Effect of La/Zn on Microstructural Evolution and Mechanical Properties of Extruded Mg-9Gd-3Y Alloy

Abstract

This study investigated the impact of La and Zn additions on the microstructure and mechanical properties (at room and high temperatures) of extruded Mg-9Gd-3Y (GW93) alloy. The incorporation of La and Zn induced the precipitation of granular second phases and LPSO phases, increasing the second-phase area fraction in the Mg-9Gd-3Y-0.6La-1Zn (3GW93) alloy to 14.6%. Within the 3GW93 alloy, the Mg₁₂La phase exhibited the following crystallographic orientation relationship with the α-Mg matrix: (020)_Mg12La//(01$\overline{1}$0)_α-Mg and [10$\overline{1}$]_Mg12La//[$\overline{2}$110]_α-Mg. The 3GW93 alloy containing both La and Zn demonstrated the highest strength under both room- and high-temperature conditions. At room temperature, its yield strength (YS) and ultimate tensile strength (UTS) were 284 MPa and 354 MPa, respectively. This represents an increase of 83 MPa in YS and 70 MPa in UTS compared to the GW93 alloy (YS: 201 MPa, UTS: 284 MPa). Similarly, at 300 °C, the 3GW93 alloy (YS: 249 MPa, UTS: 285 MPa) exceeded the GW93 alloy (YS: 182 MPa, UTS: 244 MPa) by 67 MPa in YS and 41 MPa in UTS. The enhanced mechanical properties of the 3GW93 alloy are attributed to synergistic grain refinement and dispersion strengthening effects originating from the LPSO, Mg₅RE, and Mg₁₂La phases.

1. Introduction

Owing to their exceptional strength-to-weight ratio, low density, and superior vibration-damping characteristics, magnesium (Mg) alloys demonstrate extensive potential for utilization across the aerospace, automotive, and electronics industries [1,2,3,4]. However, Mg alloys suffer from significant strength degradation at high temperatures and poor creep resistance, which severely limits their engineering applications as power transmission and drivetrain components in the aerospace, defense, and automotive industries [5,6]. The currently used Mg-Al-Zn (AZ) and Mg-Al-Mn (AM) commercial alloys undergo a sharp decline in mechanical properties when temperatures exceed 120 °C. This occurs because the thermally unstable Mg₁₇Al₁₂ phase readily softens or coarsens above 120 °C, resulting in diminished effectiveness in impeding dislocation motion and, consequently, reducing high-temperature mechanical performance [7,8]. Therefore, developing high-strength magnesium alloys with elevated-temperature resistance is pivotal for advancing magnesium alloy technologies.

The capability to maintain excellent strength and creep resistance at 200–250 °C positions Mg-Gd-Y alloys as prominent candidates in heat-resistant magnesium-rare earth (Mg-RE) alloy research [9,10,11]. The introduction of Zn creates long periods of stacking ordered (LPSO) strengthening phases [12,13]. Compared to α-Mg, LPSO phases possess a superior elastic modulus and hardness, thereby enhancing the alloy’s strength. Lamellar LPSO phases increase the critical resolved shear stress (CRSS) for basal slip while activating non-basal slip systems, consequently enhancing plasticity [14,15,16]. Due to its economic efficiency, La has become a prevalent alloying element in Mg systems. During heat treatment, the low solid solubility of La in Mg (0.03 wt.% at 500 °C [17]) promotes preferential Mg₁₂La phase precipitation along grain boundaries. The Mg₁₂La phase effectively pins grain boundaries, suppressing grain growth [18,19,20]. However, the synergistic effect between the particle-strengthening phases and the LPSO phase on the microstructure and mechanical properties of Mg alloys requires further study. Therefore, in this paper, we aim to investigate how the co-addition of La and Zn influences the microstructure and mechanical properties of the Mg-Gd-Y alloy by promoting the precipitation of the particle and LPSO phases.

This study employed the Mg-9Gd-3Y (GW93) alloy as the base alloy. By adding La and Zn elements, Mg-9Gd-3Y-0.6La (1GW93), Mg-9Gd-3Y-1Zn (2GW93), and Mg-9Gd-3Y-0.6La-1Zn (3GW93) alloys were prepared. The impact mechanisms of La and Zn alloying on the microstructure–property relationships of hot-extruded GW93 alloys were elucidated across room- to high-temperature regimes.

2. Materials and Methods

Pure Mg (99.98 wt.%), Mg-30Gd (wt.%), Mg-30Y (wt.%), and Mg-30La (wt.%) master alloys combined with pure Zn (99.99 wt.%) were melted in a resistance furnace to fabricate four alloys: GW93, 1GW93, 2GW93, and 3GW93. The preparation procedure was as follows: A crucible containing pure Mg was placed into a resistance furnace pre-heated to 500 °C. The furnace temperature was then raised to 720 °C under a protective atmosphere of CO₂/SF₆ (99:1). After the pure Mg melted, Mg-30Gd, Mg-30Y, or Mg-30La (1GW93 and 3GW93) master alloys were added to the melt. After these master alloys dissolved, pure Zn granules were introduced (2GW93 and 3GW93). Following complete melting of all alloying components, the melt was mechanically stirred for 5 min while skimming off surface dross. The molten metal was subsequently held at 720 °C for 10 min before being poured into a pre-heated (200 °C) casting mold (Ø85 mm × 300 mm). Upon solidification, an 85 mm diameter ingot was produced. To reduce errors, each component was repeatedly smelted three times, with one ingot cast each time.

Surface–ground cylindrical billets (Ø80 mm × 60 mm) were prepared from ingots and subsequently underwent homogenization treatment at 520 °C for 12 h. Homogenized billets were formed by forward extrusion using an XJ-500-ton horizontal extrusion press (YCMMC, Wuxi, China). The homogenized billets were pre-heated at 430 °C for 1 h prior to extrusion, which was performed at the same temperature with an 18:1 extrusion ratio and a 1.5 mm/min ram speed. Finally, a bar with a diameter of 20 mm and a length of approximately 1000 mm was obtained.

Microstructural characterization was performed using optical microscopy (OM) with a ZEISS-Axiolab 5 from Carl Zeiss AG (Baden-Württemberg, Germany), scanning electron microscopy (SEM) with a JSM-7800F from JEOL (Tokyo, Japan), and transmission electron microscopy (TEM) with an FEI Talos F200X from Thermo Fisher Scientific (Waltham, MA, USA). Texture analysis employed electron backscatter diffraction (EBSD). The specimens used for OM and SEM observations underwent etching in a 4% nitric acid ethanol solution (~50 s). EBSD data were analyzed using AZtecCrystal software (Version 2.1), with grains exhibiting intragranular misorientation < 2° defined as dynamically recrystallized (DRX) grains, and those exceeding this threshold as unDRXed grains. Tensile tests along the extrusion direction (ED) were conducted at 1 × 10⁻³ s⁻¹ (strain rate) using a universal testing machine (CMT5105-100 kN; SUST, Shenzhen, China) equipped with a GX-1200A temperature control system. Prior to high-temperature testing, the specimens were held at the target temperature for 10 min. The samples used for the tensile test, which were prepared in accordance with the national standards of the People’s Republic of China (GB/T228.1-2021 and GB/T228.2-2015) [21,22], are shown in Figure 1. For statistical reliability, a mechanical extensometer was used and three tests were performed per condition.

3. Results and Discussion

3.1. Microstructural Evolution

OM images of the as-extruded GW93, 1GW93, 2GW93, and 3GW93 alloys are presented in Figure 2. As revealed in Figure 2a, the GW93 alloy primarily consists of DRXed grains generated during extrusion, with sparse second phases aligned along the ED. Figure 2b–d demonstrate that the 1GW93, 2GW93, and 3GW93 alloys all exhibit a bimodal microstructure featuring alternating coarse unDRXed grains and fine DRXed grains, accompanied by regularly aligned second phases parallel to the ED.

Figure 3 displays SEM images of the extruded GW93 series alloys, with

Table 1. SEM-EDS results of points in Figure 3 (at.%).

Point	Mg	Gd	Y	La	Zn
A	3.9 ± 0.3	36.3 ± 1.5	59.8 ± 2.7	-	-
B	86.4 ± 2.9	9.9 ± 1.1	3.7 ± 0.5	-	-
C	93.4 ± 2.3	2.7 ± 0.5	2.8 ± 0.4	1.0 ± 0.2	-
D	85.7 ± 1.6	10.8 ± 0.9	2.8 ± 0.5	0.7 ± 0.2	-
E	88.8 ± 2.7	9.1 ± 0.8	1.8 ± 0.2	-	0.3 ± 0.1
F	85.0 ± 2.3	9.4 ± 1.1	3.2 ± 0.4	0.6 ± 0.1	1.8 ± 0.2

listing the EDS results (at.%) for the marked second phases (the same morphology of the second phase was collected more than 15 times). The white blocky phases in the GW93 alloy are the RE (Y, Gd)-rich phases (Point A), while bright particulate phases correspond to the Mg₅RE phases (Point B). After the sole addition of La, a grayish Mg₁₂La phase (point C) appears in the alloy at the grain boundaries [23], and La elements are also found in the Mg₅RE phase (point D). Zn introduction promotes blocky LPSO (blue markers) and intragranular lamellar LPSO (green markers), while Zn incorporates into the Mg₅RE phase (Point E). Co-addition of La and Zn enables the simultaneous incorporation of both elements into the Mg₅RE phase (Point F). It is worth noting that at points D, E, and F, the combined Gd, Y, La, and Zn content approaches 16.67 at.%—matching the RE stoichiometry of Mg₅RE. This indicates the substitution of La and Zn atoms for some Gd and Y in the Mg₅RE phase [18].

Low-magnification SEM images quantify the area fraction of second phases in the extruded Mg-9Gd-3Y(-La-Zn), with the results presented in Figure 4. When the La element was introduced, it promoted the formation of the Mg₁₂La phase and increased the area fraction of the second phases. In the 2GW93 alloy with added Zn element, the substantial rise in second-phase area fraction (from 0.9% to 9.2%) resulted from the formation of abundant blocky and lamellar LPSO structures in the alloy. The mixed addition of La and Zn mutually promoted the precipitation of the second phase of the particles and LPSO, maximizing the area fraction of the second phase of the 3GW93 alloy (14.6%).

Uniform, fine DRXed grains can be observed in the GW93 alloy (Figure 5a), with no coarse grains present, evidencing full recrystallization during extrusion. Comparative analysis indicated that the grain sizes of the extruded 1GW93, 2GW93, and 3GW93 alloys show a bimodal distribution feature. Figure 5e,i,m reveal the coexistence of fine DRXed grains and coarse unDRXed grains in the microstructure. Concurrently, Figure 5c,g,k,o display average grain sizes of 18.6 μm (GW93), 7.9 μm (1GW93), 6.6 μm (2GW93), and 4.0 μm (3GW93), with corresponding DRXed grain sizes measuring 18.6 μm, 7.8 μm, 6.5 μm, and 3.9 μm. The individual addition of either La or Zn effectively refined the grain size, while the co-addition of La and Zn produced the most pronounced grain refinement effect, with the 3GW93 alloy achieving the minimum grain size of 4.0 μm. Figure 5d,h,l,p statistically quantify the texture intensity of Mg-9Gd-3Y(-La-Zn) alloys. GW93 exhibited a predominantly <0001>∥ED grain alignment with a weak texture intensity (2.07 mrd). After the single addition of the La element, the single addition of the Zn element, and the mixed addition of the La and Zn elements, the alloy still exhibited a fiber texture of <0001>//ED, and the texture intensities were 2.88 mrd, 4.99 mrd, and 4.99 mrd, respectively. Figure 5b,f,g,n display the grain orientation distribution in the extruded Mg-9Gd-3Y(-La-Zn) alloy, where the dark green, light green, and red regions represent DRXed grains, substructures, and deformed grains, respectively. The research results show that the DRX fraction of the GW93 alloy reached 99.4%, while the recrystallization fractions of the 1GW93, 2GW93, and 3GW93 alloys with added La and Zn elements dropped to 94.3%, 89.3%, and 90.9%, respectively. This suggests that the La and Zn additions suppressed the DRX process and impeded DRXed grain growth.

To further investigate the influence of the mixed addition of La and Zn on the microstructure of the GW93 alloy, the 3GW93 alloy was characterized by TEM. The results showed that there were a large number of fine granular second phases at the grain boundaries and within the grains in the 3GW93 alloy. Based on the calibration results of selective electron diffraction (SAED), the Mg₅RE phase (fcc, a = 2.23 nm [24]) has been found both at the grain boundaries and within the grains (Figure 6a,c), and this particle is often formed in Mg-RE alloys [25,26]. In addition, there is the Mg₁₂La phase (Figure 6e,q). The nanoscale Mg₁₂La phase within the grains has a good orientation relationship with α-Mg, specifically, (020)_Mg12La//(01$\overline{1}$0)_α-Mg and [10$\overline{1}$]_Mg12La//[$\overline{2}$110]_α-Mg (Figure 6g,h). Lamellar LPSO phase and blocky LPSO phase were found inside the DRXed grains and at the grain boundaries (Figure 6j,n). The lamellar LPSO phases inside the grains were basically parallel to each other, but the parallel directions of the lamellar LPSO phases inside different grains were different, as shown in Figure 6m. The blocky LPSO phase, Mg₅RE phase, and Mg₁₂La phase distributed between grains could effectively prevent the growth of recrystallized grains. Due to the pinning effect of the second phase at the grain boundaries, the size of grain A is smaller than that of grain B, as shown in Figure 6i. Based on further analysis of LPSO phases with distinct morphological characteristics, the diffraction patterns in Figure 6l,o reveal that the blocky LPSO phase in the 3GW93 alloy adopts an 18R structure, while the lamellar LPSO phase exhibits a 14H structure. This observation is aligned with the analytical results reported in Ref. [27]. In the DRXed grains, dislocations are found to accumulate near the lamellar LPSO phase (Figure 6n), and at the same time, there are also dislocations hindered by the lamellar LPSO phase that accumulate between the LPSO phases (Figure 6p).

Recrystallization is a process in which nucleation and growth occurs through a thermal activation mechanism during the hot working of Mg alloys when the storage energy reaches the critical value needed to release it. This process could effectively alleviate the stress concentration within the microstructure [28,29]. Under the hot extrusion process conditions, the second phase broke and was uniformly distributed within the matrix. This dispersed second phase affected the DRX process through two mechanisms: First, it inhibited dislocation movement and grain boundary migration, thereby hindering the coarsening of recrystallized grains. Second, as it comprised heterogeneous nucleating particles, it significantly increased the nucleation rate and promoted grain refinement [30,31]. In this study, the introduction of the La and Zn elements formed the Mg₁₂La and LPSO phases. Micro-nanoscale particles were formed by the fragmentation of the blocky LPSO phase during the extrusion process. The micrometer-scale LPSO phase hindered dislocation movement and promoted dislocation accumulation and entanglement, and the nanoscale LPSO and Mg₁₂La phases dispersed within the grains promoted DRX nucleation through the particle-stimulated nucleation (PSN [32]) mechanism [33]. On the contrary, the nanoscale Mg₁₂La phase, Mg₅Gd phase, and LPSO phase dispersed at the grain boundaries inhibited the movement of dislocations and the migration of grain boundaries by generating a pinning effect (as shown in Figure 6r, there were still high-density dislocations and deformation regions remaining around the Mg₁₂La phase), thereby suppressing grain growth and the migration of grain boundaries of deformed grains. Therefore, the DRXed grain size of the 3GW93 alloy was smaller than that of the GW93 alloy, and the DRX fraction was lower than that of the GW93 alloy.

The elemental distribution at the 3GW93 alloy grain boundaries is depicted in Figure 7. It was observed that the elements Gd, Y, La, and Zn co-segregated at the grain boundaries, which manifested as the most obvious segregation of Gd and the weakest segregation of La. According to previous studies [34,35,36], RE elements tend to segregate at grain boundaries, owing to their atomic size difference compared to Mg. Due to their combination of high size mismatch and high solid solubility, Y and Gd exhibited the most pronounced segregation effect. The co-segregation of the Gd, Y, La, and Zn elements at grain boundaries enhanced solute drag effects, effectively stabilizing the grain boundaries. This phenomenon contributed to grain refinement and suppressed dislocation motion, ultimately improving material strength. Moreover, the segregation of these elements at grain boundaries significantly increased intergranular fracture energy, thereby enhancing both grain boundary cohesion and material ductility.

3.2. Mechanical Properties

3.2.1. Room Temperature Mechanical Properties

The room temperature tensile stress–strain curves of the as-extruded Mg-9Gd-3Y(-La-Zn) alloys are presented in Figure 8, with the corresponding mechanical properties detailed in

Table 2. Tensile properties of extruded alloys at room temperature.

Alloy	YS(MPa)	UTS(MPa)	ε(%)
GW93	201 ± 4	284 ± 2	8.9 ± 0.2
1GW93	217 ± 3	292 ± 1	8.9 ± 0.4
2GW93	255 ± 3	341 ± 3	13.8 ± 0.5
3GW93	284 ± 2	354 ± 2	7.4 ± 0.3

. La and Zn additions enhanced the mechanical properties of the GW93 alloy. The Zn-modified 2GW93 alloy demonstrated optimal comprehensive mechanical properties, exhibiting a YS of 255 MPa, a UTS of 341 MPa, and a fracture strain (ε) of 13.8%. This represented substantial improvements over the GW93 alloy (YS: 201 MPa, UTS: 284 MPa, ε: 8.9%), with a 54 MPa increase in YS, a 57 MPa increase in UTS, and a 4.9 percentage point elevation in fracture strain. The La and Zn co-modified 3GW93 alloy achieved peak room temperature strength values, with its YS (284 MPa) and UTS (354 MPa) exceeding those of the GW93 alloy by 83 MPa (41.3%) and 70 MPa (24.6%), respectively.

Grain size measurements revealed average diameters of 18.6 μm (GW93), 7.9 μm (1GW93), 6.6 μm (2GW93), and 4.0 μm (3GW93) in the extruded alloys. According to the Hall–Petch relationship [37]:

$${\sigma }_G={\sigma }_0+{kd}^{-1/2}$$

(1)

In the formula, σ_G represents the alloy’s yield strength, σ₀ is the strength constant of the single-crystal metal, k is the Hall–Petch coefficient, and d is the average grain size. Dislocation movement was effectively hindered by the increased grain boundary density, thereby improving alloy strength. With the finest grains among the tested alloys, 3GW93 achieved peak YS at room temperature. Unlike the conventional strengthening mechanism, grain refinement could simultaneously improve the plasticity of the alloy. However, the fracture strain of the 3GW93 alloy (7.4%) was lower than that of the GW93 alloy (8.9%), indicating that in addition to grain size, other factors also influenced the plasticity.

The second-phase volume fraction and particle size distribution collectively modulated DRX evolution and mechanical properties. According to the Orowan mechanism [38]:

$$∆{\sigma }_s\propto f^{1/2}{d}^{-1}\ln d$$

(2)

where Δσ_s represents the yield strength increment, f the second-phase volume fraction, and d the average second-phase size. The 3GW93 alloy exhibited the highest second-phase volume fraction (14.6%). La incorporation simultaneously refined grains and stimulated Mg₁₂La precipitation. The addition of the La element not only refined the grains, but also promoted the precipitation of the Mg₁₂La phase. The introduction of the Zn element induced the formation of the LPSO phase. Blocky LPSO phases hindered dislocation motion while promoting dislocation multiplication and accumulation. Concurrently, lamellar LPSO twisting during deformation enhanced plasticity. Meanwhile, dislocations could accumulate within the lamellar LPSO, thereby enhancing the strength and plasticity of the alloy [39,40]. The combined effects of fine grain strengthening and the dispersion strengthening induced by the LPSO, Mg₅RE, and Mg₁₂La phases improved the mechanical properties of the 3GW93 alloy.

The fracture surfaces of the GW93 and 1GW93 alloys presented classic quasi-cleavage patterns (Figure 9a,b) comprising three distinct features: cleavage planes, tear edges, and dimple structures. It was noteworthy that the cleavage plane sizes and dimple diameters of the 1GW93 alloy decreased, which is attributable to the grain-refining effect of the added La element [18]. The fracture surface of the 2GW93 alloy displays numerous fine dimples, indicating substantial plastic deformation during tensile testing, which corresponds to its high ε (13.8%) (Figure 9c). In contrast, the 3GW93 alloy exhibits fragmentation of second-phase particles on its fracture surface, with clear microcrack propagation traces observed adjacent to these particles (Figure 9d). Previous studies demonstrated that in Mg-Gd-Y alloys, crack nucleation is mainly induced by basal slip. With Zn’s addition, cracks tend to nucleate at the interface between the LPSO phase and α-Mg, then propagate along the blocky LPSO [41,42]. Furthermore, the brittle Mg₁₂La phase is also prone to stress concentration during the deformation process [18]. In the 3GW93 alloy, microcracks preferentially nucleated at the interface between the LPSO phase and α-Mg and the brittle Mg₁₂La phase, then propagated through these regions until macroscopic crack formation and fracture occurred.

3.2.2. High-Temperature Mechanical Properties

Figure 10 displays the tensile stress–strain curves of the extruded Mg-9Gd-3Y(-La-Zn) alloys at high temperatures (200–300 °C), with the corresponding mechanical properties shown in

Table 3. Tensile properties of extruded alloys at high temperatures.

Temperature	Alloy	YS (MPa)	UTS (MPa)	ε (%)
200 °C	GW93	199 ± 3	272 ± 3	16.1 ± 0.4
	1GW93	201 ± 4	285 ± 2	15.9 ± 0.6
	2GW93	248 ± 2	339 ± 5	18.5 ± 0.5
	3GW93	272 ± 3	344 ± 3	11.8 ± 0.6
250 °C	GW93	187 ± 2	259 ± 3	20.4 ± 0.5
	1GW93	199 ± 4	273 ± 5	17.6 ± 0.6
	2GW93	238 ± 3	311 ± 5	18.3 ± 0.7
	3GW93	257 ± 3	317 ± 3	16.1 ± 0.4
300 °C	GW93	182 ± 3	244 ± 3	31.3 ± 0.3
	1GW93	197 ± 2	242 ± 4	35.7 ± 0.5
	2GW93	238 ± 3	275 ± 2	38.6 ± 0.5
	3GW93	249 ± 4	285 ± 3	33.9 ± 0.6

. At high temperatures, the weakening and sliding of grain boundaries reduced their hindrance effect on dislocations, leading to a decrease in alloy strength and an increase in plasticity [43]. At all tested temperatures, the 3GW93 alloy maintained the highest strength. At 200 °C, the YS and UTS of the 3GW93 alloy are 272 MPa and 344 MPa, respectively, which are 73 MPa and 72 MPa higher than those of the GW93 alloy (YS: 199 MPa, UTS: 272 MPa), respectively. At 250 °C, the YS and UTS of the 3GW93 alloy are 257 MPa and 317 MPa, respectively, which are 70 MPa and 58 MPa higher than those of the GW93 alloy (YS: 187 MPa, UTS: 259 MPa), respectively. At 300 °C, the YS and UTS of the 3GW93 alloy are 249 MPa and 285 MPa, respectively, which are 67 MPa and 41 MPa higher than those of the GW93 alloy (YS: 182 MPa, UTS: 244 MPa), respectively. Notably, when the temperature reached 250 °C, the plasticity of the 3GW93 alloy increased significantly. Moreover, when the temperature rose to 300 °C, its plasticity exceeded that of the GW93 alloy.

When deformed at high temperatures, the sliding of the grain boundaries intensified, and the pinning effect of the second-phase particles weakened, resulting in a decrease in the resistance of dislocation movement. This was a key reason for the attenuation of the high-temperature strength of the alloy [44,45]. Meanwhile, non-basal slip was activated in the Mg alloys as the temperature increased, dislocation motion intensified, and the alloy’s plastic deformation capability improved. The GW93 alloy with La and Zn additions maintained relatively high UTS and fracture strain even in a high-temperature environment, demonstrating excellent high-temperature mechanical properties. It can be observed from the TEM analysis (Figure 6) that the second phase in the 3GW93 alloy is uniformly distributed across the recrystallized grains and grain boundaries. During hot extrusion, the uniformly distributed fine Mg₁₂La phase and LPSO phase in the 3GW93 alloy can effectively impede dislocation movement and inhibit grain boundary migration, thereby enhancing the stability of grain boundaries at high temperatures. This played a decisive role in improving the high-temperature strength of the alloy [46]. Based on the second-phase strengthening mechanism, the YS and UTS of the 3GW93 alloy with the combined addition of La and Zn exceeded those of the GW93 alloy within the range of room temperature to 300 °C.

Figure 11 shows the tensile fracture morphology at various high temperatures. The fracture morphology of the GW93 alloy at 200 °C shows a ductile fracture with denser dimples and a small number of tear edges compared to room temperature (Figure 11a). When the temperature reaches 250 °C, the tear edges disappear and the fracture morphology is dominated by dimples (Figure 11e). As the temperature rises to 300 °C, the dimples at the fracture become deeper and larger (Figure 11i). At 200 °C, a small number of microcracks extending along the particles can still be observed on the fracture surface of the 3GW93 alloy, and the number of dimples has increased compared to room temperature (Figure 11d). When the temperature rises to 250 °C, no microcracks appear at the fracture surface, and the number of dimples increases (Figure 11h). As the temperature continues to rise to 300 °C, the morphology of the dimples at the fracture surface shows that they become deeper and larger, and the Mg₁₂La phase is found at the bottom of the dimples (Figure 11l).

The fracture mechanism of the Mg-9Gd-3Y(-La-Zn) alloy at room temperature was dominated by quasi-cleavage fractures, with cracks preferentially nucleating and propagating along second-phase interfaces [42]. With the increase in temperature, the GW93 and 1GW93 alloys underwent a fracture mode transition from the typical quasi-cleavage fracture presented at room temperature to the ductile fracture dominated by the dimple structure, and at the same time, the fracture strain of each of these alloys was improved. The fracture morphology of the 2GW93 and 3GW93 alloys shows ductile fractures with denser and deeper dimples as the temperature increases. The improvement of dislocation movement under high-temperature conditions and the synergistic effect of the multi-slip system effectively enhanced the plastic deformation capacity of the alloys, ultimately achieving the transformation from fracture mode to toughness and an increase in fracture strain.

4. Conclusions

1.

The co-addition of La and Zn promoted the precipitation of the particle second (Mg₅Gd and Mg₁₂La) phase and the LPSO phase, thereby restraining DRX behavior during the thermal deformation process. As a result, the Mg-9Gd-3Y-0.6La-1Zn alloy had the smallest grain size and the most second phase.

2.

The synergistic effects of grain refinement and second-phase strengthening (LPSO + Mg₁₂La) from La and Zn addition endowed the Mg-9Gd-3Y-0.6La-1Zn alloy with superior room temperature strength, demonstrating 284 MPa in YS and 355 MPa in UTS.

3.

At high temperatures, due to the pinning effect of a large amount of Mg₁₂La and LPSO phases relative to the grain boundaries, as well as the obstruction of dislocation movement, the Mg-9Gd-3Y-0.6La-1Zn alloy exhibited relatively excellent mechanical properties at 200–300 °C. At 300 °C, its YS was 249 MPa and its UTS was 285 MPa.