Effect of La/Nd ratio on microstructure and mechanical properties of as-cast AZ91-xLa/Nd alloy

The effect of La/Nd ratio on microstructure and mechanical properties of as-cast AZ91D alloys were investigated. The alloys were produced by casting the molten metal into a preheated steel mould. Microstructures were investigated by XRD and SEM (equipped with EDS). The hardness measurement was carried out by Brinell hardness (HB) digital display tester. Compression tests were conducted at both room and high temperatures (20 °C and 150 °C). Microstructural characterizations reveal that the area fraction of the β-Mg17Al12 phase changed with the variation of La/Nd ratio. The area fraction of the β-Mg17Al12 phase decrease significantly and then increased slightly with the increase of La/Nd ratio. Besides, the results show that the addition of La and Nd contributed to the formation of needle-like phase Al11RE3. Meanwhile, the Brinell hardness and high-temperature compression properties of this alloy reach the maximum values (67 HB and 330 MPa, respectively) when the La/Nd ratio is 2/3. In this alloy, the amount of Al11RE3 phase is 51/TA and the average length is 69.55 μm. Moreover, a generation model of Al11RE3 is established and the growth mechanism of Al11RE3 is analyzed. This work thus has proved that La/Nd ratio has a great effect on the microstructure and mechanical properties of AZ91.


Introduction
Magnesium alloys are the most promising lightweight material and are undoubtedly widely used in aeronautical, 3C, and automotive industries [1,2]. The alloys have made significant inroads in interior parts as well as structural components [3,4]. The most significant magnesium alloy application of automobiles utilizes AZ type casting alloys such as Mg-9 wt% Al-1 wt% Zn alloy (AZ91) [5]. AZ91 used in automobile offers good solidification characteristics of excellent fluidity and less susceptibility to hydrogen porosity so it is remarkably easy to cast [6]. Although AZ91 alloy was developed for a suitable combination of room-temperature strength with ductility and corrosion resistance, it's poor high-temperature (>120°C) mechanical properties limit the automotive application of power train parts [7,8]. Hence, it is reasonable to expect that there are high strength and low-cost alloys based on AZ91 alloy to be developed by alloying or micro-alloying [9].
Over the past decade, many investigators have been contributed for the modified-AZ91 alloy by alloying and micro-alloying elements, such as La, Ce, Nd, Y, Ca, B, Sr, Sb, Br, Pr and so on [10]. The results indicate that light rare earth elements (La, Ce and Nd) are more effective than other elements. The effect of rare earth elements to the as-cast microstructure of magnesium alloys is mainly reflected in the following aspects: (i) The addition of light rare earth elements can reduce the solid-liquid interfacial tension. The critical forming radius of the rare earth-containing magnesium alloy is reduced, and it is easy to nucleate, which promotes grain refinement of the alloy [11]. (ii) The partition coefficient of light rare earth in Mg is less than 1, and the solid solubility decreases with the decrease of temperature, so the rare earth element can increase the supercooling degree of the solid/ liquid interface front component during solidification, which promotes the formation of completely divorced eutectic [12]. (iii) After the solidification process, some rare earth elements (like Nd) are dissolved in α-Mg, Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. which reduces the diffusion coefficient of solid solution, making it difficult to diffuse Al in α-Mg solid solution, thus inhibiting the secondary β-Mg 17 Al 12 phase [13]. At the same time, the Al element in the alloy will be consumed during the formation of the stable phase of Al 11 RE 3 , which will reduce the number of eutectic β-Mg 17 Al 12 and obtain finer dispersed particles [14].
Furthermore, Zhang et al [15] studied the mechanical properties of Mg-4Al-4La alloy and concluded that AlLa44 exhibited excellent mechanical properties due to the hindrance to grain boundary sliding and dislocation motion provided by Al 11 La 3 phase. Yang et al [16] reported that needle-like phase Al 11 La 3 was thermodynamically stable in Mg-Al alloy at temperatures lower than 1000 K. Studies [17,18] showed that Mg-4Zn alloy with addition of 1 wt% La exhibited good corrosion resistance and minor Nd significantly improved the tensile properties of ZK60 alloy. Arrabal et al [19] found that Nd refined β-phase morphology in AZ91 and improved the corrosion resistance. Chen et al [20] studied the effect of Nd on mechanical properties of rotary forging AZ71. It was concluded that mechanical properties reached the maximum value when added 1.0 wt % Nd to AZ71. In addition, Zengin et al [21] reported that the yield and ultimate tensile strengths of A4 (Mg-4 wt% Al) alloy were remarkably improved by addition of 1 wt% Ce/La. Xu et al [22] studied the addition of mischmetal (Ce, La, Pr, Nd) could enhance the mechanical properties because of the formation of Al-RE phase. Asl et al [23] reported that the presence of thermally stable Al 11 RE 3 increased the mechanical properties of the AZ91 alloy at elevated temperatures. Bayani et al [24] showed that the mechanical properties and hightemperature properties of AZ91 alloy reached the maximum value with the addition of 2.0 wt% misch metal. However, the effect of different La/Nd ratio on microstructure and mechanical properties of Mg-Al alloys is not clear.
Therefore, the objective of the present study is to investigate the effects of La/Nd ratio on the microstructure, ambient and elevated compression properties of as-cast AZ91. The total amount of La and Nd added is kept as constant at 2 wt% and the La/Nd ratio is variable. A generation model of Al 11 RE 3 is established and the growth mechanism of Al 11 RE 3 is analyzed.

Experiment procedure
The AZ91-xLa/Nd alloys were produced from the AZ91D ingot, Mg-12.8 wt% La and Mg-30 wt% Nd master alloys. The experimental alloys were melted and alloyed in a graphite crucible by an electric resistance furnace under a protection of SF 6 (1 vol%) and CO 2 (99 vol%) mixed gas atmosphere. At 750°C, the molten metal was held for 30 min to make RE dissolved completely. The melt was refined by RJ-2 flux and was held at 720°C for 20 min. Then the melt was poured into a steel mold (10 mm in thickness) which was coated and preheated to 200°C . The chemical compositions of AZ-xLa/Nd alloys (listed in Table 1) were tested using inductively coupled plasma method atomic emission spectroscopy (ICP-AES).
The microstructures of as-cast AZ91-xLa/Nd alloys were analyzed by scanning electron microscopy (SEM, FEI QUANTA-200) equipped with dispersive x-ray spectroscopy (EDX) operated at 20 kV. The specimens (10 mm × 10 mm × 20 mm) for SEM were cut from the ingots (80 mm × 100 mm × 10 mm) were polished using grades of polishing paper and then etched by 4% nitric acid in alcohol. The quantitative metallographic analysis (the amount and length) was conducted for each alloy by using SEM images with the Image-pro v5.0 software. The amount of needle-like phase is calculated from the true area (TA) of SEM micrographs (800×1000 μm 2 ). The phase identifications of experimental alloys were carried out by x-ray diffractometer (XRD) using Cu Kα radiation operated at 40 kV and 40 mA (K=1.5418 Å). Samples (Φ40 mm × 6 mm) for hardness testing were cut from the as-cast AZ91D-xLa/Nd alloys. The Brinell hardness of each specimen was measured with a 310HBS-3000 digital display hardness tester (HBS5/250/30). The average hardness of each sample was obtained from five tested values. Three cylindrical compression specimens (Φ10 mm × 25 mm) were machined from each alloy ingot to ensure the reproducibility of the data, and the compression tests were performed at a rate of 1 mm min −1 on a WDW-200 electronic universal testing machine at room temperature (20°C) and high temperature (150°C).

Results
3.1. As-cast microstructure structure Figure 1 shows the SEM images of the as-cast AZ91-xLa/Nd alloys. Figure 1(a) shows a typical microstructure of AZ91 alloy, and the grey zone indicated by α is Mg matrix. The coarse dark grey strip-shaped phase indicated by β is β-Mg 17 Al 12 which is evenly distributed along grain boundaries [25]. At room temperature, β phase can improve the hardness and compression properties of magnesium alloys [26]. However, the application of magnesium alloys has been greatly limited due to their poor high-temperature properties [27].
show the presence of a large number of needle-like phases (indicated by R) and irregular block-like phases (indicated by β) in AZ91-xLa/Nd alloys. The β phase becomes finer dispersed particle with the addition of La and Nd. The length and quantity of needle-like phase are different in AZ91-xLa/Nd alloys with increasing La/ Nd ratio. Additionally, it is important to note that the white particulate phase marked by cycles, which emerges in AZ91 with RE addition.
In order to obtain information on the behavior of RE as well as on the distribution of Al and Zn in AZ91-xLa/Nd alloys, EDS area analysis was carried out and the results are shown in figure 2. The results clearly confirm that Al element concentrates on the needle-like phase and the block-like phase (figure 2(c)). La and Nd elements congregate in the needle-like phase and the white particulate phase (figure 2(d) and (e)). With respect to Zn, the segregation to the bulk phase is less pronounced, but still present (figure 2(f)). Figure 3 shows the microstructures of different second phases in as-cast AZ91-xLa/Nd alloys imaged by backscatter electron (BSE) and corresponding EDS results. Figure 3(a) shows the elongate acicular phase and the irregular polygonal phase distribute along the boundary of the matrix. Uniform distribution of second phases in the cell boundaries may increase the mechanical properties effectively [28]. Figure 3(d) shows the EDS analysis of second phases marked by A, B, and C in figure 3(a). The EDS analysis result of A from figure 3(d) reveals the existence of the eutectic phase Mg 17 Al 12 known as β phase [29]. Figure 3(b) shows the micrograph of β-Mg 17 Al 12 which has two kinds of morphologies, the laminar shaped one surrounding the irregular massive one [30].    solidification, aluminum is enriched towards the periphery of grain due to the slow diffusion. In addition, the solidification process proceeds with the magnesium-rich end of the ternary phase diagram of Mg-Al-Zn alloy, resulting in segregation of Zn and Al in the residual liquid along grain boundary [35]. Figure 4 shows the XRD patterns of the as-cast alloys: Alloy I, Alloy II, Alloy IV, and Alloy VI. As can be seen from the diffraction peaks in figure 4(b), with the increase of the La/Nd ratio, the peak intensity of Al 11 La 3 phase becomes stronger and the peak intensity of β-Mg 17 Al 12 phase becomes weaker. It reveals that the number of Al 11 La 3 phase increases and the number of β-Mg 17 Al 12 phase decreases. Besides, the peak intensity of Al 11 Nd 3 becomes weaker as La/Nd ratio increases from 0/0 to 2/1 and disappears almost when the La/Nd ratio is more than 1/1 (shown in figure 4(c)). It reveals that the number of Al 11 Nd 3 phases in the alloy decreases significantly. Light rare earths (La, Ce, Pr, and Nd) have quite similar properties and can be referred to collectively as RE. But the formation of Al x RE y is sensitive to individual rare-earth elements [36]. Compared with Nd, La combines with Al more readily to form Al 11 La 3 due to higher electronegativity difference with Al [37]. As consequence, the number of Al 11 Nd 3 phases in the AZ91-xLa/Nd alloys less than Al 11 La 3 when the La/Nd ratio is more than 1/1. Figure 5 shows the area fraction of β phase in the AZ91D-xLa/Nd alloys. In Alloy I (without La and Nd added), the area fraction of coarse strip-shaped β-Mg 17 Al 12 phase is 27.71%. In Alloy II, III, IV, V and VI (with La and Nd added), the area fraction of β-Mg 17 Al 12 phase is 8.68%, 7.58%, 3.81%, 3.51% and 4.86%, respectively. When La/Nd ratio is 1/1 (Alloy V), the area fraction of β-Mg 17 Al 12 phase is the smallest, only 3.51%. It is indicated that the additions of La and Nd to AZ91D-xLa/Nd alloys results in the great decrease of volume and size of β-Mg 17 Al 12 phases [14]. As the La/Nd ratio increases, the area fraction of β-Mg 17 Al 12 phase decreased  significantly and then increased slightly. The morphology of β-Mg 17 Al 12 phase becomes coarsen when the La/ Nd ratio is greater than 1/1. Figure 6 shows a statistical histogram of the quantity and length of the needle-like phase calculated from the TA of SEM micrographs in the AZ91D-xLa/Nd alloys. It is observed that La/Nd ratio has a strong influence on the quantity and length of the needle-like phase in the alloys. With increasing La/Nd ratio, the amount of needle-like phase in the alloy increases and its average length decreases. The amount increases as the La/Nd ratio increases from 0/0 to 1/1 and reaches the maximum when La/Nd ratio is 1/1. The amount decreases when La/ Nd ratio is greater than 1/1. Besides, the longest length of needle-like phase in alloys increases as the La/Nd ratio increases from 0/0 to 2/3 and then decreases when the La/Nd ratio is greater than 2/3. The longest length reaches the maximum (300.18 μm) when the La/Nd ratio is 2/3. The average length of needle-like phase in alloys decreases from 78.13 μm to 68.42 μm as La/Nd ratio increases from 0/0 to 3/2. Figure 7 shows the Brinell hardness of the AZ91D-xLa/Nd alloys. Brinell hardness decreases from 66.86 to 65.03 and then increases to the peak value of 67.10 as the La/Nd ratio increases from 0/0 to 2/3. As the La/Nd ratio increases from 2/3 to 3/2, Brinell hardness decreases again to the minimum value of 64.50. It is revealed that the needle-like phase has no significant improvement on the Brinell hardness at room temperature. This may be due to the great decrease of fine granular β-Mg 17 Al 12 phase. Besides, needle-like phase with big size is hard to have contribution to hardness, as hardness testing depends on the nature of the constrained deformation around the indenter tip. Figure 8 shows the stress-strain curves of the as-cast alloys (a) at 20°C and (b) at 150°C obtained by the compression tests. It can be seen that AZ91-xLa/Nd alloys (with La and Nd added) exhibit superior compression strength and moderate elongation compared with AZ91 both at room temperature and high temperature. Besiders, when the La/Nd is 2/1 (Alloys VI) shows higher ductility than other AZ91-xLa/Nd alloys. Liu [38] reported that the strength and elongation of Mg-4 wt% Al alloys may decrease with the increasing amount of coarse and brittle intermetallics. The high ductility of Alloys VI should be related to the shortest average length of the needle-phase. AZ91-xLa/Nd alloys containing short and fine Al x RE y may increase strength and ductility. Figure 9 shows the SEM micrographs of fractured surface and the macroscopic sample after compression. All the fractures of AZ91-xLa/Nd alloys exhibit the same typical characteristics: fractures occur in a direction of 45°from the axis. AZ91-xLa/Nd alloys show no obvious signs of softening in compression experient at 150°C from the macroscopic images of the fracture. During the compression process, grains that cannot withstand the critical shear stress crack ( figure 9(a)) at the grain boundaries on the sliding surface, and then expand to fracture. The fracture is characterized by cleavage steps (figure 9(a)) and cleavage rivers ( figure 9(b)). Quasi-cleavage type fracture (shown in figures 9(c) and (d)), which blends cleavage facets with uneven surface and micropores caused by dimple rupture, indicates that plasticity of AZ91-xLa/Nd alloys is improved at 150°C. Figure 10 presents the compression results at room temperature (20°C) and elevated temperature (150°C). The compressive property at room temperature has been improved markedly as La/Nd ratio increases from 0/0  to 1/5. In alloy II, a considerable number of β-Mg 17 Al 12 phases and Al 11 RE 3 phases strengthen the AZ91D-xLa/ Nd alloys with a lower La/Nd ratio of 1/5 at room temperature. It is revealed that the interaction of β-Mg 17 Al 12 phases and Al 11 RE 3 phases strengthen the compression properties of the alloys at room temperature. Fine β-Mg 17 Al 12 phases and a right amount of Al 11 RE 3 phase distributed around the grain can strengthen the alloy for their effective impediment to grain boundary sliding. The compressive property (20°C) decreases from 363 MPa to 337 MPa as La/Nd ratio increases from 1/5 to 1/3 due to the great reduction of β-Mg 17 Al 12 phases and only a few increase of Al 11 RE 3 phases. As La/Nd ratio increases from 1/3 to 1/1, the compressive property increases to a peak value of 364 MPa. Meanwhile, the fraction of β-Mg 17 Al 12 phases continuously decreases and the number of Al 11 RE 3 phases increases. It suggested that the strengthening effect of Al 11 RE 3 phases increases as the number increases. The compressive property (20°C) decreases again (in Alloy VI) when La/Nd ratio is greater than 1/1. Because the amount of Al 11 RE 3 phase increases significantly than the other alloys and the morphology of β-Mg 17 Al 12 phases coarsens slightly in Alloy VI. However, massive Al 11 RE 3 phases will cause the matrix to be split significantly and deteriorate the compressive properties. The compressive properties of AZ91 alloy reach the peak value under the condition of La and Nd added with a right La/Nd ratio.

Compressive properties
The compressive property at elevated temperature increases from 274 MPa to 330 MPa as La/Nd increases from 0/0 to 2/3 and decreases from 330 MPa to 318 MPa as La/Nd ratio increases from 2/3 to 2/1. When La/ Nd ratio is 2/3, the compressive property at 150°C reaches an optimal value. The β-Mg 17 Al 12 phases strengthen the Mg-Al alloys effectively at room temperature, but the strengthening effect is weakened at elevated  temperature since the softening and coarsening of the β-Mg 17 Al 12 phases [27]. The needle-like phases (Al 11 RE 3 ) are very effective in strengthening Mg-Al alloys at elevated temperatures for their thermodynamically stableness [16,39]. The needle-like phases disperse with the right number and average length at grain boundaries and in the interiors of grains, which hinder the grain boundary sliding and harden the alloys [15]. Too many needle-like phases will cause the matrix to be split and induce the deterioration of the compressive properties. Therefore, AZ91D-xLa/Nd alloys exhibit excellent mechanical properties only on a condition of balanced mass ratio at La/Nd.

Discussion
The analysis above has been proved that La/Nd ratio has a strong influence on the phase morphology and mechanical properties of the AZ91D-xLa/Nd alloys. The addition of the rare earth elements La and Nd forms the thermally stable phase of Al 11 La 3 and Al 11 Nd 3 due to the difference in electronegativity between the elements. The electronegativity and electronegativity difference of the La, Nd, Mg and Al elements are listed in Table 2 [20,40,41]. Rare earth elements are more easily combined with Al to form thermally stable intermetallic compounds with La and Nd addition since the difference of electronegativity between La, Nd, and Al is greater than the difference of electronegativity with Mg In addition, the Al-rich rare earth compounds process the same crystallization behavior because they are isostructural with other Al-RE compounds during the solidification process [42,43]. At the same time, in the agglomerated region of Al, Mg elements, β-Mg 17 Al 12 is still generated.
The addition of La and Nd causes intensive constitutional supercooling in front of the solid/liquid interface, promoting the primary α-Mg phase solidification and enriching the Al elements in the amorphous liquid [44]. Meanwhile, the concentration of La, Nd and Al increases and restricts the growth of eutectic β-Mg 17 Al 12 phases [39,45,46]. A possible mechanism of the formation of the Al 11 RE 3 phases and β-Mg 17 Al 12 phases could be proposed and described briefly below.
Case I: For AZ91D-xLa/Nd alloys with lower La/Nd ratios (< 2/3), the amorphous region is similar to region I ( figure 11(a)). It would be easy to form a large aggregative zone of Nd and a small cluster of La since the content of Nd in region I is higher than La. Meanwhile, a scarce region of rare earth elements is formed around the large aggregative zone of Nd. Nd combines with Al to form Al 11 Nd 3 and stops the growth when Nd is exhausted without adequate supplement. The small cluster of La would generate Al 11 La 3 and stops the growth when La is exhausted. The needle-like phases grow short in length and less in quantity in this case.
Case II: For AZ91D-xLa/Nd alloys with higher La/Nd ratios (2v33v2), the amorphous region is similar to region II ( figure 11(b)). There are many La, Nd clusters with suitable size in the melt of region II. Al 11 Nd 3 is generated first probably and an appropriate amount of La in the alloy would be replenished to the cluster of Nd when Nd is exhausted. When La is exhausted, it would be supplemented with an appropriate amount of Nd and would be continuing to grow. In this case, needle-like phases grow up sufficiently to considerable long size and much number.
Case III: Case III with great La/Nd (=2/1) is similar to Case I. In this case, as shown in region III ( figure 11(c)), It would form a large aggregative zone of La and a small cluster of Nd. A scarce region of rare earth elements is formed around the large aggregative zone of La. La combines with Al to form Al 11 La 3 and stops the growth when La is exhausted without supplement adequately. The small cluster of La generates Al 11 La 3 and stops growing when La is depleted.
Case IV: The amorphous region of AZ91D-xLa/Nd alloys would include some regions like region IV ( figure 11(d)). La and Nd in these regions accumulate inside the ma/trix and the formation of the needle-shaped phase in combination with the Al element along the boundary would grow from the periphery toward the center. The needle-like phases may grow throughout the entire grain.
Case V: The growth of β-Mg 17 Al 12 phase in the amorphous region is similar to region V ( figure 11(e)). Al combined with Mg around the boundary of the matrix would generate β-Mg 17 Al 12 phase with a limited average size because of the consumption of Al with rare earths.