J.Cent.South Univ.(2025) 32: 706－726
1 Introduction
Mg alloys are known as the lightest structural materials, and are widely used in the field of 3C electron, automobile and aerospace due to their low density and high specific strength [1-3]. Although wrought Mg alloys have many attractive properties,their applications are still limited due to their strong basal texture [4, 5]. It is well known that the critical resolved shear stress (CRSS) of basal slip is far lower than that of non-basal slip systems at room temperature [6, 7]. Thus, limited by the special hexagonal close-packed crystal structure, only two independent slip systems are activated at room temperature for conventional rolling, which results in an extremely strong basal texture [8, 9]. The strong basal texture gives rise to high anisotropy in the plate and increases the difficulty of strain accommodation in the thickness direction at room temperature [10, 11].
To reduce the anisotropy of Mg alloys, some studies have focused on different processes to weaken the texture. TIAN et al [12] obtained sheets with weakened basal texture by width-limited rolling, and the plasticity was significantly improved. Moreover, rare earth elements such as Y and Ce can cause double peak texture, which leads to a significant decrease in texture intensity [13, 14]. However, the complex process and expensive rare addition increase the consumption of test, which limit large-scale industrial usage of Mg alloys [4, 15]. Thereby, how to develop a new method with low cost is urgent. As is well known, with temperature increasing, the CRSS of non-basal slips is significantly reduced, which is beneficial to activate the non-basal slips [16]. The activation of non-basal slip further causes the grain orientation to be deflected, that is, the c-axis of the grain is no longer predominantly parallel to the normal direction [17]. Therefore, high temperature deformation can be used to weaken texture and reduce anisotropy without additional consumption. WANG et al [18] indicated that the in-plane anisotropy of ZK60 specimen was significantly weakened at 300 ℃ due to the activation of non-basal slip. WANG et al [19] concluded that at higher temperature, the increase of non-basal slips activity led to the decrease of normal direction anisotropy of AZ31 magnesium alloy. With the increase of temperature, the slips are activated due to the decrease of CRSS [20], and the difference of CRSS between basal slip and pyramidal slip decreased [13]. Meanwhile, there is no doubt that high-temperature rolling is also a potential effective process to obtain excellent magnesium alloy sheets with a weakened texture and great mechanical properties. However, rolling equipment at high temperature should be required and the energy cost as well as the loss of rolling mills is high. Therefore, to start the non-basal slips induced by thermal activation at higher temperatures with low cost is necessary and the process should be designed.
Thus, in this paper, a new two-step thermal pre-rolling technology is developed to achieve the non-basal slip on AZ31 Mg alloys sheet at higher temperatures. After that, normal rolling is followed to save costs. The microstructure evolution, texture behaviors and the mechanical response are investigated systematically. The detail mechanisms are discussed deeply.
2 Experimental procedure
Commercial AZ31 alloy sheets (Mg-3 wt.% Al-1 wt.% Zn) with a thickness of 4 mm were used as the investigated material in this work. The dimensions of the initial square samples are 80 mm×80 mm. Before rolling, all of the samples were annealed at 350 ℃ for 1 h, named as-received sample. The as-received samples were pre-rolled from 4 mm to 2.2 mm by 1 pass of 45% thickness reduction at different pre-rolling temperatures of 350 ℃ (350 ℃-PRI) and 450 ℃ (450 ℃-PRII). Then the pre-rolled samples were further rolled to 1 mm by 5 passes at 300 ℃, and the reduction of per pass was 15%. For comparison, other samples without pre-rolling were normal rolled by 5 passes at 300 ℃, and the reduction of per pass was 25%. During each pass, the samples were reheated for 5 min again to the rolling temperature. The samples after each pass are named the style of “temperature-pass”, such as 350 ℃-1P and 300 ℃-1P. Moreover, in order to obtain the performance of each pass sample, some samples were selected to anneal at 300 ℃ for 1 h in each pass, termed “temperature-passA” (e.g., 350 ℃-1PA, 300 ℃-5PA).
Dog-bone tensile sample with a gauge dimension of 12 mm×3 mm was machined from the samples annealed at 300 ℃ for 1 h. Then the uniaxial tensile tests were conducted at room temperature using a DNS200 electronic universal testing machine along rolling direction (RD), aligned 45°, and the transverse direction (TD). The strain rate was 1×10-3 s-1. The n-value was obtained from the uniform plastic deformation region of the true stress-strain curve. Selective tests were carried out at a permanent uniform strain without failure of 12% to measure the changes in sample dimensions (length and width) to determine the Lankford value (r-value).
The schematic diagram of rolling processes is given in Figure 1. After grinding, polishing and etching with a solution of picric acid, all samples annealed at 300 ℃ for 1 h were subjected to optical microscopy (OM; Leica DM2700 M). The microstructure evolutions were measured by OM and electron backscatter diffraction (EBSD) technologies. After grounding, polishing and electropolishing with commercial AC2 solution at a temperature of -20 ℃ using a voltage of 20 V for 90 s, EBSD analysis equipped with a scanning electron microscope (GEOLGSM-7800F FEG SEM) was used to identify the microstructure, and then an HKL channel 5 system was used to obtain the detailed analysis results.
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3 Results and discussion
3.1 Microstructure evolution before annealing
Figure 2 shows the OM and EBSD maps of the as-received sample. The microstructure is uniform and consists of equiaxed grains with an average size of 9.67 μm. No twins appear in the grain. As shown in Figure 2(b), most grains lay on basal plane, that is, the c-axis is basically parallel to normal direction (ND). Therefore, the initial AZ31 Mg alloy sheet presents a typical basal texture. Figure 2(c) gives the (0002) pole figure, which expresses a typical rolling basal texture with a relatively high texture intensity of 18.12.
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Optical micrographs after rolling are shown in Figure 3. It can be clearly seen from Figure 3(a) that a large number of twins appear in the grain of pre-rolling sample due to large reduction of 45%. In addition, recrystallized grains appear at twin and grain boundaries. However, the size of deformed grains decreases significantly, while it increases obviously on recrystallized grains with the increase of pre-rolling temperature. The volume fraction of dynamic recrystallized grains of the 450 ℃-PRII sample is larger than that of the 350 ℃-PRI sample, which is because the fine dynamic recrystallized grains begin to grow and consume the original grains. Compared with the 450 ℃-PRII sample, there are a large number of twins to coordinate plastic deformation in coarse grains of the 350 ℃-PRI sample, which indicates that the deformation mechanism is different at various pre-rolling temperatures. In the first pass rolling, the microstructure is composed of coarse grains with twins and recrystallized grains. However, the size of coarse grains decreases, and the grains become fine and equiaxial, especially in the 450 ℃-1P sample, as presented in Figures 3(d) and (e). After that, from the second pass to the fifth pass, the distribution of the microstructure becomes more uniform, and twins appear in almost every grain. For the comparison group, the microstructure evolution on directly rolling sample at 300 ℃ is different from that of pre-rolling samples, which is still not uniform and composed of fine recrystallized grains and coarsely deformed grains, as shown in Figures 3(c) and (f). To further investigate the deformation mechanism, EBSD is used for further analysis.
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Figure 4 presents the EBSD maps of select samples without annealing. As shown in Figures 4(a) and (b), coarse grains are separated by twin boundaries, but the grain fragmentation is not complete, and some large grains still exist. In addition, the recrystallization phenomenon is obvious at the grain boundaries. As shown in Figures 4(d) and (e), the number of twins inside the coarse grains increases with the rolling process continues, dividing the large grains into small grains. It can be clearly seen in Figure 4(c) that many obvious recrystallized grains can be observed at the twin boundaries of the sample without pre-rolling. Moreover, the corresponding microstructure compositions of grains after dynamic recrystallization (DRXed grains), subgrains and deformed grains are also given in Figure 4. The size of the subgrains is larger than that of the DRXed grains and deformed grains. Dynamic recrystallization (DRX) leads to grain refinement as rolling progresses. As shown in Figures 4(d) and (e), the DRXed grains consist of a large number of fine DRXed grains, while the number is significantly reduced in Figures 4(a) and (b).
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3.2 In-grain misorientation axes analysis
It is known to all that with the increase of temperature, the CRSS of non-basal slip systems reduces. Prismatic slip and pyramidal slip play an important role in the high temperature deformation process of magnesium alloy. Especially when the c-axis is parallel to the external force and in the hard orientation state, the pyramidal <a+c> slip can well coordinate the deformation of c-axis direction. In addition, previous studies have also proved that the activation of non-basal slip is closely related to dynamic recrystallization. HAN et al [21] pointed out that non-basal slips play an important role in the dynamic recrystallization. WEI et al [22] reported that the activation of slips is influenced by grain size. Therefore, the microstructure evolution may be related to the activation of non-basal slips. In order to find out the effect of non-basal slips on deformation, in-grain misorientation axes (IGMA) and texture component maps are used for further analysis.
By observing the intensity distribution of IGMA in a single grain, researchers can basically determine which slip system is activated. In theory, IGMA distribution can be divided into two types. When the IGMA intensity distribution is around <0001>, it is mainly caused by the activation of prismatic slip system [23]. However, when the intensity distribution is concentrated at <u v t 0>, which results from basal slip or pyramidal <a+c> slip [24]. Further, combined with the Schmid factor, the main activated slip system can be accurately determined. Figure 5 presents the IGMA distribution of numbered grains in Figure 4. As shown in Figure 5, several grains are selected from the denoised EBSD data to analyze their IGMA and the misorientation angle in this paper ranges from 2.5° to 5° [25, 26]. Although the IGMA intensity distribution of some grains spreads to (0001), the intensity distribution of selected grains is mainly concentrated around <u v t 0>. Moreover, the maximum intensity is greater than 2 and thus it is generally considered that there is a preferential IGMA. Therefore, it is speculated that basal slip and pyramidal <a+c> slip dominate the deformation.
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In order to make a qualitative judgment on non-basal slip, Schmid factor is used for further analysis. The Schmid factor maps of as-received sample, 350 ℃-PRI sample and 450 ℃-PRII sample along RD and TD are shown is Figure 6. And the Schmid factor maps of 300 ℃-1P sample, 350 ℃-1P sample and 450 ℃-1P sample along RD and TD are shown in Figure 7. With the increase of rolling temperature, the difference of CRSS between basal slip and non-basal slip decreases. Therefore, the magnitude of Schmid factor is the main factor for the activation of slip systems. As shown in Figure 6, the Schmid factor values of basal slip of 350 ℃-PRI sample and 450 ℃-PRII sample along RD are 0.215 and 0.251, respectively. However, the values of pyramidal <a+c> slip of 350 ℃-PRⅠ sample and 450 ℃-PRII sample along RD are 0.395 and 0.375, respectively. It is clearly evident that the Schmid factor (SF) of pyramidal <a+c> slip is higher than that of basal slip. Moreover, the same trend is seen in TD. In the first pass, the same trend continues. In summary, the Schmid factor value of basal slip is lower than that of that of pyramidal <a+c> slip. Combined with the analysis of IGMA and Schmid factor, pyramidal <a+c> slip plays an important role in the deformation process.
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Figure 8 shows the texture component maps of 350 ℃-1P sample. It is obvious that the basal texture and pyramidal texture dominate the overall texture component, while prismatic texture accounts for only a very small part, which indicates that pyramidal texture has a great influence on the deformation of pre-rolling sample. The activation of pyramidal <a+c> slip will weaken the texture and make the deformation easier. Therefore, the pre-rolling process is likely to promote the pyramidal slip and change the deformation mechanism.
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3.3 Dynamic recrystallization mechanism
3.3.1 Discontinuous dynamic recrystallization
In the hot temperature rolling process, dynamic recrystallization behavior is complex, and different DRX mechanisms may play an irreplaceable role [27]. Continuous dynamic recrystallization (CDRX) and discontinuous dynamic recrystallization (DDRX) are two main recrystallization mechanisms, which may act independently or simultaneously [28, 29]. To clarify the recrystallization behavior during deformation, enlarged EBSD data were investigated.
Region II in Figure 4(c) is enlarged in Figure 9. In Figure 9(b), grains G1 and G12 are two parent grains of large size, and a chain of small recrystallized grains named G2-G11 can be clearly seen at the serrated grain boundaries. Figure 9(e) presents the misorientation profile along L2 marked in Figure 9(a), where the curve has two sharp peaks. Peak misorientation angles are 42.6° and 39.5°, respectively. Moreover, Figure 9(c) shows the (0002) pole figure of selected grains. From the pole figure, parent grains G1 and G10 are obviously deviated from the polar center about 38.5°. In general, serrated boundaries have high local stain gradients, which become a potential nucleation site for discontinuous dynamic recrystallization (DDRX) of grain boundary bulging [30, 31]. It can be clearly seen in Figure 9(b) that the parent grains show obvious serrated boundaries that generate DRXed grains, which is the feature of DDRX [31, 33]. Furthermore, DDRX helps to produce new grains with different orientation, which contributes to weak texture [32, 33]. Figure 9(d) presents the misorientation profile along L1. Neither point-to-origin misorientation nor point-to-point misorientation along L1 exceeds 10°, which indicates that the misorientation angle gradient is relatively low inside the parent grains.
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3.3.2 Continuous dynamic recrystallization
The region III in Figure 4(c) is enlarged in Figure 10. As shown in Figure 10(b), the coarse grain is divided into several smaller grains by internally generated twins, which is marked with a green line, but possibly due to the large strain, recrystallized grains induced by twins are not fully detected in EBSD. However, observing the region corresponding to region III in the IPF map and DRXed grains map, the small recrystallized grains can be seen. Theoretically, the orientation angle of the _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M001c.jpg)
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Figure 11 shows the enlarged region I in Figure 4(a). As shown in Figure 11(a), a grain is divided into several subgrains by some LAGBs. However, it can be predicted that the misorientation angle between G2 and G4 will gradually increase and develop into recrystallized grains as the deformation continuous. Figures 11(d) and (e) present the misorientation angle map along L4 and L5. The misorientation angle of the point-to-point and point-to-origin curves does not exceed 15° either along grain interior L4 or across grain boundary L5. In addition, it is evident that the misorientation angle from point-to-origin is gradually increasing, which is characteristic of CDRX [32]. As can be seen from the pole figure (Figure 11(c)), there is no significant difference in the orientation of all grains.
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3.4 Influence of double twin
It is well known that twinning can cause grain orientation changes so that the grains are in a favorable orientation to slip and further weaken the texture [37, 38]. Double twins are easily produced when the deformation temperature is higher than 300 ℃ [39, 40]. However, when the temperature is at 450 ℃, the effect of the double twin decreases. In the present work, due to the large strain of unannealed samples, it was difficult to distinguish double twins in the band contrast map. However, it can be clearly seen in the optical microscopy map of Figure 3(a) that the 350 ℃-PRI sample has a large number of double twins compared with the 450 ℃-PRII sample. Figure 12 shows the magnified view of region IV in Figure 4(d). Figure 12(d) shows the line profile for the microstructure angle along L6 in Figure 12(a). The red point-to-point line presents the misorientation angles across the double twin T, which are 39.6° and 36.5°, respectively. The existence of a
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In summary, as shown in Figures 3(a), (b) and 11, CDRX mainly occurs in the coarse grains’ boundaries or twins generated within the coarse grains. However, due to the higher temperature of pre-rolling at 450 ℃, the fraction of twins is less than that of sample pre-rolled at 350 ℃. As the deformation continues, more twins are generated in the microstructure, especially in the 350 ℃ sample, thus the large grains are divided by twins. Therefore, the intersection of twins provides more nucleation site for DRX. As shown in Figures 3(d) and (e), DDRX induced at twin boundaries and CDRX generated inside twins are very common in subsequent normal rolling. As discussed in Section 3.2, the activation of non-basal slips has an important impact on the microstructure evolution. The activation of pyramidal <a+c> slip in the samples pre-rolled at 350 ℃ is better than that of the samples pre-rolled at 450 ℃, which leads to more obvious recrystallization in the samples pre-rolled at 350 ℃, while the samples pre-rolled at 450 ℃ may focus more on the growth of DRXed grains.
3.5 Texture evolution during rolling
Figure 13 shows the (0002) pole figures of the selected samples. It can be clearly seen in Figure 13(c) that the 300 ℃-1P sample exhibits a very high texture intensity of 20.64, which is higher than that of the as-received sample. Because the basal slip also plays a role in normal rolling. Therefore, it shows a typical basal texture, where the c-axis is parallel to ND, which results in poor plastic formability [37]. In contrast, from the pre-rolled pass to the first pass, the texture intensity of the 350 ℃-1P sample decreases significantly from 17.14 (Figure 13(a)) to 10.70 (Figure 13(d)) and that of the 450 ℃-1P sample decreases from 18.03 (Figure 13(b)) to 15.89 (Figure 13(e)). The former drops by as much as 37.6%, while the latter only shows 11.9%. After larger reductions of 45% and 15%, the texture intensities of the 350 ℃-1P and 450 ℃-1P samples are lower than that of the 300 ℃-1P sample, which is subjected to 25% reduction rolling. This suggests that pre-rolled process can dramatically weaken the basal texture, especially for the 350 ℃-1P sample, which is related to the activation of the pyramidal <a+c> slip determined by IGMA and Schmid factor. Texture weakening has been proven to be an effective method to improve the formability of magnesium alloys [38]. For this study, the texture distribution of pre-rolled samples is more dispersed than that of the normal rolled sample, i.e., the c-axis of grains inclined from ND to RD is more obvious. It is interesting that the 350 ℃-1P sample generated a double peak texture in RD, and the texture also shows an inclination from ND to TD, as shown in Figure 13(d). However, as the pre-rolling temperature increases to 450 ℃ in Figure 13(e), the 450 ℃-1P sample does not develop an obviously split texture and does not incline toward TD, which spreads to RD accompanied by a weak split. The texture evolution can be affected by rolling temperature, rolling reduction and initial texture [43, 44]. Both double and single peaks may occur in the rolled AZ31 sheet [44, 45]. HAN et al [46] also revealed that AZ31 sheet with equal channel angular rolling shows an unusual double peak texture in RD, and the formability is better than the conventional rolled sheet with strong basal texture. Previous studies have shown that the splitting of the basal poles in RD is associated with the double twin, while the inclined basal texture toward TD is related to the activity of prismatic slip [46-49]. In general, the strain in the width direction is provided by prismatic slip [24]. Therefore, it is speculated that the double peak texture is caused by double twins. However, the inclined texture toward the TD is associated with pyramidal slip, which is different from the previous studies. The intensity distribution of IGMA and texture component have shown that the inclined texture from ND to TD is mainly caused by pyramidal slips, especially pyramidal <a+c> slip rather than prismatic slip. Compared with 450 ℃-PRII sample, 350 ℃-PRI sample shows larger pyramidal <a+c> slip Schmid factor of 0.395, which indicates that 350 ℃-PRI sample is more likely to activate pyramidal <a+c> slip. Therefore, 350 ℃-1P sample shows significant inclined texture along TD. Additionally, it can be clearly seen in Figures 6(a) and (a) that basal slip in RD of 300 ℃-1P sample exhibits a smaller SF value of 0.179, while that of as-received sample is 0.203, which means that basal slip is still at work in the 300 ℃-1P sample, resulting in a higher texture intensity of 300 ℃-1P.
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3.6 Microstructure evolution after annealing
Figure 14 shows the IPF, misorientation angle maps and grain size distribution maps of the fifth pass samples annealed at 300 ℃ for 1 h. After annealing, the deformed grains and twins disappeared completely, and the microstructures present a typical basal texture and are composed of equiaxed recrystallized grains. The average grain sizes of the 300 ℃-5PA sample, 350 ℃-5PA sample and 450 ℃-5PA sample are 7.04, 4.31 and 5.70 μm, respectively. The 350 ℃-5PA sample shows the smallest average grain size of 4.31 μm. Moreover, although the average grain size of the 300 ℃-5PA sample is 7.04 μm, it can be seen clearly from the grain size distribution map that the microstructure of the 300 ℃-5PA sample shows a large volume of fine grains, resulting in a smaller average grain size in the final calculation. Different from the 300 ℃-5PA sample, the smaller average grain sizes of 350 ℃-5PA and 450 ℃-5PA are attributed to the fact that the grain sizes of those two samples are mostly distributed at approximately 5 μm. That is, the grain size distribution of the latter two is more uniform. The misorientation angles are characterized by recrystallized grain orientations of 35.75° and 35.00° for the 350 °C-5PA and 450 °C-5PA sample, respectively, while that of the 300 °C-5PA sample is 33.71°. Due to the activation of non-basal slip systems, the grain orientation is more dispersed, which makes the deformation easier and the accumulated energy increases. Furthermore, the grain orientation dispersed after recrystallization, and the 350 ℃-5PA sample is the most dispersed.
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Figure 15 shows the (0002), _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M005c.jpg)
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4 Mechanical properties and anisotropy
Figure 16 illustrates the true stress-strain curve of sheets annealed at 300 °C for 1 h in the tensile directions of RD, 45° and TD. The curves of other pass samples are not shown due to similar trends. Regardless of the rolling temperature and reduction change, the flow stress increases with increasing strain until it reaches the maximum value and then decreases with increasing strain until fracture. Moreover, the ultimate tensile strength (UTS), yield stress (YS), r-value and n-value of the three fifth pass annealed samples are presented in Figure 17 for a better review. The YS in TD is the largest, but that in RD is the smallest, continuously decreasing from TD to RD. The regular changes in YS are reflected in the texture distribution. As shown in Figure 15, all the samples present an inclined texture from ND to RD. This indicates that the grain orientation is more dispersed in RD and the worst in TD.
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Table 1 shows the average mechanical properties of three fifth pass annealed samples. For the three fifth pass annealed samples, the average fracture elongation is 17.31%, 18.83% and 16.81%, respectively, and the 350 ℃-5PA sample shows the largest average fracture elongation. The average YS and UTS of 350 ℃-5PA sample are higher than those of 300 ℃-5PA sample and 450 ℃-5PA sample. As shown in Figure 14, the smallest grain size is obtained in the 350 ℃-5PA sample. It is generally thought that the plasticity and strength of metal materials cannot be improved at the same time, but fine grain strengthening is one of the few ways to increase both strength and plasticity [50-52]. According to the Hall-Petch formula, the smaller the grain size is, the greater the yield strength of the material is. In addition, when the material with fine grains is affected by an external force, the plastic deformation can be dispersed in more grains and is uniform, which results in the improvement of plasticity. These results show that the mechanical properties of the 350 ℃-5PA is the best. The sample subjected to 350 ℃ pre-rolling shows a smaller _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M009c.jpg)
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| Sample | YS/MPa | UTS/MPa | _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M011c.jpg) | _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M012c.jpg) | _Xml/alternativeImage/B0208A2C-E139-49a8-9CAC-347280CA9EB2-M013c.jpg) |
|---|---|---|---|---|---|
| 300 ℃-5P | 186.90 | 307.16 | 17.31 | 4.38 | 0.26 |
| 350 ℃-5P | 194.12 | 312.87 | 18.83 | 2.47 | 0.32 |
| 450 ℃-5P | 175.70 | 291.94 | 16.81 | 2.70 | 0.27 |
Figure 18 shows the ΔYS and Δr values of three fifth pass samples. As shown in Figure 18(a), the ΔYS of 300 ℃-5PA sample, 350 ℃-5PA sample and 450 ℃-5PA sample are 5.77, 3.75 and 6.20, respectively. After five passes, the ΔYS of the 350 ℃-5PA sample is the smallest, indicating that the plane anisotropy reaches the minimum compared with the 300 ℃-5PA sample and 450 ℃-5PA sample. A similar phenomenon is also reflected in Δr. As shown in Figure 18(b), the Δr of three fifth pass annealed samples are 2.28, 0.65 and 1.14, respectively. 300 ℃-5PA sample presents a larger Δr, while after the pro-rolling process, the Δr of 350 ℃-5PA sample and 450 ℃-5PA sample decreases significantly, especially for the 350 ℃-5PA sample, which indicates that pre-rolling process is an effective way to reduce the anisotropy of the AZ31 magnesium sheet. SHI et al [53] pointed out that prismatic slip is the main reason for the high anisotropy of magnesium alloy sheets, while the activity of pyramidal <a+c> slip. The r-value of sheet decreases significantly, that is, the anisotropy weakens significantly [54]. In this study, the activity of pyramidal <a+c> slip promotes the occurrence of DDRX, dispersing the grain orientation, weakening the basal texture and further weakening the anisotropy. Under the influence of the deformation condition, the activity of the pyramidal <a+c> slip of 350 ℃ pre-rolling sample is the most obvious, which results in weakened anisotropy.
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5 Conclusions
Pre-rolling is performed at 350 and 450 ℃ to investigate the anisotropy weakening of AZ31 Mg alloy sheet. After the rolling process, the anisotropy of pre-rolling samples is compared with the sample without pre-rolling. The main conclusions are summarized below.
The high temperature pre-rolling process activates non-basal slip and disperses grain orientation, which results in dynamic recrystallization more active. In addition, due to the influence of twins, the texture intensity decreases significantly and the texture distribution disperses obviously.
The texture intensities of the 350 ℃-5PA and 450 ℃-5PA samples are 10.97 and 10.61, respectively, which are lower than that of the 300 ℃-5PA sample (12.80). The 350 ℃-5PA sample shows a unique texture, where the c-axis is slightly rotated from RD to 45°, which is beneficial to weaken anisotropy.
Compared with the sample without pre-rolling and the 450 ℃-5PA sample, the 350 ℃-5PA sample shows the smallest Δr of 0.65, which indicates that the 350 ℃-5PA sample exhibits less anisotropy in the sheet. In summary, the 350 ℃ pre-rolling process is an effective method to obtain magnesium alloys with a weakened anisotropy.
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[J]. Journal of Materials Research and Technology, 2022, 19: 1627-1649. DOI: 10.1016/j.jmrt.2022.05.164.YANG Chao-yang, WANG Li-fei, XUE Liang-liang, HUANG Qiu-yan Huan, XIA Da-biao, FU Xin-wei, SONG Bo, ZHENG Liu-wei, WANG Hong-xia and KWANG Seon-Shin declare that they have no conflict of interest.
YANG Chao-yang, WANG Li-fei, XUE Liang-liang, HUANG Qiu-yan, XIA Da-biao, FU Xin-wei, SONG Bo, ZHENG Liu-wei, WANG Hong-xia, KWANG Seon-Shin. Weakened in-plane anisotropy of AZ31 magnesium alloy sheet induced by pre-enhanced non-basal slips during hot rolling [J]. Journal of Central South University, 2025, 32(3): 706-726. DOI: https://doi.org/10.1007/s11771-025-5856-z.
杨朝阳,王利飞,薛亮亮等.热轧过程促进非基面滑移开动弱化AZ31镁合金板材平面各向异性研究[J].中南大学学报（英文版）,2025,32(3):706-726.