Influence of a novel SPD technique together with heat treatment on the microstructural characteristics and hardness of Mg-13Gd-4Y-2Zn-0.5Zr alloys

To illuminate the influence of CEE-AEC on the microstructure of Mg-13Gd-4Y-2Zn-0.5Zr alloy, OM and SEM micrographs of the as-cast and 3 CEE-AEC passes samples were displayed in figures 3(a)–(d), respectively. The TEM images and corresponding EDS analyses of the phases in CEE-AEC samples were displayed in figure 4. In the as-cast alloy, lots of fine lamellar phases precipitated from the grain boundaries to the interiors. The network-like eutectic phases could be observed in the SEM micrograph, which almost precipitated on the interdendritic block phases. After 3 CEE-AEC passes, the interdendritic block phases were elongated along with the extrusion direction, and some of them were broken down into much finer pieces. A large majority of the broken block phases distributed around the the grain boundaries, which pinned the grain boundaries and inhibited the growth of grains [16]. According to the corresponding selected area electron diffraction (SAED) pattern exhibited in figure 4(b), there were thirteen additional diffraction spots existing at the positions between the (0000)_ɑ-Mg and (0002)_ɑ-Mg reflections. It proved the interdendritic block phases had a 14H-type structure. What's more, the network-like eutectic phases were also shattered into the granular-like phases during CEE-AEC process. And the shattering eutectic phases still distributed around the elongated block phases.

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Compared to the network-like phases, the distribution of the granular-like phases was more dispersive, which restrained the movement of dislocations and strengthen the matrix [17, 18]. Referring to the corresponding SAED pattern (B//[001]) and EDS analysis exhibited in figures 4(c) and (e), the white granular-like phases could be ascertain to be the β-Mg₅(Gd, Y, Zn) phases, which had a face centered cubic (fcc) structure. There were some cuboid-shaped phases in the microstruture of CEE-AEC samples. The corresponding SAED pattern (B//[011]) and EDS analysis exhibited in figures 4(d), (f) indexed that the cuboid-shaped phases were RE-riched phases which had a FCC structure and a = 0.568 nm lattice constant. Apparently, with shear strains accumulating in the muti-pass CEE-AEC process, the grain structure was greatly refined. The average grain size was measured as 4.6 μm and 3.0 μm after one and three CEE-AEC passes respectively by linear intercept method. With reference to the study results of Yan et al [8], the grain size was refined to 12.1 μm and 1.4 μm in the asymmetric cavity with 5 mm differential after one and three CEE-AEC passes respectively, which meant increasing the differential speed of asymmetric cavity (from 5 mm to 10 mm) couldn't refine the grain structure further after 3 CEE-AEC passes but could produce better refinement effects after one CEE-AEC pass. What's more, certain intragranular lamellar phases, 14H-type structure according to figure 4(a), exhibited different orientations in different grains. These different orientation lamellar phases could suppress the substrate dislocation slip and the lattice rotation [19, 20].

To identify the optimal solution-treated parameters of CEE-AEC samples, the hardness variation in different solution treatment time was exhibited in figure 5. As shown, the hardness of the CEE-AEC samples owned a relative high value compared to that of the as-cast one, which increased from 78 HV to 106 HV. After solution treatment, the hardness of the samples had a general decrease. The decreasing hardness could be attributed to the coarsening of the grain sizes in solution-treated samples, according to the Hall-Petch relationship [21]. When solution treatment maintained 8 h, the hardness of the alloy descended further compared to that of 4 h. The further descent of the hardness was related to the dissolution of the lamellar LPSO phases which precipitated in large quantities at 4 h, as shown in figure 6(b). The recent researches indicated that the lamellar 14H-LPSO phase plays a significant role in strengthening the alloys in Mg-Gd-Y-Zn systems [22, 23]. And Yamasaki et al [24] also found that the highly dispersed 14H-LPSO hard lamellas could bring about high strength for Mg-RE alloys. Hence, the dissolution of the lamellar phases would result in the further decrease of the alloy hardness. When the solution time extended to 12 h, the lamellar phases almost dissolved completely, as shown in figure 6(c). But the hardness of the alloy achieved the peak instead. Actually, with the solute atoms diffused at the Mg matrix, the degree of the lattice distortion would be gradually enhanced. The distortional lattice hindered the movement of dislocations, thereby resulting in the increasing hardness of the solution-treated samples [25]. Undoubtedly, the peak hardness signified the solute atoms in the Mg matrix has reached the saturation, the overlong solution-treated time would give rise to the growth of grains, which resulted in the decrease in hardness.

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Figure 6 illustrates the optical microstructure of the solution-treated samples at 520 °C with the time range from 4 h to 16 h. Obviously, the grain sizes of all samples had a general growth. After the solution treating for 4 h, the volume fraction of the elongated block phases declined. The lamellar phases precipitated in large quantities from the grain boundaries and even ran through the whole grains. Notably, the lamellar phases invariably precipitated at the vicinity of the dissolving block phases. With the block phases dissolving, free energy was released and numerous solute atoms diffused to the grain boundaries, which promoted the precipitation of the lamellar phase [26]. Thus, it could be summarized that the interdendritic block phases transformed into intragranular lamellar phases or intragranular lamellar phases precipitated at the expense of the interdendritic block phases during the solution treatment period. Liu et al [12] also found the transformation of the elongated stripe phase to lamellar phase and Zhu et al [27] attributed the phenomenon to the minimization of shear strain energy. In addition, with the time of solution treatment prolonged, the intragranular lamellar phase gradually dissolved into Mg matrix (see figure 6(b)) and almost disappeared in the solution treatment of 12 h (see figure 6(c)). To best our knowledge, the dissolution and precipitation of the second phases was a dynamic process. The volume fraction of the second phases started declining when the precipitation of the lamellar phases arrived the peak, and fianally maintained relatively stable after the solution-treated time of 12 h. Hence, combining hardness curves analysis in figure 5, the 520 °C−12 h was finally selected as the optimum solution parameter of the Mg-13Gd-4Y-2Zn-0.5Zr alloys of 3 CEE-AEC passes.

In order to investigate the influence of CEE-AEC on the solution treatment of Mg-13Gd-4Y-2Zn-0.5Zr alloys, the SEM micrographs of the 520 °C−12 h samples were exhibited in figure 7. As shown, the microstructure of the CEE-AEC samples had distinct difference with the as-cast ones. In the CEE-AEC samples, a large number of the second phases dissolved in to the Mg matrix after solution treatment of 12 h, including broken block phases, lamellar phases, granular-like phases and some elongated block phase. While in the as-cast alloys, minorities of the interdendritic block phases became discontinuous due to the dissolution, as indicated by red arrows in figure 7(b). And numerous RE-riched small particles precipitated around the dissolving block phases. Lots of lamellar phases also could be observed in the as-cast alloys after solution treating for 12h. However, the phenomenon that numerous lamellar phases precipitated from the grain boundaries has already appeared at the CEE-AEC samples in the solution treatment of 4 h, with reference to the figure 6(a). According to the studies of Abe et al, the precipitation of the lamellar phases includes two steps, one is the formation of the stacking faults (SFs) in Mg matrix, the other is the solid atoms (Gd, Y, Zn) diffuse to the SFs [28]. The faster diffusion rate of the solid atoms would contribute to the advance precipitation of the lamellar phases. Actually, the improvement of the diffusion rate was related to the multiplicative dislocations caused by CEE-AEC. The local strain fields occurred around the dislocations provided diffusion kinetics for the solid atoms and multiplicative dislocations provided the sufficient channels for solute atoms diffusing. The factors pointed above facilitated the advance precipitation of the lamellar phases. What's more, compared the micrograph of two samples, the volume fraction of the second phases was lower in the CEE-AEC samples. It indicated that the Mg matrix of CEE-AEC samples owned the higher element concentration compared to the as-cast ones. Higher element concentration in the Mg matrix was favorable for the precipitation of the β phases in subsequent ageing treatment.

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Figure 8 shows the hardness comparison curves of the alloys which aged at 225 °C for different time. The as-received samples included 3 CEE-AEC samples and as-cast alloys, respectively. As shown, the samples of CEE-AEC have exhibited distinct ageing hardening response in a short treating time (4 h), which obtained a relative high hardness value of 111 HV. While the hardness of the as-cast ones didn't arise apparently until 8 h. As the ageing time increased to 16 h, the ageing hardening response of the CEE-AEC samples reached the peak with the value of 128 HV. Compared to the CEE-AEC samples, the hardness of the as-cast samples arrived at the peak after ageing treatment of 64 h, with the value of 121 HV. The value of peak hardness and corresponding ageing treatment time of the CEE-AEC samples were 5% higher and 75% shorter than those of the as-cast ones, respectively. From the overall trend of two curves, the hardness curve of CEE-AEC samples arose rapidly from the initial stage to the peak stage of the ageing treatment, and remained relatively stable for a long time afterwards. While the tendency of the as-cast samples increased more slowly until the hardness reached the peak and declined more quickly in the subsequently long-term ageing treatment, which was contrary to the hardness variation tendency of the CEE-AEC ones. The rapid increment of the hardness in CEE-AEC samples may be related to the higher element concentrations in the Mg matrix. Higher element concentrations provided more nucleation kinetics for the precipitation of the β' phases, which contributed to a sharp increase in the hardness of the CEE-AEC samples.

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In order to investigated the microstructure evolution and ageing precipitation sequence of the samples of 3 CEE-AEC passes, TEM and corresponding SAED were conducted. Figures 9(a) and (b) show the morphology of the samples aged at 225 °C for 4 h. It's apparent to see a large number of fine spherical precipitates with clear strain contrast, indicated by the red circle in figure 9(a), distributed heterogeneously in the Mg matrix. The uneven distribution of the spherical precipitates were related to the density of the dislocations. Actually, the second phases precipitated preferentially in the areas which were filled with large amounts of dislocations. The dense dislocations caused by CEE-AEC provided lots of favorable nucleation positions, which resulted in a concentrated distribution of these spherical precipitates. Referring to the corresponding SAED pattern (B//[1–100]_α) exhibited in figure 9(c), additional diffraction intensity emerging at the positions of 1/2 (11–20)_ɑ-Mg reflections. Such a microdiffraction feature demonstrated that the crystal structure of fine spherical precipitates was in accord with the hexagonal D0₁₉. These precipitates with D0₁₉ structure were named β'' phases in Mg-RE systems, they had the lattice constant a = 2 × a_ɑ-Mg ≈ 0.64 nm, c = c_ɑ-Mg ≈ 0.52 nm. The orientation(relationship between the β'' phase and a-Mg matrix was [0001]_β'' //[0001]_ɑ-Mg, [11–20]_β'' // [11–20]_ɑ-Mg [29, 30]. Compared with initial β'' phase, these spherical precipitates were more fine, which meant the beginning of the transformation from spherical precipitates (β'') to lens-shaped precipitates (β'). But β'' phase still precipitated predominantly during the early-aged period. As illustrated in figure 9(b), there were grain boundary precipitates (GBPs) accompanied with the clear precipitate free zone (PFZ) forming in the samples which aged at 225 °C for 4h. With reference to the corresponding SAED pattern (B//[111]) in figure 9(b), the GBPs were identified to own a fcc structure which were consistent with the equilibrium β phases (Mg5RE, fcc structure, ${\rm{F}}\overline{4}3{\rm{m}},$ a = 2.23 nm [31]). The mean size of the GBPs and width of PFZ were about 278 nm and 0.62 μm, respectively. Such coarse GBPs and wide PFZ were generally observed in the Mg-Gd-Y-Zn-Zr alloys aged at 225 °C for about 256 h. To the best of our knowledge, the formation of the PFZ was resulted by the concentration reducing of solute atoms and vacancies near the grain boundaries. In the CEE-AEC processed alloy, large amounts of dislocations gathered in the grain boundaries with the strains accumulating. The local strain fields occurring around these dislocations would induce solute atoms and vacancies diffusing to the grain boundaries spontaneously, which not only remarkably accelerated the precipitation and formation of GBPs and PFZ but promoted them coarsening and widening rapidly.

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After ageing treating for 16 h, the CEE-AEC samples arrived at the peak-aged stage. the morphology of the samples were indicated in figures 9(d) and (e). The SAED patterns shown in lower left corner of figures 9(f) and (i), incident electron beam paralleled to [0001]_ɑ, overlapped three symmetric orientation diffraction patterns (0°, 120°, 240°) of β' precipitates, as shown in figure 10 [27]. Actually, figure 9(f) exhibited the 240° orientation diffraction pattern in [0001]_α zone, while figure 9(i) exhibited the 0° orientation diffraction pattern in [0001]_α zone. As shown in figure 9(d), spherical precipitates existing in the early-aged period have completely transformed into lens-shaped phases with the ageing time prolonged to 16 h. The lens-shaped phases exhibited a triangular arrangement structure at the condition that the incident beam paralleled to [0001]_ɑ, which was the symbol that the ageing hardening response arrived the peak [30, 32]. In the corresponding SAED pattern recorded from the [0001]_α zone axis, additional diffraction intensity presented to the positions of 1/4 (−1010)_ɑ-Mg, 2/4 (−1010)_ɑ-Mg and 3/4 (−1010)_ɑ-Mg reflections. Such a microdiffraction feature demonstrated the precipitation of β' phase. The β' phase had a base centered orthohombic (bco) structure with the lattice parameters a = 2 × a_ɑ-Mg ≈ 0.64 nm, b = 8 × d(−1010)_ɑ-Mg ≈ 2.22 nm and c = c_ɑ-Mg ≈ 0.52 nm. The orientation relationship between the β' phase and a-Mg matrix was [001]_β' //[0001]_ɑ, (100)_β' //(2–1–10)_ɑ [33–35]. What's more, there were some coarse plate-shaped phases precipitating on the {1–100} planes and interlacing at a certain angle. The crystal structure of the phases, fcc structure with the lattice constant a = 2.22 nm, was identified by the microdiffraction pattern in figure 9(e). Hence, the coarse plate-shaped phases could be deduced to be the equilibrium β phases, which generally precipitated in the over-aged period of as-cast Mg-RE alloys. The advance precipitation of the β phases may be related to the severe plastic deformation. Actually, the deformation storage energy reserved in the process of CEE-AEC provided adequate nucleation kinetics for the formation of β phases, and the dense dislocations in Mg matrix provided the favorable positions for the nucleation of β phases. Both of them contributed to the accelerating precipitation of the equilibrium β phases. Although the equilibrium β phases were observed in the Mg matrix, the metastable β' phases kept in a large majority at the peak-aged stage, according to the SAED pattern presented in figure 9(f). The orientation relationship between the β phase and a-Mg matrix was [110]_β //[0001]_ɑ, (1–11)_β //[11–20]_ɑ and (−112)_β //[1–100]_ɑ [30, 32].

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With continuous ageing treatment, the CEE-AEC samples reached the early over-aged stage at 256 h. Figures 9(g)–(i) presented the morphology of the samples at early over-aged stage and corresponding SAED pattern recorded from [0001]_ɑ zone axis. Obviously, the lens-shaped β' phases have already formed the continuous network with the prolonged ageing time. There were some rhombic-shaped β₁ phases precipitating at the necks of the decomposing β' phases, as indicated in figure 9(h). The formation of the β₁ phases was considered as the result of the minimization of the shear strains around the β' phases. Some parts of the β₁ phases attached to the β' phases and others paralleled to the {1–100}_ɑ planes. When β₁ phases grew to a certain size, in situ transformation would happen between metastable β₁ phases and equilibrium β phases. Finally, the equilibrium β phases with a coarse plate shape would precipitate extensively on the {1–100}_ɑ planes in a long-term ageing treatment. The precipitation schematic diagram of the β phases was exhibited in figure 11, which illuminated the precipitation process from the mestable β'' phase to the equilibrium β phase. As shown in figure 9(i), the intensity of diffraction spots arose obviously compared to that in figure 9(f). The more evident diffraction contrasts could be attributed to the growth of β' phases and the formation of β₁ phases. Although the β₁ phases has formed at 256 h-aged treatment stage, the β' phases still precipitated predominantly, according to the diffraction feature of figure 9(i). This precipitation characteristics also demonstrated that the β' phases owned excellent thermal stability and the enhanced hardness related to the β' phases would maintain a long time at ambient temperature.

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