in Figures 2 and 3, respectively. As shown in Figure 2, the microstructure mainly shows dendrite morphology with the secondary phases distributed in interdendritic spacing and along grain boundaries. As seen from Figures 2(a) and 3(a), as-cast AZ80–2Sn alloy consisted of primary α-Mg matrix and eutectic phase. These eutectic phases were identified to β-Mg17Al12 phase and γ-Mg2Sn phase by XRD and EDS analyses (Figure 4). As a second phase, Mg17Al12 shows semicontinuous and reticulated shape and distributes mainly at grain boundaries. Furthermore, there are numerous Mg2Sn precipitate particles dispersing in Mg17Al12 eutectic phases. Figure 2(b)–(d) shows the metallographic structures of the as-cast alloys by additions of Y as 0.5, 0.95 and 1.93 mass%, respectively. It can be seen that there have been significant effects on the metallographic structure of alloys with Y addition. By adding Y, the amount of β-Mg17Al12 is decreased and the dimension of that is reduced. Mg17Al12 eutectic phase turns to discontinuous, and the more disperse phases occur with the increase of Y content. However, there has no obvious effect on grain refinement with Y addition, which is in agreement with the previous work in the Mg-Zn-Al-Y alloy .
In addition, small polygonal particles are found in alloy 2 under SEM (Figure 3(b)); its chemical formula was identified as Al2Y (2Al + Y = Al2Y) by EDS analysis (Figure 5). As shown in Figure 1, the XRD also shows the weak peaks of Al2Y phase. Further addition of Y element leads to the rapid increment of Al2Y particles in the alloy. Once the content of Y exceeds 0.5 mass%, Al2Y phase particles become more and more coarsening (Figure 3(c)–(d)).
Effects of yttrium content on the room-temperature mechanical properties
The results of the tensile tests of the investigated alloys at room temperature are shown in Figure 6. It can be seen that the tensile strength and elongation of the AZ80–2Sn alloys are improved with the addition of Y. Alloy 2 shows the optimal mechanical property compared with alloy 1, where the tensile strength is improved by 18.65%, but the improvement of YS seems unobvious. The further increment of Y content leads to the decline of UTS and YS. The elongation attains the maximum value of 8.23% when the content of Y reaches to 0.5 mass%, but further increment of Y content leads to the rupture elongation that decreased slightly.
Figure 7 shows the SEM images of tensile fracture morphology of the investigated alloys at room temperature. The fracture pattern of alloy 1 (Figure 7(a)) indicates quasi-cleavage steps and tear ridges, with the quasi-cleavage being the main characteristic. Alloy 1 with the characteristic of brittle fracture has low elongation value. When the content of Y increases to 0.5 mass%, some dimples can be observed on the fracture morphology, as shown in Figure 7(b), and the corresponding elongation is 8.23%. When the content of Y is up to 0.95 mass%, the fracture morphology of the alloy 3 clearly shows a few local dimples, tear ridges and the formation of quasi-cleavage fracture characteristic (Figure 7(c)). When the content of Y is up to 1.93 mass%, the fracture morphology of alloy 4 demonstrates more brittle fracture models with tear ridge and cleavage step, as shown in Figure 7(d). In addition, some coarse polygonal particles can be seen in alloys 3 and 4. Based on the above analyses, it can be found that the fracture mechanism of the investigated alloys has not been changed with the addition of Y, the fracture morphologies mainly consist of some dimples and tear ridges, and some secondary micro-cracks are also observed on the fracture morphology (Figure 7(d)). This demonstrates that the fracture mechanism of Y-containing AZ80–2Sn alloys is quasi-cleavage fracture mode.
XRD analysis and microstructure observations by SEM show that as-cast alloy 1 consisted of α-Mg matrix, β-Mg17Al12 phase and γ-Mg2Sn phase. As a second phase, Mg17Al12 shows semicontinuous and reticulated shape and distributes mainly at grain boundaries. Furthermore, there are numerous Mg2Sn precipitate particles dispersing in β-Mg17Al12 eutectic phases. The as-cast alloy 1 exhibits poor mechanical properties, which can be mainly ascribed to the presence of the coarse network β-Mg17Al12 eutectic phases at grain boundaries. The β-Mg17Al12 phase with a body-centred cubic crystal structure is incoherent with the magnesium matrix with hexagonal close-packed crystal structure, which results in the fragility of the Mg17Al12/Mg interface [14, 15].
Y and Mg have same hexagonal close-packed crystal structure with similar lattice constants, and they have similar atomic radii . The lattice constants of Y and Mg are a = 0.36500, c = 0.57410 nm and a = 0.32094, c = 0.52104 nm, respectively . The atomic radii of Y and Mg are 0.182 and 0.162 nm, respectively. According to the theory of “size and structure matching”, Y can act as heterogeneous nucleation nucleus for the α-Mg phase, which promotes the nucleation rate and refines grains. As a result, Y should have a good effect on grain refinement . But the grain refinement of Y addition on the AZ80–2Sn alloy could not be found in this research, because there were little separate Y phases in alloys 2, 3 and 4. According to the binary alloy phase diagram of Mg–Y, Y and Mg can form the compound of Mg24Y5, in which the content of Y is about 40 mass% [4, 18]. According to the ternary alloy phase diagram of Mg–Al–Y, Al and Y can form compounds of Al2Y and Al3Y, in which the content of Y is about 27% and 32 mass%, respectively [4, 18]. In this research, the Al2Y compound was identified, but any binary compounds of Mg–Y and ternary compounds of Mg–Al–Y were not found. The melting point of Al2Y is 1485℃  and the crystallization temperature of Al2Y (980℃) is much higher than the eutectic reaction temperature of L (450℃) →α-Mg + Mg17Al12. Thus, Al2Y phase will preferentially form and concentrates at the front edge of α-Mg phase. Obviously, during the solidification process, Al2Y phase is difficult to act as heterogeneous nucleation nucleus for the α-Mg phase due to its cubic crystal structure and large lattice constant (a = 0.786 nm). Therefore, this can explain why there has no obvious effect on grain refinement with Y addition.
As shown in Figure 6, the tensile strength and elongation of the AZ80–2Sn alloys are improved with the addition of Y. The improvement of tensile strength is mainly ascribed to the refinement of Mg17Al12 phase and the dispersion strengthening. During the solidification process, a majority of the Al in Y-containing AZ80–2Sn alloy is depleted because of the formation of Al2Y, which can reduce the volume fraction of Mg17Al12 phases. The decrease in the volume fraction of Mg17Al12 phase also leads to the increase in tensile strength [14, 15]. During the process of further cooling, the growth of Mg17Al12 phase is restricted greatly due to the enrichment of Y atoms in solid–liquid interface. Therefore, Mg17Al12 phase is modified from semicontinuous network to fine dispersed phases. The fine Mg17Al12 phases and Al2Y phases can hinder the slippage of dislocation, which can improve the room-temperature properties of experimental alloy. Similarly, the improvement of elongation may be attributed to the refinement of Mg17Al12 phase. Research on deformation mechanism in squeeze-cast AZ91–X (Y, Nd and Sr) alloys indicates that micro-crack originates from the interface between the coarse network Mg17Al12 eutectic phase and α-matrix . By adding Y, the amount of Mg17Al12 is decreased and the dimension of that is reduced. Thus, the elongation of experimental alloy is slightly improved when the content of Y reaches to 0.5 mass%.
However, the mechanical properties of experimental alloys are decreased at room temperature when the content of Y exceeds 0.5 mass%. The decreases are ascribed to the following reasons. When the content of Y exceeds 0.5 mass%, Al2Y phase becomes coarsening. There are large amounts of coarse polygonal Al2Y phases that appear in the alloys and weaken the dispersion strengthening effect. In addition, the coarse polygonal intermetallic phase particles have been found to have significant impact on the micro-cracks formation and the flaws propagation. This point can be supported by the fracture morphology observation of alloys 3 and 4 (Figure 7(c)–(d)); some coarse polygonal particles can be observed lying in the dimple, and these polygonal particles are proved to be Al2Y phase by EDS analysis. This indicates that some micro-cracks originate from the interface between the hard Al2Y phases and α-matrix in the case of severe strains. Thus, all of these factors can contribute to decreased ductility and strength of alloys 3 and 4.
Mg17Al12 eutectic phase in as-cast AZ80–2Sn alloy shows semicontinuous and reticulated shape and distributes mainly at grain boundaries. Furthermore, there are numerous Mg2Sn precipitate particles dispersing in Mg17Al12 eutectic phases. The AZ80–2Sn alloys with variable Y contents all contain Al2Y phase. By adding Y, the amount of Mg17Al12 is decreased and the dimension of that is reduced. But there has no obvious effect on grain refinement with Y addition.
The tensile tests at room temperature indicate that the best comprehensive properties are obtained when the addition of Y is 0.5 mass%. When the content of Y is above 0.5 mass%, the tensile strength and elongation become decreasing.
The improvement of tensile strength is mainly ascribed to the refinement of Mg17Al12 phase and the dispersion strengthening. The decrease is ascribed to the coarsening of Al2Y phases when the content of Y exceeds 0.5 mass%.
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About the article
Published Online: 2015-01-14
Published in Print: 2015-12-01
Funding: This work was financially supported by the National Basic Research Program of China (2013CB632200) and the sharing fund of Chongqing University’s large-scale equipment (20131063009).