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Yifan Song, Xihai Li, Jinliang Xu, Kai Zhang, Yaozong Mao, Hong Yan, Huiping Li, and Rongshi Chen, Effect of annealing treatment on the microstructure and mechanical properties of warm-rolled Mg–Zn–Gd–Ca–Mn alloys, Int. J. Miner. Metall. Mater., 31(2024), No. 10, pp.2208-2220. https://dx.doi.org/10.1007/s12613-023-2812-5
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Magnesium and its alloys are the lightest structural metals. They have high specific strength, high specific stiffness, satisfactory electromagnetic shielding capability, good damping properties, and good machinability, which have broad application potential in the aerospace, rail transportation, electronic information, and automotive industries [1–6]. However, magnesium, as a hexagonal close-packed structure, has few independent slip systems [7–8]. Albeit there are various deformation modes such as basal slip, nonbasal slip, tension twinning, and compression twinning, initiating these deformation modes simultaneously is difficult due to the large difference in critical resolved shear stress (CRSS) [9–11], which cannot satisfy the five independent slip systems required for homogeneous deformation based on the von Mises criterion, resulting in poor ductility and formability. Even the commercial wrought magnesium alloy Mg–3Al–1Zn (wt%, AZ31), which has the best plasticity, still cannot overcome these problems [12–17], which largely hinders the application of magnesium alloys in several fields.
At present, two strategies, thermomechanical processing and microalloying, have been developed to address the poor plasticity of magnesium alloys [18–25]. Among them, rare earth (RE) elements, as alloying elements with high solid solution in magnesium alloys, can change the lattice constant and reduce the stacking fault energy and CRSS of nonbasal slip, thereby effectively refining the grain, weakening the texture tension, and even altering the texture type, which has drawn increasing interest from scholars at home and abroad [26–32]. Many high-plasticity magnesium alloy systems have been designed, such as Mg–RE [33–35] and Mg–Zn–RE [36–37], which have outstanding mechanical properties at room temperature and high temperature. Moreover, the microstructure of the alloy can be enhanced by subsequent annealing heat treatment, such as grain size, texture, and second phases, to adjust the mechanical properties and processing performance of the material [38]. Zhang et al. [39] observed that the Mg–3Al–0.6Ca–0.2Gd (wt%) alloy annealed at 350°C for 1 h presented good comprehensive mechanical properties and low anisotropy. Guo et al. [40] examined the effects of recrystallization and phase refinement during the annealing process on the mechanical properties of Mg–7.8Li–0.8Zn (wt%) alloy and revealed that annealing treatment greatly improves the elongation of rolled plates, with the highest elongation reaching 34.6% and relatively high strength. Basu et al. [41] found that annealing at different temperatures for 60 min can change the microstructure of hot-rolled Mg–1Gd (wt%) and Mg–1Dy (wt%) alloys, thereby improving the mechanical properties of magnesium alloys.
Previously, we designed a series of Mg–Zn–Gd high-plasticity magnesium alloys through Gd–RE microalloying, which exhibited good formability at room temperature due to their large elongation, high strain hardening rate, and low anisotropy [42–46]. Subsequently, we improved the strength while maintaining the outstanding plasticity of these alloys by introducing a reinforcing phase with the addition of Ca. However, the recrystallized grain size, texture, and strength–plasticity matching of the alloy are sensitive to rolling and annealing processes because of the relatively low degree of alloying.
Based on this, we designed a Mg–Zn–Gd–Ca–Mn alloy via microalloying and employed the warm rolling process to obtain sufficient residual stress for static recrystallization, with the aim of obtaining a fine microstructure and good balance of mechanical properties. The effect of the annealing process on the microstructure, texture, and mechanical properties of the warm-rolled sheet was investigated to provide a theoretical basis and reference for the industrial production process of these alloys.
A Mg–1.8Zn–0.8Gd–0.1Ca–0.2Mn (wt%) alloy ingot was prepared by melting pure magnesium, zinc, gadolinium, manganese, and Mg–30wt%Ca master alloy in a heat-resistant mild steel crucible. Chemical compositions were determined using inductively coupled plasma atomic emission spectroscopy, and the results are summarized in Table 1. Slabs with dimensions of 400 mm × 200 mm × 20 mm were cut from the ingot, homogenized at 400°C for 10 h, and then cooled in air. The slabs were heated to 320°C for 20 min, rolled to a final thickness of 3 mm with a reduction of 15%–20% per pass along the same direction, and reheated at 320°C for 15 min after each pass. The samples were cut to appropriate sizes using a wire-cut electrical discharge machine and annealed at 250, 300, 350, 375, 400, 425, 450, 480, or 500°C for 1 h. The samples were cooled to room temperature in air.
| Zn | Gd | Ca | Mn | Mg |
| 1.77 | 0.75 | 0.12 | 0.21 | Bal. |
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The microstructure was observed at the rolling direction–transverse direction (RD–TD) and rolling direction–normal direction (RD–ND) planes by optical microscopy (OM), scanning electron microscopy (SEM) with an energy-dispersive spectrometer (EDS), transmission electron microscopy (TEM), and electron backscatter diffraction (EBSD). EBSD characterization was conducted using a Hitachi Se3400 SEM instrument at 20 kV. Orientation imaging microscopy was performed at a step size of 0.9 μm, and the data were analyzed with TSL-OIM Analysis software. Macrotexture was determined at the RD–TD plane using a Rigaku D/max-2400 X-ray diffraction (XRD) machine with single Cu Kα radiation and a measurement angle of 0°–70°. XRD data were corrected with Mg powder and processed with DIFFRAC plus TEXEVAL software. TEM was conducted using a ThermoFisher TalosF200X G2 instrument. For OM and SEM observations, polished specimens were etched with a mixture of 2 g picric acid, 12 mL acetic acid, 10 mL ethanol, and 30 mL distilled water. The average grain size of the alloys was calculated using the equivalent spherical diameter method. Samples for texture analysis were prepared by mechanical grinding and electropolishing with an AC2 electrolyte solution.
Tensile specimens with a gauge length of 30 mm, width of 7 mm, and thickness of 3 mm were cut from the sheet along the RD and TD. Tensile properties were analyzed using an MTS E45.305 universal testing machine at a constant speed of 1.5 mm·min–1. Three specimens were tested for each sample.
The microstructures of the Mg–Zn–Gd–Ca–Mn alloy in the rolled and annealed states are displayed in Figs. 1 and 2, respectively. Fig. 1 presents the microstructure of the alloy in the rolled state and recrystallization stage. From Fig. 1(a), the grain size of the rolled sheet is inhomogeneous, with an average grain size of approximately 35 μm. Because of the relatively low rolling temperature, most of the grains have twins, which are irregularly distributed and orientated at an angle to the RD and cross each other. At an annealing temperature of 250°C (Fig. 1(b)), static recrystallization starts to take place locally, forming a fine recrystallized grain structure with a recrystallization volume fraction of approximately 40% and an average grain size of approximately 18 μm. However, there are still some uncrystallized grains with a few twins inside. With the annealing temperature increasing to 300°C (Fig. 1(c)), the twins within the grains basically vanish, and the recrystallization ratio increases to 80%–90%, with an average grain size of approximately 10 μm. After annealing at 350°C (Fig. 1(d)), complete recrystallization occurs, and all the recrystallized grains are equiaxed, with an average grain size of approximately 15 μm. Fig. 2 displays the microstructure of the recrystallized grain growth stage of the alloy at annealing temperatures of 375–500°C. With the continuous increase in annealing temperature, the recrystallized grain size grows normally and gradually increases from 25 μm at 375°C to 42 μm at 425°C, as shown in Fig. 2(a)–(c). At temperatures above 450°C, as shown in Fig. 2(d)–(f), the grains grow rapidly, and the average grain sizes are 165 μm at 450°C, 229 μm at 480°C, and 298 μm at 500°C.
Fig. 3 presents the grain size and recrystallization ratio statistics with increasing temperature. For Mg–Zn–Gd–Ca–Mn alloy, the annealing temperatures can be divided into three ranges based on its static recrystallization behavior and grain size. The recrystallization begins at 250°C, and the recrystallization ratio gradually increases with increasing temperature. Complete recrystallization and grain normal growth occur at 350°C, and the grains are homogenous and fine. The grain growth range can be further divided into normal growth from 375°C, where the grains grow slowly, and rapid growth, where the grains increase rapidly and significantly from 42 μm at 425°C to 165 μm above 450°C. Below the temperature of 425°C, the Mg–Zn–Gd–Ca–Mn alloy recrystallizes at a relatively slow rate of grain size growth compared with the Mg–Zn–Gd alloy, but above 425°C, the recrystallized grains grow significantly. This phenomenon may be associated with the type and concentration of the second phase and the solute atoms due to the addition of Ca and Mn. The grain size growth curve was fitted with Origin software to obtain the recrystallized grain growth formula:
| L=−297.87/[1+(T/454.38)25]+311.45 | (1) |
where T refers to the annealing temperature, and L refers to the recrystallized grain size.
Fig. 4 exhibits the SEM images of the rolled state and the morphology of the second phases in the Mg–Zn–Gd–Ca–Mn alloy. From Fig. 4(a), some large bulk and long string phases are distributed along the RD in the rolled state alloy. Among them, the bulk phases can be divided into angular cubes (10–15 μm, Fig. 4(b)) and some irregular blocks (4–12 μm, Fig. 4(c)). The long string phase is composed of particle phases of different sizes (Fig. 4(d)). Furthermore, many spherical and rodlike particle phases are present inside the grains (Fig. 4(e)). The morphologies of the spherical and rodlike phases and their corresponding electron diffraction patterns are shown in Figs. 4(f) and 5. The diameter of the spherical phase is approximately 150–200 nm, the diameter of the rodlike phase is 200 nm, and the length is approximately 600–700 nm, where the number ratio of the spherical phase to rodlike phase is approximately 7:3, and the two phases seem to be the same type phase from the SAED patterns. From the EDS analysis results (Table 2), the regular cube phase is the Mg(Gd,Ca) binary phase, the irregular block phase comprises Zn and Gd, which is Mg3Gd2Zn3 (W phase), and Ca is also identified in some of the irregular block phases, which may result from the coexisting Ca2Mg6Zn3 and W phase; the long string phase comprises Zn, Gd, and Ca, which is the MgZn(Gd,Ca) ternary phase; the spherical and rodlike particle phases contain Zn and Gd with a Zn/Gd atomic ratio of 2.4, and they are determined as the W phase, which is consistent with Zhao et al. [47].
| Phase | Mg | Zn | Gd | Ca |
| Block phase | 37.63 | 40.78 | 21.59 | 0 |
| 51.98 | 31.44 | 16.46 | 0.12 | |
| Cube phase | 8.14 | 0 | 79.93 | 11.93 |
| Long string phase | 47.83 | 32.78 | 19.11 | 0.29 |
| Spherical phase | 87.96 | 8.56 | 3.58 | 0 |
| Rodlike phase | 93.06 | 4.95 | 1.98 | 0 |
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Figs. 6–8 show the changes in the morphology and quantity of the second phases during the annealing process. Fig. 6 shows the SEM images of the second phase in the recrystallization stage. After annealing at 250°C for 1 h, the size of the long string phase is decreased compared with that of the rolled state, while there is no significant change in the block phase and the fine particle phase within the grains, as shown in Fig. 6(a) and (b). After annealing at 300°C for 1 h, in Fig. 6(c) and (d), the number and size of the long string phase are further reduced, and the size of the block phase and the number of particle phases are also reduced. Among them, the number of rodlike phases in grains accounts for approximately 20% and the number of spherical phases accounts for approximately 80%. After annealing at 350°C for 1 h (Fig. 6(e) and (f)), the number and size of all the long string and block phases are further decreased, and the number of particle phases inside the grain is also further decreased, where the number of rodlike phases accounts for approximately 10%, and the number of spherical phases accounts for approximately 90%.
Fig. 7 shows the SEM of the second phase evolution at the normal growth stage of recrystallized grains. From Fig. 7(a) and (b), after annealing at 375°C, the number and size of the block and long string phases are further decreased. The rodlike particle phase inside the grains almost vanishes, while the spherical phase is reduced slightly. After 400°C (Fig. 7(c) and (d)), all the block, long string, and particle phases inside the grains are further reduced, and the number of spherical phases decreases slightly. After annealing at 425°C, almost half of the spherical phase vanishes, as presented in Fig. 7(e) and (f).
Fig. 8 exhibits the second phase evolution during the rapid grain growth stage. After annealing at 450°C, the smaller-sized phases in the long string phase vanishe, and thus, the long string phase becomes discontinuous, and the spherical phase within the grain entirely disappears. On the other hand, a few cube phases start to form, as shown in Fig. 8(a) and (b). After annealing at 480°C, the long string phase basically disappears, as shown in Fig. 8(c) and (d); the number and size of the cube phase are further increased after annealing at 500°C, as shown in Fig. 8(e) and (f).
Based on the above analysis, the second phases in the rolled state of the alloy are primarily large-sized block and long string phases distributed along the RD and spherical and rodlike particle phases within the grains. The large-sized block and string phases are the MgZnGd ternary phase, which was essentially the unresolved eutectic phases in the as-cast alloy that were broken during heat treatment and rolling. The spherical and rodlike particle phases within the grains are possibly precipitated during rolling deformation. The evolution of the morphology, size, and quantity of these phases throughout the annealing process are thought to be associated with the annealing temperature and phase type. Overall, it can be divided into three stages. (1) In the recrystallization stage of 250–350°C, the number of block, long string, and spherical phases slightly decreases, while the rodlike phase decrease significantly, probably because the gradual solid solution of Zn and Gd into the matrix makes the rodlike phase morphology tend to become a spherical phase. In this temperature range, the grain size remains mostly the same due to the presence of several second phases of different sizes around the grain boundaries and within the grains, which hinders the migration of grain boundaries. (2) At 375–425°C, the number of block and long string phases further decreases, the rodlike phase disappears, and the fine MgZnGd ternary spherical phase within grains decreases sharply, even to the extent that some internal grains are free of spherical phases. The grains grow significantly due to the decrease in the second phase. (3) At 450–500°C, the spherical particle phase vanishes entirely, and there is no hindrance to grain boundary migration, which leads to rapid grain growth. Meanwhile, a few large regular cube phases, which are Mg(Gd,Ca) phases formed during high-temperature heating, increase. In conclusion, the pattern of grain size change of the present alloy during annealing is ascribed to the stability of the second phase, especially the fine phase within the grains. Therefore, a suitable annealing temperature for this alloy is below 350°C to ensure a fine microstructure.
Fig. 9 displays the macrotexture of the Mg–Zn–Gd–Ca–Mn alloy sheet. Fig. 9(a) exhibits that the rolled state alloy has a circular nonbasal texture comprising two strong TD-split texture components with a peak texture strength of 2.1 and two weak RD-split texture components, similar to that of the Mg–Zn–Gd–Ca alloy by Yan et al. [44] and Wei et al. [46]. Therefore, the addition of Ca and Mn did not influence or obliterate the TD-split texture-weakening effect of Gd. Similarly, the addition of Mn did not influence the formation of the RD-split texture components of Ca. Besides RE elements, it has been demonstrated that alloying elements such as Ca and Mn can also play a role in texture weakening. The texture weakening mechanism is associated with the particle-stimulated recrystallization nucleation mechanism induced by second-phase particles containing Ca or Mn and the reduction of the lattice constant [15,48–49].
After annealing at 250°C (Fig. 9(b)), the texture type is roughly the same as that in the rolled state, and the texture strength does not change much. After annealing at 300°C (Fig. 9(c)), the intensity of the two TD-split texture components increases, whereas the −RD-split texture component weakens and the +RD-split texture component vanishes. At this time, the texture changes from a circular-shaped texture to a double-peaked texture typical of Mg–Zn–RE. After annealing at 350°C (Fig. 9(d)), the +RD-split texture component also disappears, and the intensity of the TD-split texture slightly weakens, with a peak intensity of 1.96. Generally, with increasing annealing temperature, static recrystallization transforms the RD texture component to the TD texture component. In the recrystallized grain growth stage, the angle of the TD texture component increases, and the peak texture intensity decreases.
The inverse pole figure (IPF) map, grain size distribution, and microscopic pole figure of the Mg–Zn–Gd–Ca–Mn alloy are illustrated in Fig. 10. In Fig. 10(a) and (d), the rolled state alloy comprises coarse deformed grains and several fine dynamically recrystallized grains of 2–5 μm, with an average grain size (Dave) of 18.85 μm. In addition, there are a large number of low-angle grain boundaries (LAGBs) within the grains and high internal energy storage (Fig. 10(g)). When annealing at 300°C, static recrystallization takes place, and the large deformed grains disappear. A large number of statically recrystallized grains appear with an overall average grain size of 11.13 μm, whose size is slightly larger than that of the dynamically recrystallized grains by 6–8 μm. The recrystallized grains have a more random grain orientation, and some LAGBs still exist. The overall energy storage is significantly reduced because of static recrystallization, with only a single grain showing high energy storage (Fig. 10(h)).
The microscopic pole diagrams of the alloy in different states are exhibited in Fig. 10(j)–(l). The intensity of the texture gradually weakens from 6.838 in the rolled state to 3.807 after annealing at 350°C with increasing annealing temperature. The microtexture is gradually discrete, and the texture evolution pattern is similar to that of the macroscopic texture. However, its intensity is significantly higher than that of the macroscopic texture, which may be associated with the fact that the scanning area of EBSD is much smaller than that of XRD.
The misorientation distributions of the Mg–Zn–Gd–Ca–Mn alloy are illustrated in Fig. 11(a) and (b). In Fig. 11(a), the misorientation angle of the rolled alloy is primarily distributed in the range of 2°–12.98° (<15° LAGB), accounting for up to 60% or more. After annealing at 300°C for 1 h (Fig. 11(b)), the LAGBs decrease rapidly, accounting for approximately 10%, and the high-angle grain boundaries increase to a proportion of 90%. Fig. 11(c) and (d) demonstrates the twin distributions of the Mg–Zn–Gd–Ca–Mn alloy in the rolled and annealed states at 300°C. The common twin modes in magnesium alloys are the {10ˉ12} tension twins, {10ˉ11} compression twins, and {10ˉ11}–{10ˉ12} double twins. From Fig. 11(c), there are many {10ˉ12}–{10ˉ11} double twins, a few {10ˉ12} tension twins, and very few {10ˉ11} compression twins within grossly deformed grains in the rolled state. After annealing at 300°C for 1 h (Fig. 11(d)), most of the twins disappear, and only a few {10ˉ12} tension twins and {10ˉ12}–{10ˉ11} double twins exist within individual grains.
Because of the relatively low rolling temperature, dynamic recrystallization was insufficient. Therefore, a predominantly deformed microstructure with high distortion energy, which can cause and affect static recrystallization during annealing, was obtained. The microstructural characteristics of the rolled state were studied. The EBSD microstructure of the rolled-state alloy is presented in Fig. 12. The IPF map (Fig. 12(a)) shows that the structure is irregularly shaped deformed grains with many LAGBs and twins within them. The twins (Fig. 12(b)) are mainly {10ˉ12}–{10ˉ11} double twins and a few {10ˉ12} tension and {10ˉ11} compression twins. The double twins maintain their lenticular shape under heat and stress during the rolling process but are split into several pieces because of the interaction with dislocations, while the tension twins grow into an irregular shape. It is clear that dynamic recrystallization occurs mainly along the {10ˉ12}−{10ˉ11} double twin and grain boundary-shaped nuclei. In Fig. 12(d), the crystallographic orientation on both sides of the recrystallization is essentially the same, which belongs to the same grain, and the type of recrystallization can be concluded as recrystallization along the twin crystal form nuclei. Similarly, a small amount of dynamic recrystallization is also found in the lamellar {10ˉ12} tension twins, as observed in Fig. 12(e). Also, numerous small dynamically recrystallized grains are observed near the grain boundaries.
The shear displacement of twins is much smaller than the amount of slip deformation and thus can easily occur. Also, the twins can change grain orientation so that the unfavorable crystallographic orientation for the slip becomes more favorable, thus coordinating plastic deformation. With gradually increasing rolling deformation, the dislocation slip starts and interacts with more stable double twin and grain boundaries, forming many LAGBs. The boundaries of grains and twins with high-density dislocations become the potential core of dynamic recrystallization and induce dynamic recrystallization [50–51]. The high kernel average misorientation (KAM) around grain boundaries and intragranular twins and the accumulated subgrain boundaries, as shown in Fig. 12(c), indicate that they can offer sufficient strain gradients to prompt static recrystallization. As marked by white arrows in the figure, some dynamically recrystallized grains appear near the region with high KAM values, implying that this region of grain boundaries and compression twin boundaries can also be a potential nucleation location for static recrystallization.
The room temperature tensile stress–strain curves along the RD and TD are plotted in Fig. 13, and the corresponding mechanical properties are summarized in Table 3. It can be observed that the rolled alloy shows high yield strength (YS > 210 MPa) and ultimate tensile strength (UTS > 315 MPa) but has a lower elongation and strain-hardening exponent n. This is mainly due to the large deformation and high dislocation density during the rolling process coupled with the presence of several phases that produce dislocation plugging during further deformation and a stronger work-hardening phenomenon [52–53], requiring greater force to be applied to deform the alloy.
| State | Direction | YS / MPa | UTS / MPa | Elongation / % | n |
| Rolled | RD | 213.3 | 315.9 | 6.22 | 0.04 |
| TD | 238.6 | 329.4 | 19.3 | 0.16 | |
| Annealed at 250°C | RD | 189.5 | 286.1 | 12.2 | 0.08 |
| TD | 230.3 | 310.5 | 23.2 | 0.19 | |
| Annealed at 300°C | RD | 182.1 | 271.1 | 27.4 | 0.21 |
| TD | 176.9 | 275.8 | 32.3 | 0.30 | |
| Annealed at 350°C | RD | 169.0 | 275.2 | 25.2 | 0.25 |
| TD | 155.4 | 273.8 | 31.2 | 0.33 |
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At an annealing temperature of 250°C, the YS and UTS decrease slightly while the elongation and n increase. At 300°C, the YS and UTS decrease significantly, and the elongation and n increase with great intensity. When the temperature increases up to 350°C, the elongation change is not significant, and the YS is slightly decreased because the average grain size is slightly increased compared to 300°C due to grain growth. When the temperature is higher than 375°C, rapid grain growth makes the grain size inhomogeneous, and thus, the strength and plasticity decrease [54].
Note that the alloy sheet has clearer anisotropy along the TD and RD. With increasing annealing temperature, the difference in elongation decreases, whereas the difference in YS changes little. The alloy annealed at 300°C has optimal plasticity with elongation values of 27.4% and 32.3% along the RD and TD, respectively. Meanwhile, the YS is higher than 170 MPa, and the UTS is >270 MPa, with a relatively good match of strength and plasticity. The effect of annealing temperature on strength, plasticity, and anisotropy is associated with the effect of annealing on the degree of recrystallization, grain size, texture, second phase, and fracture mode. This will be discussed in a comprehensive manner at the end.
The secondary electron (SE) and backscattered electron (BSE) SEM images of the tensile fracture of the Mg–Zn–Gd–Ca–Mn alloys are displayed in Fig. 14. In Fig. 14(a) and (d), both TD and RD direction fractures of the 300°C annealed state comprise many dimples and some tearing ridges, and all other annealing temperature fracture morphologies are similar, implying a ductile fracture in the annealed alloy. In contrast, several long string and block phases exist in the fracture morphology of the rolled state, as observed in Fig. 14(b) and (e), indicating a brittle fracture. With increasing annealing temperature, the number and size of the long string phase in the fracture decrease at 300°C (Fig. 14(c) and (f)), and the block phase is observed at the bottom of the dimples, with part of the second phase torn and many particle phase distributions observed in the fractures. The EDS analysis reveals that the long string phase and the bright white block phases A and C are the MgZnGd ternary phase, and some dark block phases B and D are the MgZnCa ternary phase. Also, the number of second phases is higher in the RD fracture than in the TD fracture.
In conclusion, the annealing process, especially the annealing temperature, significantly influences the grain size and the corresponding mechanical properties of the Mg–Zn–Gd–Ca–Mn warm-rolled sheets. The influence of the annealing temperature on the microstructure is mainly reflected in two aspects. (1) The annealing temperature greatly impacts the recrystallization fraction, grain size, and texture. Static recrystallization starts at 250°C with finer recrystallized grains and a basal deformation texture, whereas complete recrystallization occurs at 350°C with fine uniform grains and a nonbasal texture, and grain coarsening becomes apparent at 450°C, resulting in large grains. (2) The annealing temperature obviously affects the fraction and morphology of the W phase. The W phase is relatively stable in the temperature range of 250–350°C. However, the W phase gradually dissolves in the matrix at temperatures above 375°C and disappears at 450°C.
The microstructural evolution is driven by the storage energy from the warm-rolling process. At lower temperatures, partial recrystallization occurs in the stress-concentrated region, such as grain boundaries, which accumulate a high density of dislocations, and the recrystallized grains are comparatively small. At higher temperatures, the dislocations climb, and solute-atom diffusion accelerates, thus quickening the recrystallization process. On the other hand, grain boundary migration becomes active not only because of the thermal input but also because of the absence of the hindering effect of the W phase on grain boundaries at higher temperatures [55–56]. As for the texture evolution mechanism of the Mg–Zn–Gd alloy during recrystallization, the orientated nuclei and/or orientated grain growth are responsible for the transition from basal to nonbasal texture. However, this remains unclear.
From the above analysis, the rolled and 250°C annealed alloys have a large number of deformed grains, few recrystallized grains, and no apparent change in texture type and strength, and there are several block and long string phases (W phase) in the alloys. At this time, the alloy has higher strength and lower plasticity due to the high residual stress and large block W phase, which are prone to cracking. At 300–350°C, the finer uniform grains and the double-peaked texture are the benefits of plasticity improvement as the grain boundaries can accommodate deformation, and soft orientation has a high Schmid factor for basal slip and extension twining. Moreover, as Zn and Gd gradually dissolve in the matrix, the quantity of block, long string, and spherical phases slightly decreases, which decreases the stress concentration and cracking of the second phase. Thus, the elongation of the alloy along the RD and TD directions reaches maxima of 27.4% and 32.3%, respectively. At 375–425°C, the number of all phases inside the grains sharply reduces, and the grains noticeably grow larger. Hereafter, both the strength and plasticity of the alloy decrease based on the Hall–Petch relationship. Generally, the strength and elongation of the alloy are associated with the grain size, the presence of particles, and the grain orientation based on the Hall–Petch relationship, the Schmid factor for deformation modes, and the crack nuclei and propagation theory.
Therefore, for Mg–Zn–Gd–Ca–Mn alloys with low alloying elements, the sensitivity of the above microstructure to annealing temperature requires precise control of the annealing temperature to obtain the desired remarkable mechanical properties.
A high-plasticity Mg–1.8Zn–0.8Gd–0.1Ca–0.2Mn alloy was developed by multicomponent microalloying, and the influence of the warm-rolled and annealing processes on the grain size, second phase, texture, and room temperature mechanical properties of the sheet was analyzed. The main conclusions can be drawn.
(1) For Mg–1.8Zn–0.8Gd–0.1Ca–0.2Mn alloy by warm rolling at 320°C, the deformation microstructure comprises some fine dynamic recrystallized grains and large grains with several LAGBs, double twins, and compression twins. It is accompanied by some larger-sized block phases (W phase) and long strings (W phase) along RD that remain from the casting process, as well as spherical and rodlike particle phases within the grains (MgZnGd ternary phase).
(2) The annealing temperature significantly influences the recrystallization degree, second phase, and texture of the warm-rolled Mg–1.8Zn–0.8Gd–0.1Ca–0.2Mn alloy. At the annealing temperature of 250°C, there is only a small amount of static recrystallization, with slight changes in mechanical properties. At 300–350°C, completely static recrystallization occurs, with uniform and fine grains and a slightly reduced number of second phases, demonstrating a typical bimodal RE texture. At the annealing temperature of 375–450°C, the long string, spherical, and rodlike particle phases within the grains are reduced, the grains start to grow. Above 450°C, the spherical phases within the grains all vanish, and the grains undergo rapid growth.
(3) The annealing process significantly affects the mechanical properties of warm-rolled alloy. The rolled alloy has high strength, with YS of >210 MPa and UTS of >315 MPa. Under the condition of incomplete recrystallization annealing at 250°C, the strength decreases slightly and the plasticity increases. In the stage of complete recrystallization at 300–350°C, the optimum strength–plasticity comprehensive mechanical properties are present, with YS of 182.1 and 176.9 MPa, UTS of 271.1 and 275.8 MPa, and elongation values of 27.4% and 32.3% in the RD and TD, respectively. Therefore, the strength and plasticity of this material can be regulated by annealing based on the demand. Meanwhile, the annealing temperature should also be precisely controlled during production to avoid fluctuations in the properties. Furthermore, there are still some larger-sized phases in this alloy that affect its mechanical properties, and this material has room for improvement.
This work was financially supported by the National Natural Science Foundation of China (Nos. 52271107 and 52205392), the Natural Science Foundation of Shandong Province (No. ZR2021ME241), and the Bintech-IMR R&D Program (No. GYY-JSBU-2022-012).
All authors declare that they have no financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Figures(14) / Tables(3)
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| Zn | Gd | Ca | Mn | Mg |
| 1.77 | 0.75 | 0.12 | 0.21 | Bal. |
DownLoad:
CSV
| Phase | Mg | Zn | Gd | Ca |
| Block phase | 37.63 | 40.78 | 21.59 | 0 |
| 51.98 | 31.44 | 16.46 | 0.12 | |
| Cube phase | 8.14 | 0 | 79.93 | 11.93 |
| Long string phase | 47.83 | 32.78 | 19.11 | 0.29 |
| Spherical phase | 87.96 | 8.56 | 3.58 | 0 |
| Rodlike phase | 93.06 | 4.95 | 1.98 | 0 |
DownLoad:
CSV
| State | Direction | YS / MPa | UTS / MPa | Elongation / % | n |
| Rolled | RD | 213.3 | 315.9 | 6.22 | 0.04 |
| TD | 238.6 | 329.4 | 19.3 | 0.16 | |
| Annealed at 250°C | RD | 189.5 | 286.1 | 12.2 | 0.08 |
| TD | 230.3 | 310.5 | 23.2 | 0.19 | |
| Annealed at 300°C | RD | 182.1 | 271.1 | 27.4 | 0.21 |
| TD | 176.9 | 275.8 | 32.3 | 0.30 | |
| Annealed at 350°C | RD | 169.0 | 275.2 | 25.2 | 0.25 |
| TD | 155.4 | 273.8 | 31.2 | 0.33 |
DownLoad:
CSV
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