Magnesium (Mg) and its alloys have the lowest density among all structural metals and high specific strength and stiffness, and thus hold promise to be used for light-weighting in a wide range of applications [[1], [2], [3], [4]]. However, Mg alloys with a hexagonal close-packing (HCP) crystal structure generally exhibit poor rollability and formability. Mg alloy sheets with large dimensions and fine-grained structures are usually fabricated by hot rolling. Therefore, extensive efforts have been devoted to optimizing thermomechanical processing parameters to prevent the formation and growth of cracks and thus to improve rollability. Conventionally, multi-pass hot rolling with small thickness reductions (20%-30% at each pass) is applied to control cracking [5]. Such rolling processes are usually accompanied by intermediate annealing to reheat the sheets to desired temperatures. Consequently, the conventional rolling process of Mg alloy sheets is economically and timely inefficient for commercial applications.

Several researchers [[6], [7], [8]] reported that the strong and ductile Mg alloy sheets with homogeneously distributed refined grains could be achieved to a single pass reduction of 80% by applying a high strain-rate rolling process (2.9-9.6 s^-1). Our previous research [9] has also shown that the high strain-rate rolling could successfully produce AZ31 Mg alloy sheets with a uniform grain structure and fine grain size in a single pass with high reductions (>80%). Nevertheless, small edge cracks could be detected after high strain-rate rolling [6,10,11]. It is known that crack tips usually possess intense stress concentration [12], which, on one hand, could satisfy the energy requirement for crack propagation, and on the other hand, can promote dynamic recrystallization (DRX) [13,14].

Mg alloys with relatively low stacking fault energy present the stronger inclination of DRX during deformation. Various DRX mechanisms have been reported in Mg alloys, including twinning-induced DRX [5,[15], [16], [17]], subgrain related continuous DRX [18] and discontinuous DRX occurring by the bulging of high angle grain boundaries [1] upon plastic deformation. During high strain-rate rolling, twinning-induced DRX and shear bands related DRX have been reported as the predominant recrystallization mechanisms to avoid the formation of early cracks [5,19].

From the thermodynamic viewpoint [[20], [21], [22], [23], [24]], twice the surface energy density ((${{\gamma }_{\text{s}}}$) should be equal or less than the work per area ($w$) (($2{{\gamma }_{\text{s}}}-w\le 0$). Therefore, when the external work could not be accommodated by plastic deformation and satisfies the thermodynamic condition, the new surfaces of crack could form. Several researchers [6,15] pointed out that twinning and DRX were the most important activities for relieving high strain energy and thus preventing cracking upon rolling. Therefore, it is anticipated that the DRX weakening the stress concentration around the edge crack tips may restrict the growth of edge cracks. Dislocation emission from the sharp crack increases the curvature radius of the crack, followed by the clear crack blunting [[25], [26], [27]]. The crack blunting plays an important role in the fracture process, improving the cracking resistance upon rolling [28]. The process of crack blunting can influence the dislocation emission and distribution in front of the crack tip. The different deformation substructures between the blunted crack tip and the sharp crack tip could cause different microstructural evolution including dislocation slip and recrystallization behavior. The aim of the present study is to understand the underlying mechanism responsible for the formation of primary and secondary edge cracks upon high strain-rate rolling of Mg alloy.

The initial material is an as-cast AZ31 (3.19 wt.% Al and 0.81 wt.% Zn) alloy subjected to homogenization at 420 ℃ for 12 h. The sample for high strain-rate rolling has a dimension of 123 mm × 50 mm × 12 mm. Four effective strain rates (10 s^-1, 15 s^-1, 20 s^-1, and 29 s^-1) were applied to roll AZ31 Mg sheets in a single pass with a reduction of 82%. Before rolling, the samples were pre-heated to test temperatures (300 ℃, 350 ℃, and 400 ℃), held for 10 min. In order to arrest the microstructures, the as-rolled sheets were water quenched as soon as possible after the high strain-rate rolling. In order to analyze the edge crack, the microstructures around the edge crack were detected by optical microscopy (OM), electron backscatter diffraction (EBSD), and transmission electron microscopy (TEM).

The microstructures along the edge crack tips under different rolling conditions are illustrated in Fig. 1. It shows that the microstructures in the vicinity of the edge crack contain uniform and refined grains. The microstructures away from the edge crack have initial coarse grains at the effective strain rate of 10 s^-1 (see Fig. 1(a1)). The crack-front profile is usually considered to be infinitely sharp [25]. However, the majority of crack tips in the current study possess a finite curvature radius, referred to as crack blunting. Fig. 2 shows inverse pole figure maps, rotation axis distributions, misorientation angle distributions and KAM maps of the edge cracks under different effective strain rates at 300 ℃. In the initial coarse grains, necklace-like structures with the morphology of twins accompanied by fine recrystallized grains (in Area F) are detected at 10 s^-1. Aiming to analyze the formation of necklace-like structure, the rotation axis distribution and misorientation angle distribution for Area F of Fig. 2(a1) are shown in Fig. 3. The misorientation angles between grains in necklace-like structures and the initial coarse grain mainly keep the values of 30°-45° about the <-12 to 10> axis. These are boundaries between double twins and matrix (double twin-matrix with a characteristic disorientation of 38°<-12 to 10 >). This indicates that the formation of necklace-like structures is attributed to the twinning-induced DRX [15,16,29].

At the effective strain rate of 10 s^-1, a sharp crack is detected, while at effective strain rates of 15 s^-1, 20 s^-1 and 29 s^-1, the blunted edge cracks are observed. The kernel average misorientation (KAM) values indicate the local misorientation level in terms of the density of geometrically necessary dislocations [5]. The microstructures along the edge crack boundary all contain refined grains, and the KAM values along the edge crack boundary are relatively low (see Fig. 2(b1)-(b4)), which is attributed to the strain energy consumption by DRX. The KAM values at the sharp edge crack tip are estimated to be 0.270 (see Fig. 2(b1)), while the comparatively high KAM values at the blunted edge crack tip are calculated to be 0.776, 1.060, and 0.857 at strain rates of 15 s^-1, 20 s^-1 and 29 s^-1, respectively (see Fig. 2(b2)-(b4)). Therefore, the KAM values of the blunted crack tip are around three times the KAM values of the sharp edge crack tip. This suggests that dislocations pile up at the blunted crack tips.

The TEM images in front of the blunted edge crack tip and the sharp edge crack tip are compared in Fig. 4. In Fig. 4(a1) and (a2), the dislocation density is high, and a massive amount of dislocations pile up on grain boundary near the blunted edge crack tip (see the white arrow in Fig. 4). The dislocation density at grain triple junctions is even higher than that of grain boundaries. For the sharp edge crack, it can be seen from Fig. 4(b1) and (b2) that there are rare dislocations around the sharp edge crack tip. Physically, an atomically sharp crack will be blunted by shifting on the atomic plane. The crack blunting is due to the dislocation emission from the stressed crack tip in the situation where the grain size is larger the 20 nm [30,31]. The stressed crack tip emits dislocations, which leaves an equal and an opposite dislocation at the boundary of the stressed crack tip. This equal and opposite dislocation dissipates with the continuous deformation, followed by the increase of the curvature radius of the crack and crack blunting. For both the sharp edge crack tip and the blunted crack tip, DRX is almost complete. Theoretically, the occurrence of DRX consumes substantial dislocations, whereas the dislocation density of the blunted edge crack tip is relatively high (see Fig. 4(a1), (a2)). It is suggested that the blunted edge crack tip continues dislocation emission after DRX until being blunted to a certain extent, and emitted dislocations are hindered by fine DRX boundaries.

It is generally known that the DRX mechanism is diverse with the temperature at low strain rate deformation (10^-5 s^-1 to 10^-1 s^-1) [32]. However, it was reported that the achievement of high strain-rate rolling at 250-400 ℃ with strain rates of 2.99.6 s^-1 was attributed to the preferred initiation of twinning and the subsequent DRX inhibit competitive crack initiation in ZK60 [[6], [7], [8]]. They also indicated that deformation twins played a significant role in DRX at 250-400 ℃ during high strain rolling. Therefore, the twinning-induced DRX is the predominant DRX mechanism at high strain-rate rolling. After the twins are all occupied by the recrystallized grains, the transition of prevail DRX mechanism may occur [29]. The edge cracks nucleate and propagate when the plastic deformation cannot be accommodated. The process of crack blunting can affect the dislocation emission and distribution in front of the crack tip. For the sharp crack tip which is too late to emit dislocation, the intense stress concentration favors twinning. Therefore, the DRX mechanism in front of the sharp and blunted crack tips may be different.

The misorientation angle distribution for area A around the sharp crack tip is shown in Fig. 2(c1). It can be seen that the high misorientation angles (65°-90°) have high Rel. frequency (12%) in contrast to low misorientation angles (15°-35°) with a Rel. frequency of 5.9%. The most frequently detected boundaries in Fig. 2 (c1) are grains with the misorientation angles of 2°-15° about the <-12-10> axis, 15°-35° about the <-12-10> axis, 35°-50° about the <-12-10> axis and 65°-90° about the <-12-10> axis and <10-10> axis. These are the boundaries between: (i) double twins and matrix (double twin-matrix with a characteristic disorientation of 38°<-12-10 >), (ii) extension twins and the matrix (extension twin-matrix with a characteristic disorientation 86°<-12-10 >) and (iii) two extension twins of different types (extension twin-extension twin with a characteristic misorientation of 7.4° <10-10 >) [15]. It can be concluded that most of the recrystallized grains nucleate at twin boundaries around the sharp edge crack tip, indicating that twinning-induced DRX is the dominant mechanism. However, the rotation axis distributions and misorientation angle distributions profiles for areas B-D around the blunted crack tips in Fig. 2(c2)-(c4) show that the majority of grains have relatively low misorientation angles (<40°), and the boundaries with rotation axis about <-12-10> has only a small fraction, which suggests that most of the recrystallized grains are not induced by twinning but transformed from subgrains. It can be seen from Fig. 2(a2)-(a4) that small isolated DRX grains A1-C1 and isolated recrystallized area E indicate that DRX can take place in isolated grains [33]. In this case, a new DRX grain is developed near the primary grain boundary. The dislocation cells (see the ellipse area in Fig. 4) and dislocation walls (see the red arrow in Fig. 4) are detected inside recrystallized grains. According to the above analysis, it is concluded that the blunted edge crack keeps emitting dislocations after DRX and the emitted dislocations are beneficial to the formation of subgrains. Clearly, subgrains with low angle grain boundaries gradually transform to recrystallized grains with high angle grain boundaries (>15°), leading to the formation of isolated new recrystallized grains. Therefore, the prevailing DRX which is induced by subgrain formation causes the grain refinement in front of the blunted edge crack tip. Especially at high temperatures, the fast dislocation rearrangement occurs, leading to the formation of low-angle boundary, finally promoting the subgrain-induced DRX [32]. As the crack tip is blunted, the stress ahead of the crack tip can be relaxed [34]. Hence, the local stress of the blunted edge crack tip may not reach the critical stress for twinning. This may be one of the reasons why the DRX mechanism is different from the sharp edge crack tip.

It is observed in Fig. 1 that the long and thin cracks connect with the blunted edge crack. Ohr et al. [35] suggested that when the crack was blunted to a certain extent, the blunted crack hardly propagated until the nucleation of a new sharp crack. An internal back stress due to the dislocations emitted from the crack tip accommodates the stress intensity as a result of an applied load, contributing to the inhibition of further crack propagation, followed by an increase of fracture toughness of materials [28]. This means that the blunted cracks with large radii are rarely propagated again. Hence, the present blunted edge cracks hardly continue growing, and the long and thin cracks can be categorized as the secondary crack.

The voids are detected at grain triple junctions, as shown in the rectangle area and square area in Fig. 1(a1) and (a2). The thin crack tends to connect with the voids formed at the triple junctions (see the rectangle and square areas in Fig. 1(a1) and (a2) and the circle in Fig. 2(b3)). The wedge cracking [36] and voids [15,37] are usually observed at grain boundaries especially in grain triple junctions. This is because dislocations pile up against grain boundaries and accumulate a high stress concentration, in which case the nucleation of voids is required to release such a high stress concentration [38]. The voids grow in a plastically deforming matrix until they link with one another causing final fracture [39]. The KAM values at grain triple junctions and around voids in the fine-grained area is relatively high (see the circles in Fig. 2(b2) and (b3)), and severe dislocation pile-ups are found at grain triple junctions (see Fig. 4(a2)). This implies that grain triple junctions could provide enough strain energy for the requirement of the formation of voids. In front of the blunted crack tip, the KAM values are high (see Fig. 2(b2)-(b4)) and substantial dislocations are detected (see Fig. 4(a1) and (a2)). Therefore, the dislocation pile-up on grain boundaries especially in the triple junctions in front of the blunted crack tip contributes to the nucleation, growth and coalescence of voids, inducing the formation of sharp cracks. These sharp cracks tend to link with the blunted crack, leading to the generation of secondary cracks of the blunted crack.

A schematic plot of the formation of primary crack and the generation and growth of the secondary crack for the blunted edge crack is illustrated in Fig. 5. The deformation is first accumulated in the initial coarse grains located in front of the edge crack tip, dislocations are then emitted from the edge crack, the stress concentration around the edge crack tip is relaxed accordingly, and finally the local stress cannot reach the critical stress for twinning. Meanwhile, these emitted dislocations accelerate the formation of subgrains, and then subgrains with low misorientation angles transformed into recrystallized grains with high angle grain boundaries. This causes the grain refinement in front of the edge crack tips as well as the onset of crack blunting. The dislocations emitted from the blunted edge crack pile up in front of the boundaries of refined grains, resulting in the increase in local work hardening and stress concentration along grain boundaries especially in the triple junctions. Voids start to nucleate when the stress concentration reaches the critical stress for the formation of voids. The growth and connection of voids lead to further crack propagation. The crack linking to the blunted edge crack tip finally causes the generation and growth of the secondary crack.

Upon high strain-rate rolling, the stressed edge crack tip emits dislocations leaving an equal and an opposite dislocation at the edge crack tip boundary, followed by edge crack blunting. The emitted dislocations promote the formation of subgrains, and the stress concentration around the crack tip is relieved during the process of crack blunting, resulting in the subgrain-induced DRX, which plays a key role in grain refinement around the blunted edge crack tip. At the sharp edge crack tip, twinning-induced DRX is the predominant mechanism. The motion of emitted dislocations is inhibited by the DRX grain boundaries, which leads to an increase in local work hardening and also an intense stress concentration along grain boundaries especially in the triple junctions. This causes the formation, growth and coalescence of voids, which eventually results in the nucleation of the secondary cracks.