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Korean Journal of Metals and Materials > Volume 60(9); 2022 > Article
Zhang, Kim, Liao, Kim, Choi, Lee, and Pan: Enhancing mechanical properties of Mg-Gd-Y-Zn alloys via microalloying with Ce and La

Abstract

This study investigated the microstructure and mechanical properties of Mg-1Gd-1Y-1Zn (at.%) alloys containing designed amounts of Ce or La. The Mg5(Gd,Zn) phase formed in the as-cast Mg-Gd-Y-Zn-Ce/La alloys and disappeared after a homogenization treatment at 500°C for 24 h. The addition of Ce and La resulted in the formation of Ce(Mg,Zn)12 and La(Mg,Zn)12 phases, respectively. Except for that, the Ce or La addition had no significant effect on the morphology, volume fraction, and type of the long-period stacking ordered (LPSO) phases in the Mg-Gd-Y-Zn alloy. The grain size decreased with increasing microalloying content because the heavy Ce and La atoms impeded atomic migration across the boundaries. The solute drag effect led to the formation of the rare earth texture in the extruded Mg-Gd-Y-Zn-Ce/La alloys, whose extent decreased with increasing microalloying content. The mechanical strength was improved by the addition of Ce or La at the sacrifice of ductility. In particular, La exhibited a stronger reinforcement ability than Ce when it was added to the Mg-Gd-Y-Zn alloys. Among the investigated chemical compositions, the Mg-1Gd-1Y-1Zn-0.3La alloy exhibited the highest strength because it had the finest grains, the highest volume fraction of the second phase, and the weakest texture intensity. Furthermore, the alloys showed an unusual yield asymmetry due to the difference in the deformation mode of the LPSO phase.

1. INTRODUCTION

Magnesium (Mg) and its alloys have received increasing attention from the aerospace, automotive, and 3C industries, largely because of their light weight, good shock absorption performance, and high specific strength [1-4]. However, the low engineering strength and strong tensile-compressive yield symmetry of extruded Mg alloys hinder their widespread application [5-7]. As one of solutions to these drawbacks, researchers have investigated Mg-Gd/Y-Zn [8,9] and Mg-Gd-Y-Zn alloys [10-12] with long-period stacking ordered (LPSO) strengthening.
Ce is one of the cheapest rare earth (RE) elements and is widely used in Mg alloys to improve their microstructure and mechanical properties [13]. For example, the addition of Ce improved the yield strength of Mg-Zn alloy through the formation of a new Mg-Zn-Ce phase [14]. Similarly, the yield strength of ZK60 alloy increased by ~50 MPa with the addition of 1.5 wt.% Ce [15]. The strengthening arises from the easy precipitation of Ce-rich phases in Mg-Ce alloys during solidification, because of its low maximum solubility. Jeong et al. [16] reported that particles in the second phase changed from Mg-Zn phase to Ce(Mg1-xZnx)11 with the addition of Ce to Mg-Zn-Zr alloys. Besides this strengthening, the ternary compounds promoted dynamic recrystallization by particle-stimulated nucleation (PSN), which led to a reduction in grain sizes and the weakening of the basal fiber texture.
La has similar physical and chemical characteristics to those of Ce, with low solid solubility and the easy formation of stable phases in Mg alloys [17]. However, La in Mg alloys has a stronger strengthening effect compared with Ce [18]. Zengin and Turen [19] reported that the tensile strength of the Mg-6Zn alloy reached 360 MPa with the addition of 1 wt.% of La. Du et al. [20] proposed adding La to weaken the texture intensity, by the existence of La solute atoms in the matrix. Furthermore, the addition of La into Mg-Zn alloys results in the formation of the Mg-Zn-La phase with an orthorhombic structure, which plays an important role in optimizing microstructure and mechanical strength [19].
A limited number of studies have investigated the effect of Ce and La on the Mg-Gd-Y-Zn alloy system, even though such microalloying has been proven to improve the material performance of Mg-Zn, Mg-Zn-Zr, and Mg-Zn-Y alloys, as noted. In particular, Kim and Kawamura [21] confirmed the enhanced mechanical properties of Mg-Zn-Y alloys following microalloying with Ce and La at both room and elevated temperatures. Their work suggests Mg-Gd-Y-Zn alloys could be potentially improved by a similar microalloying with Ce and La. Therefore, the present study investigated the effect of microalloying in Mg-1Gd-1Y-1Zn alloys with different contents of Ce and La.

2. EXPERIMENTAL PROCEDURE

The nominal compositions of the investigated alloys were Mg-1.0Gd-1.0Y-1.0Zn-xCe/La (x = 0, 0.1, 0.2, and 0.3). The actual chemical compositions of these alloys were determined by 800CCDE X-ray fluorescence spectrometry. The results and the corresponding designations of the prepared alloys are listed in Table 1. The alloys were prepared from commercially pure Mg (99.99wt.%), pure Zn (99.99wt.%), Mg-30%Gd (wt.%), Mg-30%Y (wt.%), Mg-30%Ce (wt.%), and Mg-30%La (%) master alloys in an electric resistance furnace with a mild steel crucible under the protection of CO2 and SF 6 mixed gas. The distribution and composition of the alloys were improved by stirring the mixed melt for 10 min using a steel bar and holding it at 720°C for 15 min. The melt was poured into steel molds to produce ingots with diameters of 80 mm and 150 mm in length. The cast ingots were then homogenized at 500°C for 24 h, followed by quenching in water at 25°C. Finally, these ingots were extruded into bars with diameters of 25 mm at 430°C. The extrusion ram speed and extrusion ratio were 1.5 mm·s−1 and 11:1, respectively.
The constituent phases of the as-cast and homogenized alloys were investigated by X-ray diffraction (XRD; Rigaku D/MAX-2500PC) using a copper target in the scanning angle range of 20° to 70° with a scanning speed of 2°·min−1. The microstructures were characterized with a field-emission scanning electron microscope (FE-SEM; JEOL JSM-7800F) equipped with an energy-dispersive X-ray spectroscope (EDS) and an electron backscattered diffraction (EBSD) system (HKL Channel 5) using a fixed scan step size of 0.5 μm. For EBSD analysis, the samples were electropolished with an AC2 electrolyte. Transmission electron microscopy (TEM; FEI) was performed at 200 kV. All the micrographs were taken from a central section parallel to the extrusion direction (ED). The volume fraction of different phases was determined using Image J software.
The design of tensile and compressive specimens was determined based on the ISO 527-1-2012 standard. The tensile tests used a rod-type dog-bone specimen with a gage diameter of 5 mm and a gage length of 25 mm, while the compressive tests used a cylindrical specimen with a diameter of 8 mm and height of 12 mm. These mechanical tests were performed using a CMT5105 material testing machine at a strain rate of 10−3 s−1 at room temperature. The mechanical properties were measured in triplicate to obtain the mean values.

3. RESULTS

3.1. Microstructural characterization of Mg-Gd-Y-Zn-Ce alloys

Figure 1 presents SEM micrographs and XRD line profiles of the as-cast Mg-Gd-Y-Zn-Ce alloys. The Z10 alloy contained only one type of second phase, as marked by the yellow arrows, which is the LPSO phase. In contrast, the addition of Ce gave rise to two types of second phase with an alternating distribution. The dark gray (the yellow arrows) and light gray (the red arrows) phases in the SEM micrographs correspond to the LPSO phase and the intermetallic compounds, respectively [22]. In contrast to the SEM micrographs, the XRD analysis detected only two types of peaks, for α-Mg and the LPSO phase. The absence of XRD peaks for the intermetallic compounds may be due to their low volume fraction; increasing the Ce content from 0.1 to 0.3 at.% only increased the volume fraction of the intermetallic compounds from 1.58 to 4.21%.
Figure 2 shows the microstructural evolution of the Mg-Gd-Y-Zn-Ce alloys, induced by the homogenization treatment at 500°C for 24 h. Homogenization rarely changed the composition of the second phase in the Z10 alloy. The Ce-added alloys formed two types of LPSO morphologies: lamellar LPSO (the blue arrow) and block LPSO (the yellow arrow). Comparing Figure 2 with Figure 1 suggests the lamellar LPSO phases formed late, after the high-temperature heat treatment [23]. The intermetallic compounds in the as-cast Mg-Gd-Y-Zn-Ce alloys disappeared after homogenization, whereas the new Mg-Zn-Ce phase (the green arrow) formed inside the block LPSO phase.
The volume fraction of Mg-Zn-Ce phases increased from 2.12 to 7.28% as Ce content increased from 0.1 to 0.3 at.%. The XRD analysis of the ZCe1 and ZCe2 alloys only revealed peaks for the α-Mg and LPSO phases. The absence of Mg-Zn-Ce peaks was attributed to the low volume fraction, similar to the aforementioned case of the intermetallic compounds. However, the ZCe3 alloy exhibited Mg-Zn-Ce peaks due to the increased volume fraction of 7.28%.
Figure 3 shows the microstructural evolution of the Mg-Gd-Y-Zn-Ce alloys subjected to extrusion. SEM and EBSD analyses were intended to investigate the second phase and grain structure, respectively. For this purpose, SEM micrographs were taken on a section with an average number of second phases, whereas the inverse pole figure (IPF) maps tried to contain the least number of second phases. The composition and the volume fraction of the constituent phases rarely changed after extrusion. The Mg-Zn-Ce phases still remained in the LPSO phase after extrusion. However, the morphology of the block LPSO phase changed from a net structure to an elongated structure aligned with ED, and the lamellar LPSO phase nearly disappeared after the process.
The Z10 alloy exhibited an incompletely recrystallized structure. In contrast, the Ce-added alloys presented a microstructure with nearly complete recrystallization. The distribution of the grain size is balanced, indicating a high degree of dynamic recrystallization (DRX) in the Mg-Gd-Y-Zn-Ce alloys. The average grain size (AGS) was determined to be 3.08 μm for Z10, 8.59 μm for ZCe1, 6.85 μm for ZCe2, and 4.36 μm for ZCe3 alloys. The AGS values of the Ce-added alloys were higher than that of the Ce-free Z10 alloy. However, compared with the Ce-added alloys, the value decreased with increasing Ce content, suggesting that the Ce addition induced an effective grain refinement in the Mg-Gd-Y-Zn-Ce alloying system. This is further discussed in Section 4.
Figure 4 presents the (0001) pole figures and corresponding inverse pole figures of the extruded Mg-Gd-Y-Zn-Ce alloys. The Z10 alloy exhibited the typical <01−10>// ED fiber texture, and the c-axis of the majority of grains was perpendicular to the TD. The texture of the Ce-added alloys was identified as the weak RE texture, indicating that the c-axis of certain grains lost their perpendicularity with respect to ED. The maximum pole intensity of the Ce-added alloys decreased with increasing Ce content, resulting in the lowest value of 5.59 for the ZCe3 alloy. Indeed, the pole figure of the ZCe3 alloy was distinguished from those of the other alloys. It shows two preferred orientations; one is almost perpendicular to the ED and another is almost parallel to the ED.
Figure 5 shows the TEM analysis of the extruded ZCe3 alloy possessing the LPSO and Mg-Zn-Ce phases. The trails of the lamellar and block LPSO phases were confirmed in the bright-field TEM micrograph. Selected area electron diffraction (SAED) analysis revealed the 14H-type LPSO phase in the ZCe3 alloy. This is consistent with the literature [24,25] which concluded that the 14H-type LPSO phase is easily formed with prolonged high-temperature heat treatment. The micrograph also confirmed that a kink deformation (the blue circle) formed under high stress during the extrusion [26].
Meanwhile, it was difficult to distinguish the Mg-Zn-Ce phase from the block LPSO in the bright-field micrograph, because of their similar morphology. The authors employed EDS surface scanning maps to resolve this issue. The main constituents of the LPSO phase were Mg, Zn, Gd, and Y, while those of the Mg-Zn-Ce phase were Mg, Zn, and Ce. The EDS analysis confirmed the Mg-Zn-Ce phase based on this deduction, whose chemical composition was determined to be Mg-3.50%Zn-5.39%Ce (at.%). The phase structure was identified to have a special orthorhombic structure [27].

3.2. Microstructural characterization of the Mg-Gd-Y-Zn-La alloys

Figure 6 presents the SEM micrographs and XRD line profiles of the as-cast Mg-Gd-Y-Zn-La alloys. The SEM images characterized four types of constituent phases: α-Mg, LPSO (the yellow arrow), intermetallic compound (the red arrow), and Mg-Zn-La phase (the green arrow). The morphologies of the LPSO and Mg-Zn-La phases looked almost the same, with a characteristic dark gray block. In contrast, the intermetallic compound was confirmed to have a light gray bone-like structure. The volume fraction of the intermetallic compound was low (0.89%), and barely changed with the La content. However, the volume fraction of the Mg-Zn-La phase rapidly increased with increasing La content; the fraction was only 0.67% in ZLa1 in contrast to 10.23% in ZLa3. This low amount of Mg-Zn-La phase in the ZLa1 alloy resulted in the undetected peaks in its XRD line profile. This is also applicable to the absence of XRD peaks in the intermetallic compounds in the investigated alloys.
Figure 7 shows the microstructural evolution of the Mg-Gd-Y-Zn-La alloys after a homogenization treatment at 500°C for 24 h. The intermetallic compounds were dissolved after the heat treatment, as shown in the case of the Ce-added alloys. The volume fraction of the Mg-Zn-La phase (the green arrow) moderately increased, while that of the LPSO phase (the yellow arrow) barely changed. This is also supported by the increased intensity of Mg-Zn-La peaks in the XRD line profile. Consequently, the intermetallic compounds were transformed into the Mg-Zn-La phases during the high-temperature homogenization.
Figure 8 shows SEM and EBSD micrographs of the extruded Mg-Gd-Y-Zn-La alloys. Figure 9 shows the (0001) pole figures and corresponding inverse pole figures. The extrusion process hardly changed the morphology or the volume fraction of the second phases. The AGS values were measured to be 7.39 μm for ZLa1, 5.24 μm for ZLa2, and 3.86 μm for ZLa3 alloys, respectively. The results indicate a negative correlation between the AGS and La content. It should be note that the La-added alloys possessed a unique texture, distinguished from those of the Ce-added alloys. In particular, ZLa1 alloy exhibited the typical <0001>//ED basal. The texture intensity decreased from 13.26 to 6.86 as the La content increased from 0.11 to 0.30 at.%, suggesting that the increasing La content enfeebled the texture after extrusion.
Figure 10 provides the detailed TEM analysis of the extruded ZLa3 alloy. The block LPSO phase appears as a rectangular strip with a length of 0.50-0.82 μm and a width of 0.20-0.29 μm. The Mg-Zn-La phases exhibited a morphology of large irregular blocks with a diameter higher than 3 μm. The TEM-EDS analysis confirmed Mg, Zn, Y, and Gd were the major constituents of the LPSO phase, and found Mg, Zn, and La formed the Mg-Zn-La phase. The EDS line profile of the latter confirmed the consistent result of Mg-1.88%Zn-6.37%La (at.%). The high-resolution TEM micrograph and the corresponding SAED pattern confirmed the type of the LPSO phase to be 14H. The thickness of a single stacking structure was approximately 3.62 nm, based on this micrograph. The basal plane of the LPSO phase was parallel to [10−10]α and the stacking direction along [0001]α was consistent with reported information of the LPSO phase [28]. The SAED pattern of the Mg-Zn-La phases confirmed an orthorhombic crystal structure.

3.3. Mechanical properties of the extruded alloys

Figure 11 shows the room-temperature mechanical properties of the extruded alloys. Z10 alloy exhibited a tensile yield strength (TYS) of 262 MPa, compressive yield strength (CYS) of 302 MPa, ultimate tensile strength (UTS) of 360 MPa, and elongation to failure of 6.5%. These numbers are in good accord with the reported data [22]. Both types of microalloying were effective for mechanical strengthening, at the sacrifice of ductility. It should be noted that the La addition resulted in the better strengthening than the Ce addition. The TYS range of the La-added alloys (286-326 MPa) was higher than that of Ce-added alloys (268-286 MPa), as was CYS range (343-387 MPa and 324-353 MPa, respectively). Conversely, the elongation deteriorated more in the former than in the latter. The ZLa3 alloy exhibited the best mechanical properties, with a TYS of 326 MPa, CYS of 387 MPa, UTS of 391 MPa, and ductility of 4.6%. The low ductility was inevitable, since the hard Mg-Zn-Ce/La particles readily supported premature crack propagation during the deformation process [19].
Meanwhile, the investigated alloys exhibited a higher CYS than TYS, leading to a yield asymmetry higher than 1. This trend is in contrast to other Mg alloys (e.g., Mg-Sn-Mn, Mg-Al-Mn, and Mg-Al-Zn), whose CYS were lower than TYS [29,30]. The ZLa3 alloy possessed a yield asymmetry of 1.19; this value was slightly higher than Z10 (1.15), but the second lowest among the Md-Gd-Y-Zn-Ce/La alloys. Therefore, this alloy alleviated the yield asymmetry, in addition to having the best mechanical properties among the investigated candidates.

4. DISCUSSIONS

EDS analysis was conducted to characterize the intermetallic compounds in each alloy, and the results are presented in Table 2. Mg and Gd were the two major elements in the intermetallic compounds, with a Mg/Gd ratio of ~5. Li et al. [31] identified the Mg 5 Gd phase in as-cast Mg-11Gd-3Zn-0.6Zr (wt.%) alloy, which disappeared after a homogenization treatment at 510°C for 20 h. The low volume fraction of this phase gave rise to the absence of the corresponding XRD peaks, as aforementioned. Nevertheless, the intermetallic compounds in the as-cast alloys were confirmed to be Mg 5 (Gd,Zn) through EDS analysis and the literature [32].
Figure 12 compared the XRD line profile of the REMg12 phase with those of the ZCe3 and ZLa3 alloys. The diffraction peaks of the Mg-Zn-Ce and Mg-Zn-La phases appeared to be almost the same as those of the CeMg12 phase (the black columns), although they were slightly shifted towards higher 2θ values. Pavlyuk et al. [33] reported the formation of the Ce(Mg,Zn)12 phase in the Ce-Mg-Zn alloys with lattice parameters of a = b = 1.012-1.027 nm and c = 0.572-0.588 nm. The cell parameters of the CeMg12 phase were a = b = 1.030 nm and c = 0.595 nm [34]. The differences in the cell parameters of the Ce(Mg,Zn)12 and CeMg12 resulted from dissolved Zn atoms, because the atomic radius of Zn (134 pm) was smaller than that of Mg (160 pm). In other words, the lattice distortion caused by Zn resulted in the lower cell parameter of the Ce(Mg,Zn)12 phase as compared with the CeMg12 phase. This is also supported by the XRD data; the dissolved Zn atoms led to a shift in the diffraction peaks. Meanwhile, the TEM-EDS analysis of the Mg-Zn-La phase suggested that the La(Mg,Zn)12 phase had a consistent morphology and XRD pattern [35]. Therefore, the Mg-Zn-Ce and Mg-Zn-La phases were regarded as Ce(Mg,Zn)12 and La(Mg,Zn)12, respectively.
Although the solubilities of Ce and La in the Mg matrix are very low at room temperature, they increased considerably at 500°C. The thermodynamic calculation was performed for Mg-1Gd-1Y-1Zn-xCe/La using Pandat 2013 software with the PanMg 2013 database. The calculation yielded a solubility of Ce and La in the Mg matrix of 0.033 and 0.013 at.%, respectively, at 500°C. The lower solubility of La compared with Ce resulted in the higher volume fraction of the La(Mg,Zn)12 phase as compared with the Ce(Mg,Zn)12 phase after homogenization.
Figure 13 demonstrates the DRX behavior of the extruded Mg-Gd-Y-Zn-Ce/La alloys. The red, blue, and yellow areas represent the deformed, recrystallized, and substructured grains, respectively. Compared with Z10 alloys exhibiting incomplete recrystallization, the Ce/La-added alloys almost completed DRX, with a fraction of DRXed grains higher than 87%. According to Wang et al. [36], the block-shaped LPSO phase effectively promotes DRX, owing to the PSN effect. The Ce(Mg,Zn)12 and La(Mg,Zn)12 phases with a diameter over 1 μm could also promote DRX through a similar mechanism. The La(Mg,Zn)12 phase in the Mg-Gd-Y-Zn-La alloys accounted for a higher volume fraction as compared with the Ce(Mg,Zn)12 phase in the Mg-Gd-Y-ZnCe alloys, thereby explaining the higher frequency of DRXed grains in the former.
Both the Ce- and La-added alloys possessed coarser grains compared to the Z10 alloy, as mentioned above. The Ce(Mg,Zn)12 and La(Mg,Zn)12 phases in these alloys acted as a nucleation core for heterogeneous nucleation, promoting the nucleation and growth of grains through the PSN mechanism. However, the nucleated grains can hinder the growth of adjacent grains by restricting the migration of grain boundaries. This results in the higher degree of DRX with coarser grains in the microalloyed samples, as compared with the Z10 alloy.
In contrast to the typical basal texture of Mg alloys [37], the basal planes of certain RE-containing Mg alloys are not parallel to the ED but have a random orientation [38]. Stanford and Barnett [39] referred to this abnormal texture in Mg-RE alloys as “RE texture”. The size misfit between RE (e.g., Y and Gd) and Mg atoms induces a solute drag effect, impeding atomic migration across the boundaries, which is believed to cause the RE texture [40-42]. This hypothesis supports the positive correlation between the amount of RE content and the intensity of the RE texture; this is consistent with the present results, which found the basal texture intensity decreased with increasing contents of Ce or La, as shown in Figures 4 and 9. The high fraction of DRXed grains may also contribute to the texture weakening
The 18R-type LPSO phase is typically formed in as-cast Mg-Zn-RE alloys, which is then transformed into 14H-type after high-temperature heat treatment for an adequate time [11,43]. The SAED patterns of the LPSO phases in the ZCe3 and ZLa3 alloys indicated the formation of the 14H-type LPSO phase, consistent with the literature [22]. It is thus clear that the addition of Ce or La has a negligible effect on the morphology and the type of the LPSO phase.
The microstructure is usually a key factor, controlling the mechanical properties of an alloy. The Hall-Petch equation explains the negative correlation between grain size and yield strength. This equation explains the strengthening tendency among the Ce- or La-added alloys, where the increasing microalloying contents reduced the AGS value and resultantly enhanced its mechanical strength. However, the Hall-Petch equation fails to explain the low yield strength of the Z10 alloy composed of fine grains. This implies there is a second factor affecting the mechanical properties of the investigated Mg-Gd-Y-Zn-Ce/La alloys, other than the grain size.
The second-phase reinforcement played an important role in the enhanced mechanical properties of the investigated alloys. For a quantitative analysis, Table 3 lists the volume fraction of each phase measured from at least three SEM micrographs. The addition of Ce or La had a significant effect on the volume fraction of the Ce(Mg,Zn)12 and La(Mg,Zn)12 phases, while it barely affected that of the LPSO phase. Furthermore, the composite structure of LPSO phase and intermetallic compound exhibited a higher strength than the LPSO phase alone, which explains the higher yield strength of ZCe1 and ZLa1 as compared with Z10, despite the finer grains in the latter.
According to the Orowan equation [44], increasing the volume fraction of the second phase increases yield strength. Recalling the thermodynamic calculation, the solubility of La in the Mg matrix at 500°C (0.013 at.%) was lower than that of Ce (0.033 at.%), resulting in the higher volume fraction of the La(Mg,Zn)12 than that of the Ce(Mg,Zn)12 phase in the respective alloys. Hence, the La addition leads to a better strengthening effect than the Ce addition to the Mg-1Gd-1Y-1Zn alloy. This also explains the highest mechanical strength of the ZLa3 alloy, which has the highest volume fraction of the second phase, combined with fine grains.
TYS is higher than CYS in conventional Mg alloys [29,30]. The compressive and extension twins often participate in plastic deformation, coordinating deformation during the tensile and compressive tests, respectively [29]. The critical shear stress of compressive twinning is higher than that of extension twinning, leading to the higher TYS [45]. However, this deduction only applies to Mg alloys governed by the basal texture. For the RE-textured alloys, the deformation modes of the tensile and compressive tests differed only slightly, due to the extraordinary grain orientation. Therefore, the difference between TYS and CYS was reduced in this case. It is also noted in LPSO-containing extruded alloys that the LPSO phase easily underwent slip deformation during tensile tests, while kink deformation usually occurred during compressive tests [46]a. As a result, the effective kind-band strengthening contributed to the higher CYS versus TYS in the investigated Mg-Gd-Y-Zn-Ce/La alloys.

5. CONCLUSIONS

The present study investigated the microstructural evolution and mechanical properties of Mg-1Gd-1Y-1Zn alloys microalloyed with Ce or La after casting, homogenization, and extrusion. The following conclusions were drawn from the experimental results.
(1) The Mg 5 (Gd,Zn) phase formed in the as-cast Mg-Gd-Y-Zn-Ce/La alloys and disappeared after homogenization at 500°C for 24 h. La(Mg,Zn)12 phases were found in the cast, homogenized, and extruded Mg-Gd-Y-Zn-La alloys, while the Ce(Mg,Zn)12 phases were found in the homogenized and extruded Mg-Gd-Y-Zn-Ce alloys.
(2) The composite structure of the LPSO and Ce(Mg,Zn)12/La(Mg,Zn)12 phases provided more powerful reinforcement than the LPSO phase alone. Although Z10 alloy possessed the finest grains, the microalloying of Ce or La led to the effective strengthening, due to the formation of the Ce(Mg,Zn)12 or La(Mg,Zn)12 phases. Increasing the microalloying contents led to higher strength due to the increased volume fraction of the second phase and the decreased size of DRXed grains.
(3) The ZLa3 alloy exhibited the optimal mechanical properties at room temperature. This is attributed to its having the weakest texture intensity, the highest volume fraction of the second phase, and the finest grains among the Ce- or La-added alloys.

Fig. 1.
Microstructural characterization of the as-cast Mg-Gd-Y-Zn-Ce alloys: SEM micrograph of (a) Z10, (b) ZCe1, (c) ZCe2, (d) ZCe3, and (e) corresponding XRD line profiles. Figure 1a: image courtesy of.
kjmm-2022-60-9-654f1.jpg
Fig. 2.
Microstructural characterization of the homogenized Mg-Gd-Y-Zn-Ce alloys: SEM micrograph of (a) Z10, (b) ZCe1, (c) ZCe2, (d) ZCe3, and (e) corresponding XRD line profiles. The insets in Figures 2b-d are magnified images showing the Mg-Zn-Ce phase.
kjmm-2022-60-9-654f2.jpg
Fig. 3.
SEM micrographs and EBSD IPF maps of the extruded Mg-Gd-Y-Zn-Ce alloys: (a) Z10, (b) ZCe1, (c) ZCe2, and (d) ZCe3. Figure 3a: image courtesy of [22].
kjmm-2022-60-9-654f3.jpg
Fig. 4.
(0001) Pole figures and corresponding inverse pole figures of the extruded Mg-Gd-Y-Zn-Ce alloys: (a) Z10, (b) ZCe1, (c) ZCe2, and (d) ZCe3.
kjmm-2022-60-9-654f4.jpg
Fig. 5.
Bright-field TEM micrograph and the corresponding EDS elemental mapping of the extruded ZCe3 alloy. The SAED patterns of the LPSO and Mg-Zn-Ce phases are presented as well.
kjmm-2022-60-9-654f5.jpg
Fig. 6.
Microstructural characterization of the as-cast Mg-Gd-Y-Zn-La alloys: SEM micrograph of (a) ZLa1, (b) ZLa2, (c) ZLa3, and (d) corresponding XRD line profiles.
kjmm-2022-60-9-654f6.jpg
Fig. 7.
Microstructural characterization of the homogenized Mg-Gd-Y-Zn-La alloys: SEM micrograph of (a) ZLa1, (b) ZLa2, (c) ZLa3, and (d) corresponding XRD line profiles.
kjmm-2022-60-9-654f7.jpg
Fig. 8.
SEM micrographs and EBSD IPF maps of the extruded Mg-Gd-Y-Zn-La alloys: (a) ZLa1, (b) ZLa2 and (c) ZLa3.
kjmm-2022-60-9-654f8.jpg
Fig. 9.
(0001) Pole figures and corresponding inverse pole figures of the extruded Mg-Gd-Y-Zn-La alloys: (a) Z10, (b) ZLa1, (b) ZLa2 and (c) ZLa3. Figure 9a: image courtesy of [22].
kjmm-2022-60-9-654f9.jpg
Fig. 10.
Bright-field TEM micrograph and the corresponding EDS elemental mapping of the extruded ZLa3 alloy. SAED patterns of the LPSO and Mg-Zn-Ce phases were presented as well.
kjmm-2022-60-9-654f10.jpg
Fig. 11.
Engineering stress-strain curves of the as-extruded alloys: (a) tensile flow curves of Mg-Gd-Y-Zn-Ce alloys, (b) compressive flow curves of Mg-Gd-Y-Zn-Ce alloys, (c) tensile flow curves of Mg-Gd-Y-Zn-La alloys, and (d) compressive flow of Mg-Gd-Y-Zn-La alloys
kjmm-2022-60-9-654f11.jpg
Fig. 12.
XRD line profile of the REMg12 phase compared with those of ZCe3 and ZLa3 alloys.
kjmm-2022-60-9-654f12.jpg
Fig. 13.
DRX grain distribution of the extruded alloys: (a) ZCe1, (b) ZCe2, (c) ZCe3, (d) ZLa1, (e) ZLa2, and (f) ZLa3.
kjmm-2022-60-9-654f13.jpg
Table 1.
Chemical compositions of the investigated alloys (at.%).
Alloys Mg Gd Y Zn Ce La
Z10 Bal. 1.00 1.02 0.97 - -
ZCe1 Bal. 0.93 0.94 0.98 0.09 -
ZCe2 Bal. 0.91 0.92 0.98 0.17 -
ZCe3 Bal. 0.92 0.92 0.97 0.26 -
ZLa1 Bal. 0.93 0.94 0.96 - 0.11
ZLa2 Bal. 0.98 0.96 1.01 - 0.19
ZLa3 Bal. 0.92 0.93 0.98 - 0.30
Table 2.
Chemical compositions of the investigated alloys based on the EDS analysis (at.%).
Alloys Mg Gd Y Zn Ce La
ZCe1 74.9 14.4 5.0 4.0 1.7 -
ZCe2 74.3 14.1 4.5 4.2 2.9 -
ZCe3 74.1 13.2 5.3 4.8 2.5 -
ZLa1 74.8 14.6 5.0 5.1 - 0.4
ZLa2 75.2 15.4 5.3 3.2 - 0.9
ZLa3 73.9 14.9 5.3 4.8 - 1.1
Table 3.
Volume fractions of different phases and DRXed grains in the extruded alloys.
Alloys DRXed Grains α-Mg (%) LPSO (%) Ce(Mg,Zn)12 (%) La(Mg,Zn)12 (%)
Z10 51.0 76.70 23.30 - -
ZCe1 87.5 75.86 22.56 2.12 -
ZCe2 87.3 73.17 23.79 4.68 -
ZCe3 90.5 71.81 23.98 7.28 -
ZLa1 93.4 73.73 23.78 - 3.42
ZLa2 94.5 70.64 22.49 - 7.56
ZLa3 92.8 65.81 23.16 - 12.82

REFERENCES

1. S. You, Y. Huang, K. U. Kainer, and N. Hort, Journal of Magnesium and Alloys. 5, 239–253, (2017).
crossref
2. K. K. Alaneme and E. A. Okotete, Journal of Magnesium and Alloys. 5, 460–475, (2017).
crossref
3. Y. Wang, F. Li, Q. Wang, Q. Chen, X. W. Li, and W. B. Fang, Metals and Materials International. 28, 823–832, (2022).
crossref pdf
4. M. Ra’ayatpour, M. Emamy, and J. Rassizadehghani, Metals and Materials International. 28, 679–694, (2022).

5. J. H. Lee, B. J. Kwak, T. Kong, S. H. Park, and T. Lee, Journal of Magnesium and Alloys. 7, 381–387, (2019).
crossref
6. J. A. Liu, J. Wang, M. L. Yao, and Y. M. Yao, Metals and Materials International. 27, 1613–1619, (2021).
crossref pdf
7. Y.-K. Kim, M.-J. Kim, Y.-J. Hwang, S. K. Kim, H.-K. Lim, and K.-A. Lee, Korean Journal of Metals and Materials. 59, 365–373, (2021).
crossref pdf
8. J.-K. Kim, S. Sandlöbes, and D. Raabe, Acta Materialia. 82, 414–423, (2015).
crossref
9. M. Yamasaki, M. Sasaki, M. Nishijima, K. Hiraga, and Y. Kawamura, Acta Materialia. 55, 6798–6805, (2007).
crossref
10. S. Huang, J. Wang, F. Hou, X. Huang, and F. Pan, Materials Science and Engineering: A. 612, 363–370, (2014).
crossref
11. D. Zhang, Z. Tan, Q. Huo, Z. Xiao, Z. Fang, and X. Yang, Materials Science and Engineering: A. 715, 389–403, (2018).
crossref
12. J. Kim, H. Liao, X. Ou, Z. Zhang, K. Kang, T. Lee and F. Pan, Metals and Materials International 2022).

13. Y. Chai, C. He, B. Jiang, J. Fu, Z. Jiang, Q. Yang, H. Sheng, G. Huang, D. Zhang, and F. Pan, Journal of Materials Science & Technology. 37, 26–37, (2020).
crossref
14. L. Liu, X. Chen, F. Pan, A. Tang, X. Wang, J. Liu, and S. Gao, Materials Science and Engineering: A. 669, 259–268, (2016).
crossref
15. H. Yu, Y. M. Kim, B. S. You, H. S. Yu, and S. H. Park, Materials Science and Engineering: A. 559, 798–807, (2013).
crossref
16. H. Y. Jeong, B. Kim, S. G. Kim, H. J. Kim, and S. S. Park, Materials Science and Engineering: A. 612, 217–222, (2014).
crossref
17. S. Golmakaniyoon and R. Mahmudi, Materials Science and Engineering: A. 620, 301–308, (2015).
crossref
18. Y. Z. Du, X. G. Qiao, M. Y. Zheng, K. Wu, and S. W. Xu, Materials & Design. 85, 549–557, (2015).
crossref
19. H. Zengin and Y. Turen, Materials Chemistry and Physics. 214, 421–430, (2018).
crossref
20. Y. Du, M. Zheng, X. Qiao, W. Peng, and B. Jiang, Materials Science and Engineering: A. 673, 47–54, (2016).
crossref
21. J. Kim and Y. Kawamura, Materials Science and Engineering: A. 573, 62–66, (2013).
crossref
22. H. Liao, J. Kim, J. Lv, B. Jiang, X. Chen, and F. Pan, Journal of Alloys and Compounds. 831, 154873–154873, (2020).
crossref
23. X. Wu, F. Pan, R. Cheng, and S. Luo, Materials Science and Engineering: A. 726, 64–68, (2018).
crossref
24. W. Liu, J. Zhang, C. Xu, X. Zong, J. Hao, Y. Li, and Z. Zhang, Materials & Design. 110, 1–9, (2016).
crossref
25. B.-J. Lv, J. Peng, L.-L. Zhu, Y.-J. Wang, and A.-T. Tang, Materials Science and Engineering: A. 599, 150–159, (2014).
crossref
26. K. Hagihara, M. Yamasaki, Y. Kawamura, and T. Nakano, Materials Science and Engineering: A. 763, 138163–138163, (2019).
crossref
27. W. P. Yang, X. F. Guo, and Z. X. Lu, Journal of Alloys and Compounds. 521, 1–3, (2012).
crossref
28. D. Egusa and E. Abe, Acta Materialia. 60, 166–178, (2012).
crossref
29. H. Liao, J. Kim, T. Liu, A. Tang, J. She, P. Peng, and F. Pan, Materials Science and Engineering: A. 754, 778–785, (2019).
crossref
30. P. Peng, X. He, J. She, A. Tang, M. Rashad, S. Zhou, G. Zhang, X. Mi, and F. Pan, Materials Science and Engineering: A. 766, 138332–138332, (2019).
crossref
31. K. Li, Z. Chen, T. Chen, J. Shao, R. Wang, and C. Liu, Journal of Alloys and Compounds. 792, 894–906, (2019).
crossref
32. S. Yin, Z. Zhang, X. Liu, Q. Le, Q. Lan, L. Bao, and J. Cui, Materials Science and Engineering: A. 695, 135–143, (2017).
crossref
33. V. Pavlyuk, B. Marciniak, and E. Różycka-Sokołowska, Intermetallics. 20, 8–15, (2012).
crossref
34. H. Shi, Q. Li, J. Zhang, Q. Luo, and K.-C. Chou, Calphad. 68, 101742–101742, (2020).
crossref
35. M. L. Huang, H. X. Li, H. Ding, J. W. Zhao, and S. M. Hao, Journal of Alloys and Compounds. 612, 479–485, (2014).
crossref
36. Y. Wang, F. Zhang, Y. Wang, Y. Duan, K. Wang, W. Zhang, and J. Hu, Materials Science and Engineering: A. 745, 149–158, (2019).
crossref
37. B. J. Kwak, S. H. Park, Y. H. Moon, J. H. Lee, and T. Lee, Materials Science and Engineering: A. 788, 139496–139496, (2020).
crossref
38. P. Minárik, D. Drozdenko, M. Zemková, J. Veselý, J. Čapek, J. Bohlen, and P. Dobroň, Materials Science and Engineering: A. 759, 455–464, (2019).
crossref
39. N. Barnett and M. R. Barnett, Materials Science and Engineering: A. 496, 399–408, (2008).
crossref
40. J. D. Robson, Metallurgical and Materials Transactions A. 45, 3205–3212, (2014).
crossref pdf
41. N. Stanford, G. Sha, J. H. Xia, S. P. Ringer, and M. R. Barnett, Scripta Materialia. 65, 919–921, (2011).
crossref
42. C. D. Barrett, A. Imandoust, and H. El Kadiri, Scripta Materialia. 146, 46–50, (2018).
crossref
43. Y. M. Zhu, M. Weyland, A. J. Morton, K. Oh-ishi, K. Hono, and J. F. Nie, Scripta Materialia. 60, 980–983, (2009).
crossref
44. S. Choi, G. Kim, J. P. Kim, S. H. Kim, S. B. Son, and S.-J. Lee, Korean Journal of Metals and Materials. 59, 515–523, (2021).
crossref pdf
45. C. M. Cepeda-Jiménez and M. T. Pérez-Prado, Acta Materialia. 108, 304–316, (2016).
crossref
46. H. Gao, K.-i. Ikeda, T. Morikawa, K. Higashida, and H. Nakashima, Materials Letters. 146, 30–33, (2015).
crossref
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