1 Introduction

Mg alloys, as one of the lightest metal structure materials, are frequently utilized in the aerospace, communications, electronics, transport and machinery manufacturing fields due to their superior specific stiffness and specific strength, efficient electromagnetic shielding performance, effective vibration and noise reduction performance and ease of recyclability [1, 2]. However, the hexagonal close-packed structure of Mg and its alloys results in relatively low strength and poor room temperature ductility, which greatly limits their practical industrial application [3].

Alloying with RE is one of the most effective methods for improving mechanical properties of Mg alloys [4]. Among them, Mg-Gd and Mg-Gd-Y alloys have been widely studied due to their excellent mechanical performance [5, 6]. With the addition of 15Gd (wt.%) in Mg, the tensile yield strength (TYS) increased from 131 to 290 MPa, while the elongation (EL) decreased from 35% to 18% [7]. The addition of Gd facilitated solid solution strengthening and precipitation strengthening, but the dispersed Mg₅Gd phase within the grains tended to cause stress concentration and thus decreased the ductility. In addition, with the addition of 1Y to Mg-4Gd-0.5Zr (wt.%), the TYS substantially increases from 88 MPa to 134 MPa, while the EL only slightly decreases from 44.6% to 39.0% [8]. The addition of Y not only improves the strength by refining grains, but also ensures ductility by facilitating the activation of non-basal slip. The strength of the Mg-RE alloy can be further improved after ageing. After ageing treatment, the TYS of the forged Mg-13Gd alloy reaches 430 MPa when the EL is 3.0%. The precipitation of the dense β' phase on the prismatic plane effectively impedes the movement of basal dislocations, thereby greatly improving the strength but decreasing ductility [9]. Therefore, the addition of Gd and Y to Mg can improve the mechanical properties via fine grain strengthening, solid solution strengthening and precipitation strengthening. However, high strength often requires a high Gd content, which is not conducive to the development of lightweight materials and energy savings.

For the development of energy saving and lightweight materials, light RE Nd has been developed in Mg-Gd-Y alloys [10]. The addition of Nd in Mg-Gd-Y alloys leads to significant precipitation strengthening [11, 12]. Due to the pronounced precipitation strengthening, the aged Mg-8Gd-3Y-1Nd-0.5Zr alloy can have a TYS as high as 415 MPa but an EL of only 2.0% [13]. Obviously, an increase in the RE content improves the strength but decreases the ductility of the alloy [14, 15]. It was reported that the strength of aged Mg-xGd-3Y-0.5Zr alloys increased from 188 MPa to 245 MPa with increasing Gd content from 6 wt.% to 10 wt.%, but the EL decreased from 8.7% to 3.4% [16]. Furthermore, In Ref. [17], it was reported that with increasing Nd and Gd contents, the amount of second phases in the alloy increases significantly, resulting in an increase in the strength but a decrease in the ductility of the alloy. However, in this work, the ductility of the Mg-9Gd-2Y-3Nd-0.2Zr alloy after peak ageing treatment was significantly better than that of the Mg-9Gd-2Y-1Nd-0.2Zr alloy. Therefore, it is necessary to compare the microstructural characteristics of Mg-9Gd-2Y-1Nd-0.2Zr alloy and Mg-9Gd-2Y-3Nd-0.2Zr alloy to elucidate the reason for the improved ductility of Mg-9Gd-2Y-3Nd-Zr alloy after ageing treatment.

In this study, hot extrusion and ageing treatment experiments were designed for Mg-9Gd-2Y-1Nd-0.2Zr and Mg-9Gd-2Y-3Nd-0.2Zr alloy, and the microstructural characteristics and mechanical properties of the two alloys were compared in detail. Significant grain boundary precipitate characteristics were observed in the extruded Mg-9Gd-2Y-3Nd-0.2Zr alloy. The reason for the improved ductility of the aged Mg-9Gd-2Y-3Nd-0.2Zr alloy was explained by comparing the precipitation characteristics at the grain boundary. It is hoped that these results will shed some light on the improvement in the ductility of aged Mg-RE alloys.

2 Experimental methods

Mg-9.00Gd-1.99Y-1.30Nd-0.18Zr (GWN921K, wt.%) and Mg-8.97Gd-1.95Y-2.87Nd-0.18Zr (GWN923K, wt.%) alloys were prepared from high purity (99.99%) Mg and Mg-30%Gd/Y/Nd/Zr master alloys under a mixed Ar and SF₆ protective atmosphere in an electric furnace. Then, the ingot was cut into cylindrical samples with dimensions of d85 mm×150 mm. The homogenization-treatment was conducted at 525 ℃ for 24 h to eliminate the eutectic structure. The ingots and molds were preheated at 470 ℃ for 2 h prior to extrusion. Then, the ingots were extruded into bars with a diameter of 18 mm at ram speed of 2 mm/s. After extrusion, the extruded samples were aged at 200 ℃ for 100 h.

Hardness testing was carried out on an HV-10 type Vickers micro indenter with a load of 4.9 N. Tensile testing was performed on an MTS 10 KN instrument along the ED at a speed of 2 mm/min. Dog-bone shaped specimens for tensile testing were prepared with a diameter of 5 mm and a gauge length of 25 mm. Three parallel specimens were taken for each set of tensile tests to ensure data accuracy. Samples for optical microscopy (OM) observation were prepared by mechanical grinding and polishing, and then they were etched for 30 s in a solution of 4 vol% nitric acid in ethanol. The preparation method of the samples used for scanning electron microscopy (SEM) observation was the same as that used for the OM sample. SEM observation was performed with a ZEISS EVO MA10 scanning electron microscope. The chemical composition of the second phase particles was analyzed by an equipped energy dispersive spectrometer (EDS). Samples for electron backscatter diffraction (EBSD) analysis were prepared by mechanical grinding, and then electropolished using the commercial electrolyte 70% perchloric acid cooled to -40 °C at a voltage of 25 V for 120 s. EBSD observation was performed on a ZEISS EVO MA10 scanning electron microscope equipped with an Oxford EBSD accessory, and HKL Channel 5 software was adopted for data analysis. Transmission electron microscopy (TEM, Tecnai G2 20 type operated at 200 kV) was used for observation. Thin foils for TEM analysis were prepared by punching discs with a diameter of 3 mm, mechanically grinding the discs to 50 μm, and then twin-jet electropolishing them in a solution of 3% perchloric acid, 1% nitric acid and 96% ethyl alcohol at 30 V and -35 °C.

3 Results

3.1 Microstructure of the initial alloys

Figure 1 shows the microstructures of the cast and homogenized GWN921K and GWN923K alloys and the corresponding EDS analysis results. Figure 1(b) shows an enlarged view of the yellow rectangle in the microstructure of the as-cast GWN921K alloy. The as-cast microstructures of the studied alloy consist of a large number of reticulated eutectic phases distributed in the Mg matrix and white cubic phases rich in RE elements (as indicated by the yellow arrows in Figure 1(b)). After homogenization, the reticulated eutectic phases distributed along the grain boundaries dissolved into the matrix, leaving a large number of irregularly shaped granular phases inside the grains and at the grain boundaries, which was more obvious in the GWN923K alloy, as shown in Figures 1(d), (e), (f).

Figure 1全屏查看

Microstructures of the as-cast (a, b, c) and homogenization treated (d, e, f) samples: (a, b, d) GWN921K alloy; (c, e, f) GWN923K alloy

3.2 Microstructure of extruded alloys

Figure 2 shows OM micrographs of the central part of the extruded specimens parallel/perpendicular to the ED, as well as the corresponding back-scaterred electron (BSE) micrograph along the ED. A comparison of the OM microstructures of the GWN921K and GWN923K alloys reveals that the microstructures of the alloys are similar, with both being dominated by DRXed grains, as shown in Figures 2(a), (b) and (d), (e). In addition, a large amount of the residual phase is broken into fine particles along the ED, characterized by a streamline distribution. This feature is more pronounced in the GWN923K alloy, indicating that it is closely related to the content of Nd in the alloy. In terms of grain size, the GWN923K alloy exhibited finer grains, suggesting that the distribution of the particle phase has a hindering effect on grain growth. The corresponding BSE micrographs show that, in addition to the dense distribution of bulk particle phases mainly along the ED, there are also many fine particle phases diffusely distributed in the alloy. All of the second phases exhibit three characteristic morphologies, nearly spherical, square and rod-like. According to the EDS results, these phases are essentially enriched in Gd, Nd and Y, which are typical of the Mg₅RE phase formed after extrusion of Mg-RE alloys. In addition, a small amount of Zr-rich phase in the form of rods was found at the grain boundaries of the GWN921K alloy, indicating that the residual clustered Zr-rich particle phases were extruded into rods along the grain boundaries during extrusion.

Figure 2全屏查看

OM and SEM images of the extruded alloys: (a, b, c) GWN921K; (d, e, f) GWN923K; (g, h) Corresponding EDS results of points in (c) and (f)

Figure 3 shows the orientation image maps (OIM), the grain size distribution and the {0001} pole figure and inverse pole figure of the studied alloys. As shown in Figures 3(a), (d), the same color represents grains with similar orientations, high angle grain boundaries (HAGB, misorientation angle >15°) are indicated by black, and low-angle grain boundaries (LAGB, misorientation angle with 2°-15°) are indicated by white. Both the extruded GWN921K and GWN923K alloys formed a homogeneous and approximately equiaxed DRXed microstructure. The average grain size (AGS) of the GWN921K alloy was 21.2 μm and that of the GWN923K alloy was 12.6 μm.

Figure 3全屏查看

Orientation image maps (a, d), grain size distribution maps of the extruded alloys (b, e), (0001) pole figures and inverse pole figures (c, f): (a, b, c) GWN921K; (d, e, f) GWN923K

The textural characteristics in Figures 3(c) and (f) show an abnormal extrusion texture with the {0001} plane perpendicular to the ED with textural intensities of 5.74 and 6.07, respectively. It is also noteworthy that in the OIM of the GWN923K alloy, some black precipitates are found along the ED, as shown by the white arrow in Figure 3(d). The fine second phase in Mg alloys can act as nucleation sites for DRX during hot deformation, inducing DRX behavior through a particle-stimulated nucleation (PSN) mechanism. Clearly, the smaller grain size of the GWN923K alloy compared to that of the GWN921K alloy is necessarily attributable to the influence of the second phase during hot extrusion.

Figure 4 shows bright field (BF) images, and corresponding selected area electron diffraction (SAED) and high-resolution transmission electron microscope (HRTEM) images of the precipitate in the extruded GWN921K and GWN923K alloy. The precipitates are mainly distributed inside the grains and are barely visible at the grain boundaries in the extruded GWN921K alloys, as shown in Figure 4(a). However, in the extruded GWN923 alloy, many micron-scale granular and triangular precipitates were observed at the grain boundaries. According to the SAED results, the granular precipitates were the Mg₅RE phase, and the triangular precipitates were the Mg₁₂RE phase. The Mg₅RE phase is more common in Mg-RE alloys [18-20], so the HRTEM of the Mg₁₂RE phase was characterized here in detail. Figure 4(f) shows HRTEM images of the interfacial regions between Mg₁₂RE phase particles and Mg matrix, together with the corresponding fast Fourier transform (FFT) results of Mg₁₂RE phase particles and Mg matrix region. When the incident beam is parallel to the [111]_Mg12RE and direction, the (121)_Mg12RE is found to be parallel to the . The interplanar spacings for the (121)_Mg12RE and planes are approximately 0.43 nm and 0.25 nm, respectively [21]. The orientation relationship between the Mg₁₂RE and Mg matrix can be described as [111] _Mg12RE// or (121) _Mg12RE // with a semi-coherent interface. Block-shaped particles are formed during casting, and fragmented during extrusion [22-24]. This multiscale second phase structure effectively impedes dislocation slip, and plays a crucial role in improving the mechanical properties of Mg alloys.

Figure 4全屏查看

TEM images of (a) GWN921K alloy, (b-f) GWN923K alloy: (a, b, d) Bright field; (c) Corresponding SAED pattern of Mg₅RE; (e) Corresponding SAED pattern of Mg₁₂RE; (f) HRTEM of Mg₁₂RE

3.3 Age-hardening response and mechanical properties

Figure 5 shows the hardness curves of the extruded GWN921K and GWN923K alloys aged at 200 ℃. The hardness of the GWN923K alloy in the as-extruded state is 94HV, which is higher than that of the extruded GWN921K alloy. This is attributed to the finer grain size of the extruded GWN923K alloy. The curves show that the alloys in the extruded state all exhibit a pronounced age-hardening response. In the initial stage, the hardness increases significantly with ageing time, reaches the peak hardness, remains stable for a period, and then starts to decrease slowly after 70 h. The peak hardness of the extruded GWN921K and GWN923K alloys was 127HV and 130HV after ageing for 50 h and 45 h, respectively, which was 38% higher than that of the extruded alloy.

Figure 5全屏查看

Hardness curves of the studied alloys aged at 200 ℃

The engineering tensile strain curves of the studied alloys under the as-cast, homogenized, extruded and peak-aged (T5-200 ℃) conditions are shown in Figure 6. The values of tensile mechanical properties are summarized in Table 1. Prior to extrusion, the TYS of the GWN923K alloy was greater than that of the GWN921K alloy, but significantly less ductile than that of the GWN921K alloy. The higher strength is attributed to the strengthening of the diffuse second phases and their pinning effect on the grain boundaries. The poor ductility of the homogenized GWN923K alloys is due to the formation of agglomerates of a large number of second phases, which promotes crack initiation.

Figure 6全屏查看

Typical tensile stress-strain curves of the studied alloys at room temperature

Table 1全屏查看

Mechanical properties of the studied alloys under different conditions

Alloy		UTS/MPa	TYS/MPa	EL/%
As-cast	GWN921K	232±2	149±2	7.1±0.8
	GWN923K	224±2	147±1	5.2±0.2
Homogenized	GWN921K	233±1	127±2	10.6±0.6
	GWN923K	222±2	149±2	4.7±0.1
Extruded	GWN921K	329±2	202±5	12.8±1.0
	GWN923K	313±1	212±2	13.7±1.3
T5-200 ºC	GWN921K	397±2	275±2	2.6±0.7
	GWN923K	388±2	281±1	6.2±0.8

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After extrusion, the ultimate tensile strength (UTS), TYS and EL of the GWN921K alloy were 329 MPa, 202 MPa and 12.8%, while the UTS, TYS and EL of the GWN923K alloy were 313 MPa, 212 MPa and 13.7%, respectively. The slight decrease in the UTS of the GWN923K alloy is attributed to the increase in Nd content which promotes precipitation and thus weakens the solid solution strengthening effect. The increase in TYS is attributed to fine grain strengthening and precipitation strengthening. Ageing heat treatment of Mg-RE alloys results in a substantial increase in strength but decreases ductility. The UTS, TYS and EL of the peak-aged GWN921K alloy are 397 MPa, 275 MPa and 2.6%, respectively. Notably, the peak aged GWN923K alloy achieves a ductility of 6.2% when the UTS and TYS are increased to 388 MPa and 281 MPa, respectively. An increase in the Nd content slightly reduced the UTS of the alloy, but increased the yield strength and ductility, which improved the balance between the strength and ductility of the peak-aged GWN923K alloy.

3.4 Microstructure of peak aged alloys

Mg-RE alloys undergo peak ageing treatment to significantly increase their strength. Figure 7 shows the HADDF-STEM and TEM results of the peak-aged GWN921K and GWN923K alloys. The direction of the incident electron beam is parallel to [0001]α_-Mg. After peak ageing treatment, a large number of nano-scale phases were formed inside the grains. β' phase precipitation on prismatic planes can effectively hinder dislocation slip on the basal plane of Mg alloys, thus improving their mechanical properties [25]. Compared with the GWN921K alloy, the precipitated β' phase in the GWN923K alloy is denser and finer. This means that the GWN923K alloy exhibits a stronger age-strengthening response. Furthermore, in Figure 7(c), the β' phase in the GWN921K alloy crosses the grain boundaries, whereas a typical PFZ is formed along the grain boundaries in the GWN923K alloy. This difference originates from the large amount of precipitation distributed at the grain boundaries of the GWN923K alloy after extrusion.

Figure 7全屏查看

HAADF-STEM and BF-TEM images of extruded (a, b, c) GWN921K and (d, e, f) GWN923K alloys after peak-aged treatment

4 Discussion

In this study, the microstructure and mechanical properties of extruded Mg-9Gd-2Y-1Nd-0.2Zr and Mg-9Gd-2Y-3Nd-0.2Zr alloy were investigated, and it was found that a large number of grain boundary precipitates were formed in the extruded Mg-9Gd-2Y-3Nd-0.2Zr alloys. The grain boundary precipitates inhibit grain growth on the one hand, and promote the formation of PFZs during the ageing treatment on the other hand, which improves the room temperature tensile ductility of the peak aged alloy, and results in good comprehensive mechanical properties in the Mg-9Gd-2Y-3Nd-0.2Zr alloy.

Generally, the low-temperature ageing treatment of Mg-RE alloys does not have any effect on their grain size and texture [26, 27]. The microstructure of the extruded alloys shows a full DRXed grain and <0001>//ED texture. The formation of the abnormal extrusion texture can be broadly attributed to the segregation of solute atoms at the grain boundaries, the preferred growth of grains and the activation of pyramidal <c+a> slip [28-30]. It was noted that the formation of this texture favors the ductility of the alloy tensile along the ED [31]. Furthermore, according to the classical Hall-Petch theory, it is known that the strength of the alloy gradually increases with grain refinement [32]. The fine DRXed grain boundaries can effectively impede the movement of dislocations, which results in a large number of dislocations being plugged, thus improving the strength [33-35]. Moreover, the plastic deformation of fine grains under external force can be distributed to more grains, which makes the plastic more uniform, thus improving the ductility [36, 37]. Therefore, grain refinement is conducive to improve the strength and ductility of alloys. The effect of texture on the mechanical properties can be characterized by the Schmid factor (SF). Figure 8 shows SF distribution histograms and SF maps for the basal <a>, prismatic <a>, pyramidal <a>, and pyramidal <c+a> slip systems for the GWN921K and GWN923K alloys when the loading direction was parallel to the ED. It is clear that the increase in Nd content has a small effect on the average SF of the different slip systems. That is, the contribution of the texture to the strength and ductility of the alloys after the increase in Nd content in this study is almost negligible. In addition to the grain size, the characteristics of the precipitates in the studied alloy are particularly important in influencing the mechanical properties.

Figure 8全屏查看

SF maps and corresponding distribution histograms for various slip systems of extruded GWN921K and GWN923K alloys

An increase in the Nd content significantly promoted the precipitation of secondary phases in the extruded alloy, especially the Mg₅RE and Mg₁₂RE phases at the grain boundaries. Figure 9 illustrates the typical grain boundary characteristics of the peak-aged GWN921K and GWN923K alloys. When the Nd content increased from 1 wt.% to 3 wt.%, the GWN923K alloy formed a pronounced grain boundary precipitate, and a typical PFZ formed between the grain boundary precipitate and the ageing precipitated β′ phase with a bandwidth of approximately 0.13-0.15 μm. During ageing treatment, the solute atoms tend to strongly polarize toward the grain boundary precipitate, leading to the depletion of nearby solutes, and the formation of typical PFZs [38, 39]. It was also reported the formation of PFZs in a forged Mg-Gd-Y-Zr alloy, and noted that PFZs could effectively improve the deformability [40-42]. Generally, the high RE content promotes the formation of the second phase in the alloy, which agglomerates after deformation to spatially split the α-Mg matrix, causing the alloy to initiate microcracks during tensile deformation [43, 44]. However, in this study, increasing the Nd content from 1 wt.% to 3 wt.% increased the ductility of the peak aged alloy by 58% with almost no change in strength. This is attributed to the fact that the second phase on the grain boundary not only impedes grain growth, but also promotes the formation of PFZs during ageing, and strengthens the grain boundary. In addition, the formation of PFZs reduces the number of second-phase particles near the grain boundaries, which reduces crack initiation and the tendency to fracture along the grain boundaries. Therefore, to further improve the balance between the strength and ductility of Mg-RE alloys after ageing treatment, the formation of PFZs can be promoted by adding appropriate amounts of RE and combining them with deformation to regulate the distribution characteristics of grain boundary precipitates.

Figure 9全屏查看

TEM-BF results of grain boundary characteristics in peak-aged (a) GWN921K and (b, c) GWN923K alloys

5 Conclusions

In this paper, a detailed comparison of the microstructure and room temperature mechanical properties of the extruded and peak-aged Mg-9Gd-2Y-xNd-0.2Zr (x=1 wt.%, 3 wt.%) was investigated. Several conclusions are reached, as follows:

1) Compared to the as-cast and homogenized Mg-9Gd-2Y-1Nd-0.2Zr alloys, Mg-9Gd-2Y-3Nd-0.2Zr alloy exhibits high strength and low ductility, which is attributed to the formation of a second phase during the casting process.

2) After extrusion, a large number of second phases in the Mg-9Gd-2Y-3Nd-0.2Zr alloy were distributed along the extrusion direction to form the fibrous microstructure. The second phases can be recognized as Mg₅RE and Mg₁₂RE phases, which are mainly distributed on the grain boundaries.

3) The grain boundary precipitate in the Mg-9Gd-2Y-3Nd-0.2Zr alloy promoted the formation of PFZs during the ageing treatment, which results in a 58% increase in the ductility of the peak-aged Mg-9Gd-2Y-3Nd-0.2Zr alloy compared to that of the peak-aged Mg-9Gd-2Y-1Nd-0.2Zr alloy.