Effect of trace Al and Ca co-addition and solution treatment on the microstructure and mechanical properties of Mg-10Gd alloys

To reduce the density and cost while improving the ductility of Mg-Gd system alloys, a total of 2 wt% Al and Ca were added to an Mg-10Gd alloy for partial substituting Gd element. The prepared alloys were subjected to solution treatment at 500 °C for several hours, and the variations in microstructure and mechanical properties were investigated. To reveal the effect of the combined addition of Al and Ca on the microstructure and mechanical properties of the Mg-10Gd alloys, Mg-10Gd alloys containing the same amount of Al or Ca were also fabricated and characterized. By comparison, it was found that only the co-addition of Al and Ca simultaneously improved the hardness, ultimate tensile strength, yield strength, and elongation of the Mg-10Gd system alloys in both as-cast and solid solution states.


Introduction
As the lightest metallic structural materials with high specific strength, specific stiffness, good damping properties, and biocompatibility, Mg and its alloys possess a wide range of promising applications in aerospace, automotive, electronics, and medical fields [1][2][3][4][5]. However, their absolute strength and creep resistance are low, hence their application scopes are still far from those of steel and aluminum alloys. Rare Earth (RE) elements are very effective strengthening elements for Mg alloys, such as Y, Nd, Sm, Gd, and Dy [6][7][8]. They act as obstacles for dislocation movement as solute atoms and can also be used for grain boundary pinning in the form of intermetallic phases. Among these alloying elements, Gd has been broadly introduced into the Mg alloys in recent years due to its comparable large solubility in α-Mg [9]. In addition, stable intermetallic Mg 5 Gd phases are generally precipitated in the Mg-Gd system alloys contributing to excellent high-temperature mechanical properties and creep resistance [10][11][12]. However, high Gd addition to Mg alloys will increase their cost and decrease ductility. Hence, it is necessary to seek other lighter and cheaper alternative elements to partially replace the Gd element.
To improve the mechanical properties while reducing cost, several non-rare Earth elements were added to Mg-Gd series alloys forming a new intermetallic, and the morphology and distribution of the intermetallic were controlled through an appropriate heat treatment process [13][14][15]. Dai et al [16] found that when 0.8 wt%-1.3 wt% of Al was added to an Mg-10Gd alloy, the in situ generated Al 2 Gd particles provided heterogeneous nucleation sites for α-Mg grains, and restricted grain coarsening during solution treatment. Pourbahari et al [17] added Al to Mg-6%Gd-1%Zn (GZ61) to partially replace Gd, the generated intermetallic compounds in the alloy are mainly (Mg, Al) 3 Gd, Mg 5 Gd, and Al 2 Gd phases, and with the increase of Al content, the content of Al 2 Gd gradually increased, and the refinement effect of α-Mg grains became increasingly significant. When the Al content was equal to Gd content, the tensile strength and elongation of Mg-3%Gd-3%Al-1%Zn alloys increased by 57.6% and 250.0%, respectively compared with Mg-6%Gd-1%Zn. However, with the further increase of Al content, the size of Al 2 Gd further increased and the (Mg, Al) 3 Gd phase appeared with a continuous Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. reticulate structure along grain boundaries, which was detrimental to the mechanical properties of the alloy. The crystal structure of the alkaline Earth element Ca (a = 0.623 nm, c = 1.012 nm) is similar to that of Mg, and an appropriate amount of Ca can promote the nucleation of α-Mg grains, improving the constitutional supercooling of the diffusion layer at the liquid-solid interface, reducing the speed of eutectic reaction, thereby facilitating grain refinement. In addition, the Mg-Ca intermetallic Mg 2 Ca has good thermal stability, which can hinder the grain boundary slip and reduce the diffusion rate of solute at the grain boundary, thereby improving the room-temperature strength and creep resistance of the Mg alloys [18,19]. Shi et al [20] reported that after adding 0.3 wt% Ca to Mg-10Gd-0.5Zr alloys, the main phase composition did not change (α-Mg and Mg 5 Gd), and the yield strength was also nearly unchanged. When the Ca content was increased to 1.2 wt%, a fine reticulated eutectic Mg 2 Ca phase appeared in the alloys, and the tensile yield strength was increased by 17% compared with that of Mg-10Gd-0.5Zr alloys. Moderate addition of Ca can refine the grain size, but excessive addition of Ca will not only cause grain coarsening but also cause the formation of brittle eutectic compounds at the grain boundaries, deteriorating the tensile strength and elongation of the alloy [19].
Gd element has a high equilibrium solid solubility in Mg, reaching 4.53 at%/23.49 wt% at 548°C. When subjected to solid solution treatment, a supersaturated solid solution of α-Mg with Gd will be formed, in which Gd exists as replacement atoms in α-Mg causing lattice distortion of α-Mg, thereby achieving solid solution strengthening. In addition, solid solution treatment can also be used to regulate the phase composition and the morphology of intermetallic compounds in Mg-Gd-Al or Mg-Gd-Ca alloys, thereby improving the mechanical properties of Mg alloys. Zhuang et al [21] investigated the effect of solid-solution treatment on the microstructure of Mg-9Gd-xAl (x = 0.4, 0.6, 0.8, 1.0, 3.0 wt%) alloy, and found that after solid solution treatment, the short rod-like Mg 5 Gd, fishbone and petal-like (Mg, Al) 3 Gd phases dissolved into the matrix, and only the morphology of the Al 2 Gd particles remained unchanged. When the Al content was less than 1 wt%, lamellar (Mg, Al) 2 Gd phases were precipitated within the α-Mg grains. The tensile strength of the solutiontreated Mg-9Gd-0.6Al alloy was significantly increased compared with the as-cast one, and the main strengthening mechanisms were solid solution strengthening and secondary phase strengthening. Shi et al [20] reported that Mg 5 Gd phase in Mg-10Gd-1.2Ca-0.5Zr alloy disappeared after solid solution treatment, and the Mg 2 Ca phase became coarse and the tensile yield strength of the alloy slightly decreased, but both the tensile strength and elongation were increased.
In summary, the singular addition of small amounts of Al or Ca can improve the mechanical properties of Mg-Gd system alloys, but little research has been reported on the microstructure and mechanical properties of Mg-Gd system alloys with the combined addition of Al and Ca. In the present study, the effect of the combined addition of Al and Ca on the microstructure of an Mg-10Gd alloys in as-cast and solid solution state was systematically investigated compared with the Mg-10Gd based alloys containing singular Al or Ca, and the Vickers hardness and tensile properties of these alloys were also compared with each other. Finally, the strengthening and toughening mechanism of the alloying elements and the solid solution treatment were discussed.

Experimental process
The Mg-10Gd alloys with combined addition of Al and Ca (Mg-10Gd-1Al-1Ca) were prepared from pure magnesium ingot (99.9%), pure aluminum ingot (99.9%), Mg-30Gd, and Mg-30Ca master alloys in SG2-7.5-12 A type resistance furnace. When the pure Mg ingot and pure Al ingot were completely melted at 700°C, desired amounts of Mg-30Gd and Mg-30Ca master alloys were added to the melt, and the melt temperature was increased to 750°C and kept for 20 min. Finally, the melt was skimmed and cooled to 700°C again and then poured into a steel permanent mold. To prevent oxidation and combustion of the melt, the whole melting process was protected with RJ-2 covering flux under the argon atmosphere. For comparison, Mg-10Gd alloys, and Mg-10Gd alloys containing singular Al or Ca were also fabricated. The composition of the designed alloys is shown in table 1.
Cubic samples for solid solution treatment with a dimension of 10 mm × 10 mm × 10 mm were machined from the center of the ingot, and the solid solution treatment was carried out in a box-type resistance furnace. To avoid oxidation and combustion, the solution temperature is usually selected below the eutectic temperature, which is 548°C according to the phase diagram of Mg-Gd binary alloy [22], and the solution temperature of Mg-Gd based alloys is usually selected in the range of 500°C-530°C [10,21,23]. Considering the addition of Al and Ca in the alloys is small, these alloys were solution treated at the same temperature of 500°C for various hours (4 h, 8 h, 16 h, and 24 h) followed by quenching in water at 20°C. The metallographic specimens before and after solid solution treatment were ground and polished and then etched with an acetic-picral reagent (60 ml ethylene glycol, 20 ml acetic acid, 1 ml concentrated nitric acid, and 19 ml distilled water). Olympus GX71/50-2000 optical microscope and JSM-6700F field emission scanning electron microscope were used to observe the microstructure of the alloys. Tensile tests were performed on the HT-2402 material testing machine with a constant rate of 1 mm min −1 . The tensile testing for each alloy was repeated three times using dog-bone-shaped specimens with a dimension of 15 mm × 5 mm × 2 mm in the gauge section according to the ASTM-E8 standard. Vickers hardness was determined on a TUKON 2100B Vickers hardness tester with a loading force of 5 Kg for a dwell time of 15 s, and an average value was obtained through five tests at different positions of each specimen.

Results and discussion
3.1. Microstructure Figure 1 shows the XRD diffraction patterns of the four as-cast Mg alloys. It can be seen from figure 1 that the Mg-10Gd alloy is mainly composed of α-Mg and Mg 5 Gd. The diffraction peaks of Al 2 Gd and (Mg, Al) 3 Gd phases appear after adding 2 wt% Al elements to the Mg-10Gd alloy, and the peak intensity of the Mg 5 Gd phase decrease sharply due to the consumption of Gd by the Al 2 Gd and (Mg, Al) 3 Gd phases. When 1 wt% of Al and 1 wt% of Ca is added to the Mg-10Gd alloy, the diffraction peaks of Mg 2 Ca appear, while the diffraction peaks of Al 2 Gd and (Mg, Al) 3 Gd phases nearly disappear due to the decrease of Al content. When only 2 wt% Ca is added to the Mg-10Gd alloy, the amount of the diffraction peaks of Mg 2 Ca increases obviously, and some α-Mg peaks are slightly shifted toward higher 2θ due to the solid solution of Ca in the α-Mg changing the lattice parameters of α-Mg. Figure 2 shows the microstructure of the four as-cast Mg alloys, it can be seen that alloy 1 # is composed of coarse columnar dendrites, as shown in figure 2(a). When singular 2 wt% Al is added to the Mg-10Gd alloys, the microstructure is refined and the crystals are transformed into equiaxed grains with an average size of 73.05 μm, as shown in figure 2(b). When 1 wt % Al and 1 wt % Ca is added to the Mg-10Gd alloys, the α-Mg still exhibite equiaxed grains, but the average size increases slightly to 84.98 μm. When singular 2 wt % Ca is added to the Mg-10Gd alloys, the equiaxed grains are transformed into columnar dendrites, and the dendritic arms are significantly coarsened. The grain refinement of alloy 2 # and alloy 3 # can be attributed to the Al 2 Gd phase having similar lattice constants to the α-Mg, which provides heterogeneous nucleation sites for α-Mg [16]. The grain refinement effect in alloy 3# is less significant compared with that of alloy 2# mainly due to the less Al addition, causing fewer Al 2 Gd phases as effective nucleants in alloy 3#. The transition of equiaxed grains in alloy 3# to dendrite grains in alloy 4# indicates that Al is more effective than Ca as a grain refiner for Mg-10Gd alloys. As a result, the α-Mg grains coarsen with decreasing Al addition as shown in figure 2 To further identify the phase composition of the alloys, EDS analysis was performed, figure 3 presents the SEM images of the four as-cast Mg alloys, it can be seen from figure 3(a) that alloy 1 # consists of a gray-white reticulate region and bright white dotted phases. The bright white dotted phase is identified as Mg 5 Gd phase according to the atomic ratio of Mg to Gd (5:1) combining with the XRD patterns, and the gray-white reticulate regions are the Gd enriched regions. Figure 3(b) exhibits the SEM micrographs and EDS results of alloy 2 # , as observed in figure 3(b), after singularly adding Al element, a massive petal-like phase and a small number of bright-white particles appear in the alloy. The atomic ratio of Al to Gd in the bright-white particles is about 2:1, and it is inferred that the particles are Al 2 Gd phase. The atomic ratio of (Mg, Al) to Gd in the petal-like phases is about 3:1, thus it is inferred that they are (Mg, Al) 3 Gd phase. Figure 3(c) illustrates the SEM micrographs of alloy 3 # , and it is observed that both the size and number of the petal-like (Mg, Al) 3 Gd phases in alloy 3 # decrease compared with that in alloy 2 # . As the Al content decreases, the driving force for precipitation and the migration rate of the atom decrease, the size of the (Mg, Al) 3 Gd phase consequently decreases. As the migration rate of atomic slows down, the longitudinal growth rate decreases and causes the (Mg, Al) 3 Gd phase to be distributed in the form of particles or short rods. Figure 3(d) shows SEM micrographs of alloy 4 # , and two kinds of intermetallic with different morphologies were observed in the enlarged insert, i.e., gray-white rod-like phases and black reticulate phases. Similar intermetallic was also reported by Shi et al in Mg-10Gd-1.2Ca-0.5Zr alloy [20], and the gray-white rod-like phases and black reticulate phases are Mg 5 Gd and Mg 2 Ca, respectively, and they formed a ternary eutectic phase together with α-Mg through eutectic reaction [24]. Figure 4 shows the microstructure of alloy 1 # after solution treatment for different hours. The volume fraction of the second phases was statistically analyzed by ImageJ software from at least five SEM images at low magnification. The volume fraction of the second phases in alloy 1# ( figure 3(a)) is 7.89%. After solution treatment for 4 h, the volume fraction decreases to 3.15% ( figure 4(a)). When the solution time reaches 8 h, the volume fraction further decreases to 2.63% ( figure 4(b)). When the solution time is extended from 16 h to 24 h, there is no obvious change in the volume fraction of the second phases (figures 4(c) and 4(d)), indicating that most Mg 5 Gd phase had been dissolved into the Mg alloy, similar phenomenon was also reported by Shi et al. in Mg-10Gd-1.2Ca-0.5Zr alloys solution treated at 495°C for 16 h [20]. However, there still exist a small amount of white particle phases in the solution-treated alloys, as shown in figure 4(d). Figure 5 shows the SEM micrographs of the residual particle phases and EDS results. According to the EDS results, it is inferred that the particles are Mg 2 Gd phases. Under equilibrium solidification conditions, the composition of alloy 1 # should be α-Mg solid solution + Mg 5 Gd phase, but the melt was cast into a steel mold without preheating, thus the solidification proceeded at a high cooling rate in a non-equilibrium state. In this case, the solute atoms do not have enough time to diffuse, so the Mg 2 Gd phases were preferentially precipitated [25]. The Mg 2 Gd phase has a high melting point and a low diffusion coefficient [26], so it exhibits good thermal stability during the solid solution treatment. However, the Mg 2 Gd phase was not detected in the XRD patterns due to its minor amount. Figure 6 shows the microstructure of the alloy 2 # after solid solution treatment for different solution hours. After 4 h solution treatment, the (Mg, Al) 3 Gd phases are partially dissolved, the petal-like morphology disappears, and a large number of fine rod-like phases parallel to each other precipitate inside the α-Mg. As the solid solution time increases, the (Mg, Al) 3 Gd phases are further dissolved, and the size the rod-like phases increases with increasing solution time up to 16 h and then decreases. Figure 7 demonstrates the enlarged SEM images of the parallel rod-like phases after different solution times. After 4 h solution treatment, the rod-like phases start to precipitate inside the grains. When the solution time reaches 8 h, the length of the precipitated phases becomes longer and the width does not change significantly. When the solution time rises to 16 h, the size of the precipitated phase reaches the maximum with an average length of 11.69 μm and an average width of 1.29 μm. When the solution time is prolonged to 24 h, the length of the precipitated phase decreases due to partial dissolution, and the morphology of the precipitated phase was transformed into parallel aligned particles in the α-Mg interior grain. According to EDS analysis, these precipitated phases are identified as Mg 90 Al 8.2 Gd 1.8 phases. These discontinuous rods/particles phases were also reported by Bian et al in Mg-14.02Gd-2.33Zn-1.8(Al-Sr) alloy [27], and Zhuang et al [21] also found a similar morphology evolution law of the rod-like phase during solid solution treatment in Mg-9Gd-1.0Al alloy.
According to the study of Gu et al [28], the short rod-like phases in α-Mg grain of Mg-Gd-Al alloys always grow longitudinally parallel to the [0001] direction of the Mg matrix. There exists a coherent relation between the rod-like phase and the Mg with a misfit of about 0.7%, and the rod-like phases nucleate and grow along [0001] direction with a low-energy interface due to good atomic matching in the (0001) plane. Hence, the rodlike phases are parallel to each other and have one orientation in one grain. As the rod-like phases grow, the supersaturation of the matrix decreases in the late stage of solid solution, the precipitation driving force decreases, and the growth kinetics is dominated by diffusion, under the driving force from solution concentration difference between small second phases and large second phases, the Al and Gd atoms near the small second phases diffuse to the larger second phases directionally, as a result, the large rod-like phases are thickened while the small rod-like phases disappear. In the later stages of solid solution treatment, since the spherical structure has the lowest surface energy and correspondingly the least Gibbs free energy, and the particle    precipitated phase is more thermal stable, the morphology of the rod-like phase is transformed to discrete particles when the solution time is prolonged to 24 h [29]. Figure 8 shows the microstructure evolution of alloy 4 # with different solution times. It is seen from figure 8(a) that the grain size increases significantly after 4 h solution treatment, and then nearly keeps stable with increasing solution time. According to the DSC curve of the alloy 4 # as shown in figure 9, the melting point of the Mg-Gd-Ca ternary eutectic phase (α-Mg + Mg 2 Ca + Mg 5 Gd) is 495°C [24]. Therefore, during the solid  solution treatment at 500°C, the eutectic phases were partially melted and segregated to triple boundary junctions. Figure 10 shows the microstructure evolution of alloy 3 # with various solution times. As shown in figure 10, the grain size does not change obviously as the solution time rise, and the Mg-Gd-Ca eutectic phases (α-Mg + Mg 2 Ca + Mg 5 Gd) at grain boundaries are partially dissolved. With the extension of solution time, the volume fraction of the eutectic phase further decreases, and when the solution time rises to 24 h, most eutectic phases are dissolved into the matrix. It is worth noting that neither coarsening of the eutectic phase nor abnormal grain growth occur in the alloy 3 # during the solid solution treatment.
Grain growth is a very common phenomenon in the solid solution treatment of Mg alloys, and if there are many defects and high energy at the grain boundaries, Mg alloys will spontaneously reduce the grain boundary area to reduce the total grain boundary energy, causing grain growth. During the solid solution treatment at 500°C, as the eutectic phase melts, the volume fraction decreases and the pinning effect on the grain boundaries weakens, resulting in rapid grain growth in alloy 4 # . However, the Al 2 Gd phases have a high melting point of 1525°C [17], thus keeping high thermal stability in the solution treatment at 500°C, as shown in figure 10(c), effectively pining the grain boundaries and hindering the migration of grain boundaries [16], consequently, the alloy 3 # did not undergo abnormal grain coarsening during the solution treatment. Figure 11 exhibits the hardness of the four Mg alloys at different solution times. As shown in figure 11, the hardness of alloy 1 # increases with the extension of solution time, reaching a maximum of 92.6 Hv after 24 h solution treatment. The hardness of alloy 2 # decreases at first and then increases slightly with the increase of solid solution time, reaching a minimum of 69.2 Hv at 16 h, and then slightly increasing to 78.1 Hv at 24 h. The hardness of the alloy 3 # decreases from 100.3 Hv to 92.9 Hv after 4 h solution treatment, and then increases to 117.2 Hv at 24 h. The hardness of the alloy 4 # decreases monotonically from 123.8 Hv to 63.3 Hv as the solution time rises.

Mechanical properties
The hardness of the Mg-10Gd system alloys is mainly associated with both the amount of alloy solute atoms and the strength of the grain boundary [10]. In the solution treatment process of alloy 1 # , the Mg 5 Gd phase gradually disappeared as the solution time increased, and the Gd was dissolved into the Mg matrix causing lattice distortion, achieving a solid solution strengthening effect, thereby enhancing the hardness of the alloy 1 # . The (Mg, Al) 3 Gd phases and Mg 5 Gd phases in the alloy 2 # were gradually decomposed in the solid solution process, and the Gd and Al elements were diffused to α-Mg, thus the strength of the grain boundaries gradually decreased. As the concentration of Gd and Al elements increased, the rod-like Mg 90 Al 8.2 Gd 1.8 phase precipitated inside α-Mg grains and gradually grew, acting as a secondary strengthening phase. Therefore, the hardness of the alloy 2 # decreased at first and then increased with increasing solution time. Compared with the as-cast alloy 2 # , the higher hardness of the as-cast alloy 4 # mainly attributes to the continuous reticulate eutectic phases strengthening the grain boundaries. However, there was no Al 2 Gd phase in alloy 4 # , and the eutectic phases at the grain boundaries increasingly dissolved during the solid solution treatment and lost their role in pinning grain boundaries. Consequently, α-Mg grain of the alloy 4 # was coarsened abnormally, and the hardness decreased with increasing solid solution time. Compared with the alloy 4 # , the hardness of the alloy 2 # did not decrease monotonically with increasing solution time, and it can be ascribed to the following aspects: first, the Mg 90 Al 8.2 Gd 1.8 phase improved the strength of the α-Mg as a secondary strengthening phase. second, the Al 2 Gd phase has good thermal stability and did not dissolve during the solid solution treatment, and still plays the role of inhibiting grain growth and impeding grain boundary sliding. Similar to alloy 2 # , the hardness of alloy 3 # decreased at first and then increased with extension of solution time. Alloy 3# exhibits higher hardness than alloy 2 # but lower than alloy 4 # in as-cast state, indicating that Mg 2 Ca eutectic phase has a dominant role in improving the hardness of Mg-10Gd based alloys. After the solution treatment, the grain size in Mg-10Gd based alloys containing Al kept stable due to the Al 2 Gd phase restricting grain coarsening. However, the residual Mg 2 Ca phases in alloys 3# contributed to a higher hardness than alloy 2 # . Figure 12 depicts the variations of the tensile mechanical properties of four alloys as a function of solution time. As shown in figure 12(a), the ultimate tensile strength (UTS) and elongation (EL) of alloy 1 # increase with the increasing solution time, reaching a maximum after 24 h solution treatment. The increase of UTS mainly attributes to the solid solution strengthening of Gd elements and the resultant lattice distortion, impeding dislocation movement. There are massive Gd segregation regions in the as-cast alloy 1 # , and there exist defects in the segregation regions, which may cause cracking and reduce the ductility of the alloy, while the solution treatment improves the ductility by reducing the segregation level of Gd. As shown in figure 12(b), the UTS and tensile yield strength (TYS) of the alloy 2 # changes very little, but the EL increases from 10.1% to 16.4% with increasing solid solution time. The (Mg, Al) 3 Gd phase in the as-cast alloy 2 # exhibits a petal-like shape with needle-like branches, and the tips of branches are prone to causes stress concentration. The morphology of the (Mg, Al) 3 Gd phase is transformed into discrete particles after the solution treatment, which effectively reduces the stress concentration and improves the elongation of the alloy. As shown in figure 12(c), the UTS, TYS, and EL of alloy 3 # increase with increasing solution time, and the strengthening mechanism is similar to that of alloy 2 # . The UTS, TYS, and EL of alloy 4 # decrease significantly with increasing solution time, as shown in figure 12(d). This is mainly due to the abnormal grain growth and the segregation of the Mg-Gd-Ca eutectic phase at the triple boundary junctions causing stress concentration. Table 2 lists the mechanical properties of the four alloys before and after the solution treatment. Both the UTS and EL of alloy 2 # are increased significantly before and after the solution treatment compared with alloy 1 # . However, similar to the variations of hardness, the TYS of alloy 2 # is also slightly lower than that of alloy 1 # after the solution treatment. The UTS, YS, and EL of alloy 3 # are improved simultaneously compared with that of alloy 1 # . The main strengthening and toughening mechanisms are the grain refinement by the addition of Al and Ca, and no abnormal grain coarsening occurring during the solution treatment. The alloy 3 # exhibits a higher TYS and a slight lower elongation than alloy 2 # before and after the solid solution treatment. It can be attributed to following two aspects: first, not only Al 2 Gd but also Mg 2 Ca phases play the role of reinforcements strengthening the grain boundaries. Second, the less addition of Al in alloy 3 # consumed less Gd to form intermetallic, thus more Gd atoms were diffused to α-Mg during the solution treatment, causing lattice distortion of α-Mg and achieving solution strengthening effect. It is concluded that the solution treatment is  capable of improving the strength and ductility of alloys 1 # , 2 # , and 3 # except for alloy 4 # . By comparing the strength and elongation of the four alloys, it is noticed that only the combined addition of Al and Ca improved the strength and ductility of the Mg-10Gd alloy simultaneously whether in the as-cast state or the solution state.

Conclusions
The effect of the minor addition of Al and Ca on the microstructure, hardness, and tensile properties of the ascast Mg-10Gd alloys have been investigated by changing their addition ratio, i.e., single addition of Al, single addition of Ca, and combined addition of Al and Ca (1:1 in wt%). The obtained alloys were subjected to solidsolution treatment at 500°C for several hours up to 24 h, and the variations of mechanical properties were investigated combined with microstructure characterization. Three main conclusions were drawn as follows: (1)The hardness and tensile properties of the as-cast alloy 2 # were improved significantly compared with that of alloy 1 # , especially the EL was increased by 140.5%. After the solution treatment, the mechanical properties were further enhanced. However, the TYS of alloy 2 # was lower than that of alloy 1 # after 24 h solution treatment.
(2)The as-cast alloy 4 # exhibited the highest hardness but the poorest ductility. Both their hardness and tensile properties decreased with increasing solution time due to the abnormal grain coarsening.
(3)The co-addition of Al and Ca contributed to the greatly improved hardness and tensile properties of alloy 3 # compared with alloy 1# in both as-cast and solid-solution states, and the main strengthening mechanism may be grain refinement, good grain thermal stability, second phase strengthening by the rod-like phase inside the α-Mg grain, and grain boundary strengthening by Al 2 Gd and Mg 2 Ca phases.
To sum up, Al and Ca co-addition combining solution treatment is a promising method to synchronously enhance hardness, strength, and ductility of Mg-Gd based alloys while reducing cost.