Advances in Microstructural Evolution and Mechanical Properties of Magnesium Alloys Under Shear Deformation

Abstract

Magnesium (Mg) alloys are the lightest metals used in engineering structures, making them highly valuable for lightweight designs in aerospace, automotives, and related industries. Their low density offers clear advantages for reducing product weight and improving energy efficiency–key priorities in modern manufacturing. However, their unique crystal structure leads to notable drawbacks: low plasticity at room temperature, uneven performance across different directions, and inconsistent strength under tension versus compression. These issues have severely limited their broader application beyond specialized use cases. Shear deformation methods address this challenge by creating high strain variations and complex stress conditions. This approach provides an effective way to regulate the internal structure of Mg alloys and enhance their overall performance, overcoming the inherent limitations of their crystal structure. This paper systematically summarizes current research on using shear deformation to process Mg alloys. It focuses on analyzing key structural changes induced by shear, including the formation and evolution of shear–related features, real–time grain reorganization, crystal twinning processes, the distribution of additional material phases, and reduced directional performance bias. The review also clarifies how these structural changes improve critical mechanical traits: strength, plasticity, formability, and the balance between tensile and compressive strength. Additionally, the paper introduces advanced shear–based processes and their derivative technologies, such as equal–channel angular extrusion, continuous shear extrusion, and ultrasonic vibration–assisted shearing. It also discusses strategies for constructing materials with gradient or mixed internal structures, which further expand the performance potential of Mg alloys. Finally, the review outlines future development directions to advance this field: developing shear processes that combine multiple physical fields, conducting real–time studies of microscale mechanisms, designing tailored shear paths for high–performance Mg alloys, and evaluating long–term service performance. These efforts aim to promote both theoretical innovation and industrial application of shear deformation technology for Mg alloys.

1. Introduction

Mg alloys are the lightest structural metallic materials, and their high specific strength and castability make them pivotal for lightweight designs in aerospace, automotive, and electronics industries [1]. In the context of global energy crises and global carbon reduction goals, their role in cutting energy consumption and emissions has become increasingly strategic [2]. Advanced die–casting technology– which enables integrated forming of complex thin–walled parts–has expanded their use in automotive manufacturing: once limited to non–structural or semi–structural components, Mg alloys now serve in critical load–bearing structures. This integration reduces part counts by over 60% and assembly costs by 30–40%, boosting production efficiency [1,3]. Additionally, their strong electromagnetic shielding and near–complete recyclability meet the signal stability requirements of electronics and align with the sustainability and circular economy needs of the new energy vehicle sector–unlocking potential applications in 5G devices and power battery casings.

Yet despite these advantages, Mg alloys face notable barriers to wider engineering use, rooted in their hexagonal crystal structure. This structure has relatively few slip systems, leading to poor plasticity at room temperature. Conventional processing methods like rolling and extrusion often create a strong directional texture, worsening performance inconsistencies across different directions (anisotropy). Moreover, grain coarsening and a narrow processing range can trigger shear bands and cracking, further limiting the materials’ formability [1,3]. To overcome these limitations, intense plastic deformation techniques–such as equal–channel angular pressing (ECAP) and continuous shear extrusion (CSE)–have been developed. These methods refine grain size, reduce directional texture, and enhance mechanical properties through dynamic grain reorganization [3,4]. Among these approaches, shear deformation stands out for its effectiveness: by introducing high strain gradients, it activates dynamic grain reorganization, refines grains, and disrupts the strong directional texture–ultimately improving both strength and ductility [5].

Studies have shown that shear deformation triggers dynamic recrystallization (DRX), which refines grain size and improves microstructural uniformity. For example, Tian et al. [6] developed a tube continuous extrusion–shear–expansion (TCESE) process that significantly increases the strain degree in AZ31 Mg alloy tubes. This process refines grains to approximately 10 μm and enhances material hardness by reducing the intensity of the basal texture. Figure 1 presents simulated microstructures at different temperatures, revealing distinct mechanisms for shear band evolution: at room temperature (RT), shear bands mainly form from specific twin structures, while tensile twins dominate at liquid nitrogen temperature (LNT). These twins activated by shear promote the accumulation of specific dislocations, which not only redistribute local strain but also effectively stimulate DRX–driving further grain refinement and improving microstructural homogeneity [7]. Beyond grain refinement, shear deformation enables precise control of texture characteristics in Mg alloys: it weakens strong basal textures and fosters the formation of bimodal textures. Yang et al. [8] demonstrated that the asymmetric shear strain in asymmetric extrusion (ASE) alters the orientation of the basal plane, thereby enhancing both strength and ductility of AZ31 sheets. Building on this, Li et al. [9] clarified the transition in the mechanism of bimodal texture evolution: tensile twins dominate at lower temperatures, while slip effects become prevalent at higher temperatures.

While existing reviews have provided valuable broad insights into Mg alloy processing and microstructure [4,5,8,10,11,12,13,14,15], they lack a systematic analysis of the causal relationship between shear deformation parameters, specific microstructural changes, and the resulting mechanical properties. This review addresses this gap by focusing on shear–based processes as the key mechanism for regulating microstructure–moving beyond simple summaries of grain refinement to explain how shear techniques (e.g., ECAP, CSE, ultrasonic shearing) work together to improve both strength and ductility. These techniques achieve this through targeted activation of DRX, formation of bimodal textures, and construction of gradient structures. The core contribution of this review is a unified framework linking “shear process parameters–microstructural evolution–mechanical properties”, which serves as a predictive tool for designing high–performance Mg alloys (Figure 2). This represents a notable advance over previous descriptive reviews.

2. Effects of Shear Deformation on the Microstructure of Mg Alloys

2.1. Formation and Evolution Mechanisms of Shear Bands

During the plastic deformation of Mg alloys, strain localization (a direct sign of uneven strain distribution) is prone to occur [12,16]. This uneven strain not only induces shear band formation but also regulates their nucleation positions, evolution features, and propagation trends–all closely related to the alloy’s properties and deformation conditions. Nucleation usually starts from pre–existing defects or poorly oriented regions in the alloy, including grain boundaries, second–phase particle interfaces, and localized texture–softened areas. For AZ31 Mg alloy, under high–strain–rate loading, these regions become strain concentration hotspots due to uneven stress distribution, initiating shear band nucleation. This process is accompanied by the generation and accumulation of numerous twins and dislocations, laying the foundation for subsequent evolution [10].

The evolution of shear bands in Mg alloys is essentially driven by microdefects and regulated by DRX: after nucleation, the continuous buildup of twins and dislocations promotes shear band growth, causing severe plastic deformation in grains. This manifests as grain elongation and fragmentation along the shear direction, while establishing non–uniform deformation gradients. From the center of the shear band to its periphery, grains gradually transition from a fine, recrystallized microstructure to the primary matrix [13]. Importantly, shear bands are not merely damage–prone regions–their internal DRX plays a key role in optimizing Mg alloy microstructures. Rotational dynamic recrystallization (RDRX) produces fine, equiaxed grains surrounded by low–angle dislocation boundaries via subgrain rotation and grain boundary migration in deformed grains. Meanwhile, twin–induced dynamic recrystallization (TDRX) uses high–energy twin boundaries as nucleation sites, leveraging the high strain energy stored in twins to drive new grain growth. In AZ31 Mg alloy, the combined effect of these two mechanisms effectively breaks down coarse dendrites or deformation textures within the bands, resulting in a uniformly refined microstructure and alleviating localized stress concentrations [10].

The propagation of shear bands in Mg alloys is mainly determined by initial grain size, deformation temperature, strain rate, and spatial strain unevenness. Among these, grain size exerts a particularly significant influence: studies on AZ31 Mg alloy have revealed that the likelihood of shear band formation rises with increasing grain size. Specifically, coarse–grained materials tend to form narrow, localized shear bands, while fine–grained materials promote more dispersed bands–this dispersion enables more uniform strain distribution, delaying crack initiation [13]. Deformation temperature and strain rate affect propagation by altering the stress state, with high strain rates notably facilitating shear band expansion in these alloys [17]. Furthermore, under high–strain–rate deformation, insufficient heat dissipation leads to significant adiabatic heating in Mg alloy shear bands, inducing thermal softening. Research shows this thermal softening promotes DRX within grains, transforming shear bands from mere strain concentration regions into critical sites for microstructural refinement and recrystallization [18].

Overall, the formation and propagation of shear bands in Mg alloys are governed by initial grain size, deformation temperature, and strain rate, and are closely linked to specific DRX mechanisms via the Zener–Hollomon parameter, as shown in Formula (1) below.

$$\mathrm{Z}=\stackrel{·}{\mathrm{\epsilon }}\mathrm{e}\mathrm{x}\mathrm{p}(\mathrm{Q}/\mathrm{R}\mathrm{T})$$

(1)

Z denotes the temperature–compensated strain rate, where $\stackrel{·}{\mathrm{\epsilon }}$ is the strain rate, Q is the deformation activation energy, R is the universal gas constant (8.314 J/(mol·K)), and T is the absolute temperature. Coarse–grained materials tend to form localized, narrow shear bands: at high strain rates, these bands promote TDRX through strain concentration and adiabatic heating–this is particularly evident in Mg–aluminum–zinc alloys like AZ31 (Mg–3Al–1Zn), where {10–12} extension twins act as favorable nucleation sites. In contrast, fine–grained structures form more dispersed shear bands, which favor CDRX due to uniform strain distribution [13,17,18]. This relationship is primarily governed by thermomechanical conditions via the Zener–Hollomon parameter. High Z conditions–defined by strain rates above 1 s⁻¹, lower deformation temperatures, and coarse–grained microstructures–predominantly enhance TDRX through intense shear band formation and associated thermal softening. Conversely, low Z conditions (strain rates below 1 s⁻¹, deformation temperatures above 350 °C, and fine–grained structures) promote both CDRX and DDRX, driven by controlled grain boundary migration and restoration processes [17,18,19,20,21]. Alloy composition also influences these interactions. Aluminum–containing alloys such as AZ31 favor TDRX within shear bands due to enhanced twin formation, while rare–earth–containing alloys like WE43 (Mg–4Y–3RE) promote DDRX through solute–stabilized boundary mobility and particle–stimulated nucleation [5,21,22]. This interdependence between shear band features and DRX mechanisms–quantified by the Z parameter–highlights the need to optimize both thermomechanical parameters and microstructural design. Doing so enables precise control of deformation behavior and achieves targeted microstructural refinement in Mg alloys.

2.2. DRX Behavior and Grain Refinement

The key microstructural evolution mechanisms in Mg alloys under shear deformation–including DRX and stress–induced martensitic transformation–have distinct features, and their occurrence and influence depend closely on temperature, strain rate, applied deformation path, and stress state [11]. These mechanisms fall mainly into two categories: DRX subtypes (CDRX, DDRX, TDRX) (Figure 3) and stress–induced phase transformations (β–Li → α–Mg martensitic transformation, precipitate dissolution–reprecipitation, and strain–induced metastable phase formation, referred to as SIP) [23,24,25]. Among these, CDRX is the primary DRX mechanism at higher strain rates or lower temperatures (e.g., RT), achieving grain refinement through dislocation accumulation and subgrain evolution. Meanwhile, in dual–phase Mg–Li alloys subjected to severe shear (such as 71% cold rolling reduction), this transformation follows a specific orientation relationship: ${\left[111\right]}_{\mathrm{\beta }}\parallel {\left[1\overline{2}13\right]}_{\mathrm{\alpha }}$ and ${\left[110\right]}_{\mathrm{\beta }}\parallel {\left[1\overline{1}01\right]}_{\mathrm{\alpha }}$, with strain–induced metastable phases (SIPs) forming lamellar or needle–like morphologies inside the β–Li matrix [25]. In dual–phase Mg–9.5Li–1Zn alloy, activation of the $\left\{110\right\}\left[\overline{1}11\right]$ slip system in β–Li promotes more homogeneous SIP distribution, refining the β–Li matrix into cells of approximately 1.5 μm and boosting the alloy’s yield strength from 85 MPa to 227 MPa [25]. For example, Che Bo et al. observed that during continuous shear deformation of AZ31 Mg alloy, the combined effect of basal and prismatic slip supports continuous dislocation buildup, directly triggering CDRX [19].

Moreover, TDRX is another key DRX mechanism during shear deformation, speeding up recrystallization by activating twins to provide extra nucleation sites. For instance, in the combined compression–torsion deformation of AZ31 Mg alloy, {10–12} extension twins (ETs) act as DRX nuclei, promoting the formation of a refined grain structure [5]. In the novel TCESE process, TDRX, DDRX, and CDRX act sequentially alongside precipitate dissolution–reprecipitation; the newly nucleated grains undergo multi–orientation growth and metastable phase transition, ultimately forming a heterogeneous structure with excellent balanced strength and ductility (Figure 4) [26].

At elevated temperatures (180–330 °C) and low strain rates (0.01–0.1 s⁻¹), twinning further provides extra nucleation sites for DRX. For the Mg–Li–Al–Zn–Si alloy, severe shear promotes α–Mg phase spheroidization and β–Li phase DDRX, alongside the dissolution of coarse AlLi and Li₂MgAl precipitates and reprecipitation of fine Mg₂Si particles [27]. In contrast, at RT, metastable α–Mg SIPs convert to stable α–Mg grains during tensile deformation, aiding strain hardening [25]. Notably, adjusting temperature, strain rate, and deformation path allows precise control of the dominant DRX mechanism and stress–induced phase transformation behavior. This control is critical to refining grain size, optimizing phase composition, and improving the balanced strength and ductility of high–performance Mg alloys (Figure 5) [24,28,29,30].

2.3. Microstructural Evolution Induced by Twinning

Under shear deformation, the activation of {10–12} ETs and {10–11} compressive twins is a primary mechanism governing the deformation of Mg alloys. These twins play a multi–faceted role in grain reconstruction and orientation adjustment, while interacting synergistically with slip systems and recrystallization behavior. Among these twins, {10–12} ETs are the most common deformation mode in Mg alloys–they are easily activated by localized stress concentrations under shear or compressive loading at low or ambient temperatures. As strain increases, these twins expand and may eventually consume the entire parent grain, initiating the formation of a layered, homogenized microstructure and activating grain refinement. Specifically, the intersection of multiple {10–12} twin variants within a single grain forms twin–twin junctions (TTJs). These TTJs effectively hinder dislocation movement and delay recrystallization processes, thereby enhancing the material’s strength. Furthermore, the formation of {10–12}–{10–12} twin structures helps with additional grain refinement and texture modification by increasing interface density, ultimately improving the material’s overall mechanical properties [31,32,33].

In contrast, {10–11} compressive twins are activated under specific loading directions or at elevated temperatures. They work with basal and prismatic slip, enabling coordinated deformation across multiple systems and regulating the material’s overall deformation response. This is particularly noticeable in Mg alloys with non–basal textures: the simultaneous activation of multiple twinning modes can reduce the pronounced anisotropy linked to conventional textures, promoting more uniform deformation distribution under multi–axial loading [31,32].

Additionally, twin–induced crystal lattice reorientation is a key aspect of microstructural evolution. The formation of {10–12} ETs involves a ~86.3° lattice reorientation, which effectively transforms soft–oriented grains (prone to slip) into hard–oriented grains with high deformation resistance. This reorientation significantly enhances the material’s strain hardening capacity. Uniaxial compression tests on AZ31 Mg alloy have confirmed this phenomenon: as the twin volume fraction increases, a characteristic S–shaped flow curve emerges, and twin domains nearly completely replace the original parent grain structure before fracture [33,34]. Twins also exert a profound influence on DRX behavior. Thanks to their high interfacial energy and structural instability, twin boundaries serve as preferential nucleation sites for recrystallized grains. Furthermore, high–density dislocation zones near twins accumulate substantial crystallographic defects, supplying a strong driving force for CDRX [31,35]. For example, in AZ31 Mg alloy processed by the continuous bending and annealing (ECAR–CB–A) method, the initial bimodal non–subgrain structure activated numerous {10–12} twins during shear deformation. These twins significantly accelerated the nucleation and growth of recrystallized grains by facilitating grain boundary migration and dislocation rearrangement–ultimately enhancing the material’s formability and ductility (Figure 6) [32].

2.4. Distribution Behavior of Second Phases and Precipitates

Shear deformation–a typical SPD process–effectively regulates the distribution, precipitation characteristics, and stress–induced phase transformation of second–phase particles in Mg alloys. The highly localized strain fields, high shear stress, and significant strain gradients generated during this process enable deformation via dislocation climb and grain boundary slip. Simultaneously, these factors drive the dissolution of coarse precipitates (e.g., AlLi, Li₂MgAl, LPSO phases) and the β–Li → α–Mg martensitic transformation in Mg–Li–based alloys [25,27]. For the Mg–Li–Al–Zn–Si alloy, shear–induced lattice distortion enhances Mg atom diffusion, promoting the dissolution of coarse AlLi particles at 180–230 °C [27]. This process drives the directional migration and redistribution of second–phase particles, while inducing irreversible morphological changes, uniform particle refinement, and stress–driven reprecipitation of fine metastable precipitates. In the Mg–9.5Li–1Zn alloy, SIPs reprecipitate as lamellar or granular particles within the β–Li matrix; in the Mg–Li–Al–Zn–Si alloy, fine Mg₂Si particles form at grain boundaries after shear deformation [25,27]. These modifications–including phase composition changes (β–Li → α–Mg, stable → metastable precipitates), precipitate refinement, and interfacial energy rearrangement–alter the bonding state and energy distribution between particles and the matrix. Consequently, they directly govern recrystallization and phase transformation kinetics by regulating grain boundary migration rates, dislocation annihilation efficiency, and precipitate–matrix interface stability [25,27].

Studies show that shear deformation significantly improves the dispersion of second–phase particles and reduces their average size. As illustrated in Figure 7, N. FAKHAR et al.’s multi–pass DECLE processing effectively refines the microstructure of ZK60 alloy sheets: the grain size decreases from 68 μm (annealed state) to 6.0 μm and 5.2 μm after 3 and 5 passes followed by extrusion, respectively [36]. For Mg–9.5Li–1Zn alloy sheets undergoing 71% multi–directional rolling–induced shear deformation, the β–Li matrix is fragmented into ~1.5 μm grains [25]. Under a shear rate of 0.01 s⁻¹ and temperature of 270 °C, Mg–Li–Al–Zn–Si alloy sheets form uniformly dispersed fine Mg₂Si particles [27]. During the eccentric extrusion of Mg–Nd–Zn–Zr alloys, the shear component from the asymmetric stress field continuously acts on coarse second–phase particles, triggering plastic fragmentation and dispersion along the shear direction–ultimately forming uniformly distributed refined particles [37]. Notably, for Mg–based high–entropy alloys with an atomic size mismatch of 12–27%, the uniform dispersion of second–phase particles induced by shear deformation is more targeted. Unlike conventional Mg alloys–where such dispersion inhibits abnormal grain growth and enhances strength by shortening diffusion paths–this effect also effectively reduces processing inhomogeneity caused by uneven atomic diffusion. Additionally, it improves the stability of the amorphous phase, increasing its glass transition temperature by 15 K [38].

From the perspective of morphological evolution, second–phase particles initially agglomerated or coarse in the as–cast condition transform into elongated, fragmented, or dispersed forms after shear deformation. This morphological change arises from interfacial stress concentrations caused by the plastic mismatch between particles and the matrix. Once the applied stress exceeds the particles’ fracture strength, it causes them to fragment and rearrange in an oriented manner along the shear direction. This phenomenon is most evident in SPD processes like equal–channel angular extrusion and cyclic shear extrusion. In these processes, elongated particles tend to align parallel to the shear direction, forming highly oriented fibrous structures. These structures facilitate efficient load transfer along the shear direction, allowing the design of materials with controlled anisotropy and balanced strength and ductility along specific orientations.

Second–phase particles regulate microstructural evolution through two main mechanisms. First, particles hinder grain boundary migration via the Zener pinning effect–where the pinning force is proportional to the particle volume fraction but inversely proportional to particle size. This mechanism effectively stabilizes fine–grained structures and retards recrystallization. Conversely, particles act as heterogeneous nucleation sites, promoting the formation of recrystallized grains. This effect is especially notable in local high–strain regions like shear bands–areas with extremely high dislocation densities and concentrated energy. These conditions facilitate the formation of recrystallization nuclei on particle surfaces, significantly increasing nucleation rates. Thus, by regulating the spatial distribution and morphology of second–phase particles, shear deformation enables bidirectional control of microstructural evolution: it maintains fine–grain stability through physical blocking, while accelerating microstructural homogenization by inducing recrystallization. In summary, shear deformation exerts systematic effects on second–phase particles and precipitates by adjusting the applied stress–strain conditions. It not only reshapes particle morphology and spatial distribution but also plays a key role in regulating grain boundary migration and recrystallization kinetics. Future research should integrate in situ characterization techniques with multiscale simulations to clarify the interaction mechanisms between particles, dislocations, and grain boundaries, laying a theoretical foundation for the precise optimization of shear deformation processes.

2.5. Summary and Synthesis

Shear deformation–induced microstructural evolution encompasses three core aspects: grain refinement through multiple pathways, including DRX subtypes–CDRX generates subgrains in AZ31 [19], TDRX nucleates at {10–12} extension twin boundaries [5], DDRX forms heterogeneous structures [26]–and stress–induced phase transformations such as β–Li → α–Mg martensitic transformation and strain–induced metastable phase formation in Mg–Li alloys [25,27]; texture weakening via basal plane reorientation [8], bimodal texture development through twin–activated DRX [9], and ~86.3° lattice reorientation induced by {10–12} extension twins [33,34]; and second–phase regulation involving particle fragmentation, dispersion, dissolution, and reprecipitation to enhance Zener pinning and recrystallization nucleation, optimize phase composition–for example, coarse AlLi transforms to fine Mg₂Si in Mg–Li–Al–Zn–Si [27], and SIP reprecipitates in Mg–9.5Li–1Zn [25]–and alleviate processing inhomogeneity in Mg–based high–entropy alloys through targeted particle dispersion [38]. However, there remains debate about the dominant grain refinement mechanism. For instance, TDRX dominates in AZ31 under combined compression–torsion [5], while CDRX prevails in AZ31 during continuous shear [19]; in dual–phase Mg–Li alloys under severe shear, stress–induced phase transformations take the lead [25]. These inconsistencies arise from differences in deformation modes, strain paths, local stress states, and alloy systems such as single–phase AZ31 versus dual–phase Mg–Li, indicating the dominant mechanism is highly sensitive to processing parameters and material compositions rather than being universally applicable. Additionally, the dual role of second–phase particles–both boundary pinning and promoting recrystallization–lacks quantitative characterization. Their dissolution–reprecipitation kinetics under shear are not fully understood, and key unresolved issues include predictive models linking shear parameters to DRX kinetics, texture evolution, and stress–induced phase transformation behavior, as well as real–time understanding of twin–dislocation–boundary–particle interactions. Addressing these gaps requires integrating in situ characterization techniques with multiscale modeling [25,27,35].

3. Effect of Shear Deformation on Texture Evolution

3.1. Subsection Mechanisms of Weakening in Basal Plane Texture

Understanding the mechanism of basal texture weakening is a core research focus in the field of Mg alloy shear deformation. Strong basal textures lead to significant mechanical anisotropy at RT, severely restricting the formability and engineering applications of Mg alloys [39]. Conventional processes like extrusion and rolling typically involve unidirectional stress and often produce strong basal textures, where the (0002) basal planes align parallel to the main deformation direction. This orientation greatly restricts the activation of slip systems, reducing the material’s overall plastic deformability [40].

Numerous studies have shown that shear deformation weakens crystallographic textures more effectively than conventional methods. Conventional rolling or extrusion mostly activate only one slip system and usually reduce texture strength by less than 30% [40], while shear deformation can achieve a 40–60% reduction [9,32] via the combined action of twinning and DRX. It also facilitates the formation of a bimodal texture, which is unattainable with conventional processes. Of these mechanisms, {10–12} tensile twinning is a key factor: it is easily activated at lower temperatures, causing substantial grain reorientation that disrupts texture uniformity [9]. As shear strain accumulates in AZ31 Mg alloy, a large number of multi–oriented {10–12} twin bands form; their interaction promotes the formation of an alternating banded structure, as illustrated in Figure 8. This alternating structure is typically associated with a bimodal texture, which reduces overall texture intensity by dispersing preferred crystallographic orientations [9]. Furthermore, these twinning–induced microdomains not only directly weaken the basal texture by changing local crystallographic orientations but also serve as preferential nucleation sites for subsequent DRX. The random growth of these recrystallized grains further enhances texture randomization, reinforcing the weakening effect of shear deformation on basal textures.

Shear–induced continuous and discontinuous DRX are key mechanisms for texture adjustment. CDRX promotes grain refinement through subgrain growth and coalescence, while DDRX supports nucleation and growth via grain boundary migration. Both processes refine the microstructure by replacing coarse initial grains with fine, equiaxed ones, thereby altering texture distribution [41]. During multi–pass shear processing, the material undergoes repeated shear stresses from different directions, causing continuous grain reorientation and a gradual reduction in crystallographic preference. This process effectively weakens the initial strong basal texture [42]. For example, ECAP achieves uniform dispersion of grain orientations and weakens basal plane alignment via well–designed multi–pass processing routes, thereby improving mechanical properties and isotropy [43]. Furthermore, the activation of non–basal slip systems contributes significantly to texture weakening. Although pure Mg has a high critical resolved shear stress (CRSS) for non–basal slip, elevated temperatures or specific stress states help activate them. Increased activity of these slip systems enables more complex deformation coordination, reducing the dominance of basal slip [44,45,46]. For WE43 Mg alloy, texture weakening combined with a high processing temperature of 523 K maximized the prismatic slip activation fraction to 33.1% and promoted dynamic recrystallization, which accounts for 37.7% of grains in the interior region. This microstructural evolution resulted in a remarkable increase in the bending limit to 29.5 mm from a reference value of 5.9 mm. This reference value was measured on the RDRB sample, which is the RD sample tested under room–temperature bending at 298 K. It also reduced the bending yield strength to 315 MPa, ultimately improving bending performance by 64.7% compared to the RDRB benchmark.

3.2. Texture Formation and Evolution of Bimodal Textures

Conversely, when processed in liquid nitrogen, AZ31 plates treated with equal channel angular rolling followed by ECAR–CB–A undergo deformation dominated by {10–12} ETs. This promotes grain reorientation where the c–axis is normal to the transverse direction and leads to the formation of a transverse texture component. As the number of rolling passes increases, {10–12}–{10–12} double twins (DTWs) become active in the deformation process. Their involvement ultimately eliminates the transverse texture component, highlighting the impact of different twin types and their interactions [32].

Moreover, the evolution of bimodal texture arises from multiple mechanisms: basal <a> slip becomes the primary mechanism for accommodating plastic strain after the double–twin mechanism initiates [32]. As shown in Figure 9, during the initial stage of shear deformation, {10–12} extension twins nucleate extensively and gradually form banded structures as strain increases. Through the combined effect of CDRX and DDRX, a bimodal texture composed of T1 and T2 components develops. The T1 texture features basal planes tilted 30–45° from the extrusion direction (ED) and is dominated by basal <a> slip, while the T2 texture has basal planes nearly parallel to ED and is associated with <c + a> slip. These two texture components keep a steady interplanar angle of 40–43°, with their relative intensities changing dynamically with temperature: T2 dominates at 250 °C, while a more balanced distribution occurs at 340 °C [9]. This bimodal structure significantly enhances mechanical properties. At 340 °C, the tension–compression yield ratio rises from 0.72 to 0.87, elongation improves by 50% (from 12% to 18%), and yield strength anisotropy decreases from 45 MPa to 22 MPa. These property improvements stem from the coordinated activation of slip between T1 and T2 components: T1 promotes the activation of multiple slip systems to accommodate deformation, while T2 alleviates stress concentration through <c + a> slip. Along with refined DRX grains that average 5.7 μm, this coordinated mechanism achieves quantitative improvements in both isotropy and ductility [9].

Furthermore, rare earth elements (Gd, Y, Nd) effectively regulate basal texture development in Mg alloys by greatly lowering the critical resolved shear stress ratio between non–basal and basal slip systems from 70–100 to 5–10. This reduction is achieved through solute segregation and adjusting interface energy, which activates extensive non–basal slip and systematically reorients the c–axis by 30–45° from the extrusion direction. For the Mg–2Gd–0.4Zr alloy, this results in approximately 35° c–axis tilting, with basal planes oriented 45–55° relative to the extrusion direction. Zirconium further enhances texture modification via Zener pinning and particle–stimulated nucleation. Zirconium–rich particles (50–200 nm) preferentially segregate in fine–grained regions, with local concentrations exceeding 1.62 wt% compared to the matrix average of 0.35 wt%. This microstructural change promotes complete dynamic recrystallization, achieving a nearly 100% DRX fraction. The synergistic interaction between rare earth elements and zirconium reduces the maximum basal texture intensity from >10 m.r.d. in conventional alloys to 3–5 m.r.d. This improvement substantially boosts room–temperature elongation to over 50% while keeping yield strength anisotropy below 20 MPa [22].

3.3. Relationship Between Crystal Orientation Rearrangement and Shear Direction

Shear direction is a critical factor governing grain orientation patterns, with its influence particularly pronounced due to the anisotropic nature of the HCP structure and variations in slip system activity [22,47,48,49,50]. The CRSS for slip systems in HCP Mg alloys shows a clear hierarchy: basal <a> slip has the lowest value (~45 MPa), followed by prismatic <a> slip (~110 MPa), while pyramidal <c + a> slip has the highest (~170 MPa) [50]. Meanwhile, shear deformation can effectively adjust these CRSS values through mechanisms such as strain gradients and texture adjustment. For example, shear–induced tilting of the basal plane has been shown to reduce the CRSS for pyramidal <c + a> slip by 20–30% [51]. Under applied shear stress, grains undergo lattice rotation to form specific angles between the basal plane and the shear direction, adapting to the applied deformation. This rotation meets local strain demands by facilitating the activation of slip or twinning systems with lower CRSS, thereby enhancing strain accommodation. This reorientation behavior exhibits both path dependency and strain sensitivity, as it is jointly controlled by the applied shear path and accumulated strain. During processes like ECAP or simple shear, the accumulation of shear strain causes basal planes that were initially parallel to the processing direction to gradually rotate away from their initial orientation, aligning with the shear direction to varying degrees. The misorientation angle between the basal plane and the shear direction increases approximately linearly with strain until reaching a stable state in the shear stress field [52]. This process not only promotes the activation of non–basal slip systems to mitigate plastic anisotropy from a strong basal texture but also creates favorable orientation conditions for {10–12} tensile twinning. When this misorientation angle approaches the critical value for twin nucleation, localized stress concentrations can more easily overcome the energy barrier for twin formation, further enhancing the material’s strain accommodation capacity.

Following the introduction of the Shear–Assisted Twin Orientation Regulation (SATOR) process in AZ31 Mg alloy, pre–twinned samples developed a shear texture oriented near 45° during subsequent shear deformation. Optimizing the angle between the basal plane and the loading direction significantly improved the material’s ductility. VPSC simulation results (

Table 1. Material parameters used for the VPSC simulation of the wrought AZ31 Mg alloy rod [21].

	${\mathbf{\tau }}_0$(MPa)				${\mathbf{\tau }}_1$(MPa)
Mode	AR Type	Pt	AR–SA Type	3PT–SA Type	AR Type	Pt	AR–SA Type	3PT–SA Type
plinth	45	21	58	52	0	0	0	0
prism	110	70	128	130	43	30	30	50
pyramid	170	190	170	170	85	85	85	85
{1012} twin	28	19	75	65	0	0	0	0
{1011} twin	250	250	250	250	20	20	20	20
	${\mathbf{\theta }}_0$(MPa)				${\mathbf{\theta }}_1$(MPa)
Mode	AR Type	Pt	AR–SA Type	3PT–SA Type	AR Type	Pt	AR–SA Type	3PT–SA Type
plinth	220	220	220	180	140	140	140	140
prism	500	0	217	80	0	1700	20	40
pyramid	550	550	220	550	100	100	100	100
{1012} twin	0	0	0	0	0	0	0	0
{1011} twin	300	300	300	300	10	10	10	10

) further indicate that this orientation rearrangement reduces dislocation motion resistance during basal slip, enhancing basal slip activity by nearly 3.8 times. This provides additional mobile slip systems for subsequent deformation, leading to effective improvement in plasticity [50].

3.4. Summary and Synthesis

Shear deformation effectively weakens the basal texture in Mg alloys through two key mechanisms: {10–12} extension twinning induces crystallographic reorientation to disrupt initial texture components [9,32], and dynamic recrystallization generates randomly oriented grains [41,43]. Bimodal texture formation is another important result, achieved either via deformation–induced twin variants [32] or by adding rare earth elements to activate non–basal slip systems [22]. Nevertheless, there are significant discrepancies in existing studies regarding the dominant mechanisms. Li et al. [9] attributed texture weakening mainly to {10–12} extension twinning at lower temperatures, while Xie et al. [21] confirmed that slip–dominated mechanisms are dominant at elevated temperatures in WE43 alloy. The role of rare earth elements is also controversial: Zheng et al. [22] emphasized particle pinning effects in Mg–Gd–Zr alloys, while Lu et al. [50] highlighted solute–mediated changes to slip systems in AZ31. These variations underscore the context–sensitive nature of texture evolution, which arises from differences in processing parameters and material compositions. Current research limitations include the absence of quantitative predictive frameworks for texture evolution under complex shear paths, insufficient understanding of bimodal texture stability during subsequent processing, and limited insight into how shear direction interacts with initial texture to determine final orientation distributions. Future research should focus on combining experimental and computational methods to build predictive capabilities for texture design.

4. Shear Deformation Processes and Their Recent Technological Advances

4.1. Shear Deformation via ECAP

As a well–established SPD technique, ECAP employs multi–pass shear deformation to introduce a high cumulative shear strain without altering the specimen’s cross–sectional geometry, thereby significantly refining the microstructure and enhancing the overall mechanical properties of Mg alloys (Figure 10) [53].

For microstructural refinement, a Mg–3.7Al–0.7Zn–0.8Sn–0.4Mn alloy treated with four passes of ECAP through Route A achieved the best room–temperature mechanical properties: yield strength (YS) of ~225 MPa, ultimate tensile strength (UTS) of ~312 MPa, and fracture elongation (EL) of ~31.9%. This balanced improvement mainly came from the combined action of multiple mechanisms: DRX–formed submicron–scale grains strengthened the matrix following the Hall–Petch relationship; dispersion of nanoscale second–phase particles enhanced strength via the Orowan dislocation pinning mechanism; and the weakened basal texture effectively boosted plastic deformability [54]. Multi–pass ECAP at 373 K and 2.5 × 10⁻³ s⁻¹ refined Mg–9Li alloy grains from 400 μm to 360 nm via intensive shear deformation. Generated high–density dislocations developed into subgrain boundaries, dividing grains and activating dynamic recrystallization through thermomechanical coupling. This process promoted the nucleation of strain–free grains along original boundaries and shear bands, eventually forming a homogeneous ultrafine equiaxed microstructure [55] (Figure 11).

For process optimization, combining multi–pass ECAP with pre–deformation or post–processing creates a combined effect that further enhances the overall mechanical properties of Mg alloys. After linear multi–pass ECAP, the dual–fiber texture (<10–10> and <11–20>) initially formed in AZ31 Mg alloy by conventional extrusion was significantly broken down. Texture intensity decreased by over 60%, effectively reducing the material’s anisotropy [56]. This texture weakening comes from the randomization of grain orientations induced by multi–pass shearing. Each pass rotates grains according to the shear direction, and the combined effect of multiple passes gradually disperses orientations across different grains. At the same time, new grains formed through DRX have more random orientations, which collectively reduce the concentration of preferred orientations in the basal texture. Furthermore, different processing routes have different impacts on texture evolution. Route Bc involves a 90° rotation of the sample around its longitudinal axis between passes, generating a more complex three–dimensional shear stress state. This results in better grain refinement and texture control compared to Route A [54].

For example, when a combined process of pre–deformation and indirect extrusion is applied to AZ31 Mg alloy, a uniform fine–grained microstructure is achieved after four passes. However, exceeding this optimal number of passes can cause grain coarsening, which is attributed to secondary recrystallization driven by thermal energy accumulation during excessive deformation [56]. Moreover, key ECAP parameters–including processing temperature, channel angle, and friction conditions–have significant effects that vary by alloy on microstructural evolution and property control. These parameters require careful adjustment for each specific alloy system [57]. In recent years, as understanding of ECAP mechanisms has deepened, its applications have expanded considerably. Beyond traditional cast Mg alloys, ECAP technology has been successfully explored for processing powder metallurgy materials and metal matrix composites, demonstrating excellent versatility and broad extension potential [54,55,56,57,58].

4.2. CSE and Advanced Composite Processing Techniques

CSE is an emerging SPD technique that applies intense shear forces to materials during continuous flow through the combined effect of high cumulative strain and a quasi–uniform deformation field, its schematic is shown in Figure 12a. The process achieves a cumulative strain exceeding 2.0, much higher than that of conventional rolling, while maintaining a strain variation of only 5–10%–this value is significantly lower than the 15–20% reported for ECAP [59]. It effectively activates DRX and promotes grain refinement, simultaneously enhancing both the strength and ductility of Mg alloys. This provides crucial technological support for meeting material requirements in aerospace and automotive lightweight applications, leading to the development of derivative processes such as dual–path continuous shear (DE–CS) and equal–angle extrusion–shear (ES). Technically, CSE uses specially designed channel dies to apply intense shear stress to metal billets during continuous flow at elevated temperatures. This simultaneously induces massive dislocation proliferation and entanglement, accumulating a strain exceeding 2.0 that is far higher than what conventional rolling can achieve. It also triggers plastic grain fragmentation and continuous DRX. CSE enables grain refinement in Mg alloys from the micrometer to submicrometer scale. For AZ31 alloy sheets treated with multi–stage shear deformation at 340 °C, the average grain size decreases from 6.33 μm (conventional extrusion) to 3.39 μm after CSE–II processing. This microstructural refinement stems from intense shear strain, which promotes dynamic recrystallization, weakens basal texture through approximately 15° lattice rotation, and enhances dislocation strengthening mechanisms [60]. After CSE processing at 330, 370, and 410 °C, the recrystallized volume fraction increases from 65% to 90%. However, accelerated grain boundary migration at elevated temperatures causes slight grain growth in localized regions, leading to a decrease in Vickers hardness from 77 HV to 72 HV (Figure 13) [61].

This degradation highlights the importance of precise thermal control, as temperature fluctuations beyond ±5 °C induce microstructural heterogeneity. Furthermore, the complex multi–channel die geometries generate substantial stress concentrations, requiring tool steels with strength exceeding 2 GPa. These demanding conditions accelerate die wear through both abrasive particle erosion and material adhesion. Current mitigation methods include PVD coatings (CrN/TiB₂) to enhance wear resistance and multi–zone heating systems to improve temperature uniformity. However, fundamental limitations remain in predictive wear modeling and creating energy–efficient thermal management systems that can simultaneously optimize microstructural refinement and production efficiency.

To meet complex forming requirements, derivative CSE processes further expand technological boundaries. The ES process combines the high–strain deformation capability of ECAP with the continuous processing advantages of CSE (Figure 12b). It uses an asymmetrical channel design to achieve large plastic deformation in bars with diameters ranging from 10 to 50 mm. Intense shear deformation in this process effectively breaks down the strong basal texture, reducing its intensity in the Mg–1.5Zn–0.5Zr–0.5Sr alloy from 48.5 (direct extrusion) to 6.3 [62]. This texture weakening is accompanied by microstructural refinement by promoting discontinuous dynamic recrystallization (DDRX), which increases the proportion of high–angle grain boundaries while reducing grain size. The combined effects of texture randomization and grain refinement promote the activation of non–basal slip systems, allowing synergistic deformation coordination with {10–12} tensile twinning to accommodate plastic strain. In comparison, the direct extrusion and continuous shear deformation (DECS) process uses a symmetrical dual–channel die to apply bidirectional shear during sheet processing [59], establishing a three–dimensional stress field. This reduces regional strain variations from 15–20% to below 5%. It not only breaks down the primary texture concentration and activates multiple slip systems and twinning but also significantly reduces mechanical anisotropy. A comparative analysis of the processing parameters, mechanical properties, grain size, and texture strength of SPD techniques is shown in

Table 2. Comparative Analysis of Processing Parameters, Mechanical Properties, Grain size and Texture intensity by SPD Techniques [54,55,56,59,60,62].

SPD Technique	Types	Processing Parameters		Mechanical Properties			Grain Size/μm	Texture Intensity
		Extrusion Speed	Extrusion Temperatures	YS/MPa	UST/MPa	EL/%		Before	After
ECAP	Mg–3.7Al–0.7Zn–0.8Sn–0.4Mn			225	312	31.9
	Mg–9Li	–	373 K	23.6	14.3	98.4	0.36	–	–
	AZ31		370 °C	–	250	9.4	4	5.06	1.92
CES	AZ31	0.1 mm/s	340 °C	225.2	305.8	20.17	16.12	25.77	10.7
ES	Mg–1.5Zn–0.5Zr–0.5Sr	2 mm/s	200 °C	206	254	21.4	2.63	48.5	6.3
DECS	Mg–AZ31B/Al6030	3 mm/s	330 °C	–	251.4	23.5%	5.09	–	16.97

:

4.3. Ultrasonic Vibration–Assisted Shear Deformation

Ultrasonic vibration–assisted shear deformation is an emerging processing technique that has gained significant attention in recent years. It introduces high strain rates and localized heating at ambient temperatures, effectively promoting shear band formation and microstructural evolution. On one hand, using the acoustic softening effect, ultrasonic periodic stresses disturb dislocation motion, reducing the flow stress of AZ31 Mg alloy by 20–30% while converting 20–50 kHz high–frequency mechanical energy into localized energy input. The transient temperature rise and stress pulses accelerate dislocation multiplication, leading to rapid localized accumulation of strain energy within the material. When the strain energy exceeds a critical threshold, shear band nucleation initiates. This energy accumulation further disrupts grain boundary structures, promoting the initiation and development of DRX. As shown in Figure 14, ultrasonic treatment (32 μm amplitude, 20 kHz) refined AlMg₃ alloy grains from ~270 μm to ~1.52 μm through enhanced shear band formation and grain rotation. This processing increased high–angle grain boundaries to ~21% and raised the average misorientation from 6.5° to 13.8°, as subgrain boundaries transformed via dislocation rearrangement [63]. The acoustic softening effect reduced flow stress, promoting dislocation accumulation and providing thermodynamic conditions for dynamic recrystallization. This ultimately increased microhardness from ~68 HV to 116 HV. Furthermore, ultrasonic vibration generates localized transient heating of 100–150 °C through its inherent thermal effects without significantly increasing the bulk workpiece temperature. This localized temperature rise reduces flow stress, facilitates grain reorientation, and weakens basal texture anisotropy. Experimental results confirm that ultrasonic assistance offers dual advantages of thermal softening and interfacial friction reduction during single–point incremental forming of AZ31B Mg alloy. Among the tested frequencies (0–100 kHz), the 20 kHz condition delivered optimal performance by minimizing both maximum shear stress and thinning rate. At this frequency, shear stress distributed uniformly along the forming edge, with the thinning rate stabilizing at approximately 36%. Ultrasonic vibration also significantly improved interfacial friction conditions: the friction coefficient (μ) showed a strong positive correlation with maximum shear stress (R² > 0.9) and a linear relationship with thinning rate, collectively enhancing overall forming performance [64].

4.4. Shear–Induced Evolution of Microstructural Gradient Architectures

During shear deformation, inherent gradients in shear stress and accumulated strain are core manifestations of strain distribution heterogeneity that exist in different material regions. Precise control of gradient structural characteristics such as grain size gradient can be achieved by regulating shear deformation parameters, which directly adjust the degree and spatial distribution of strain heterogeneity (

Table 3. Comparative Analysis of Shear Deformation Techniques for Mg alloys [51,54,55,56,57,58,59,60,61,62,63,64,65,66,67].

Process	Key Advantages	Key Limitations/Challenges	Scalability & Industrial Potential
ECAP	Exceptional grain refinement (to sub–micron level) Effective texture weakening & randomization via multi–pass routes Proven for a wide range of Mg alloys including cast, PM, MMCs	Inherently a batch process; low production efficiency High tooling loads and strict die strength requirements Sample size and geometry constraints (typically billets) Potential for surface defects in multi–pass processing	Primarily a laboratory–scale tool for fundamental research and high–value, small components. Continuous variants (like Conform–ECAP) are emerging but not yet mature for Mg alloys.
CES& Derivatives (ES, DE–CS)	Continuous, semi–continuous processing capability Suitable for long products (rods, tubes, sheets) Good balance of grain refinement and texture control Higher production efficiency than ECAP	Complex die design and manufacturing Significant challenges in thermal management and strain distribution High tooling wear in industrial–scale operation Limited geometry flexibility of final product	Shows the greatest immediate potential for scaling to industrial production of extruded Mg profiles and sheets, particularly for automotive applications.
Ultrasonic Vibration–Assisted Shearing	Reduces flow stress (20–30%) and forming loads Improves surface quality and formability Can be integrated with other processes (e.g., SPIF) Promotes DRX and microstructural refinement	Energy–intensive; complex system integration Limited effective volume of treatment (localized effect) Challenges in maintaining consistent amplitude/frequency at scale Lack of data on long–term tooling and transducer reliability	Highly promising for niche, high–precision forming operations (e.g., micro–forming, incremental forming) rather than bulk material production.
TRSR & Asymmetric Shearing	Effective for creating beneficial gradient structures Can process wide sheets and plates Improves ductility and fatigue resistance through microstructural heterogeneity Compatible with existing rolling infrastructure	Requires precise control of roll speeds and temperatures Can introduce undesirable residual stresses if not optimized Anisotropy control is more complex than with symmetric processes	TRSR is already an industrial process for steels and Ti alloys; adaptation for Mg alloys is feasible for sheet and plate production.

). Established techniques for constructing such architectures include three–roll skew rolling (TRSR) and asymmetric shearing. Figure 15 presents the schematic diagrams of the TRSR and asymmetric shear processes, and these techniques have proven highly effective in controlling microstructural gradients in Mg alloys through precise control of the shear stress field.

TRSR is a typical SPD process dominated by shear strain. By controlling processing parameters such as skew angle, reduction ratio, and rolling temperature, TRSR induces nanocrystalline or ultrafine–grained regions in the surface layer of Mg alloys while retaining relatively coarse original grains in the core, forming a distinct surface–to–core gradient microstructure [65]. This gradient architecture simultaneously enhances the material’s hardness and strength while significantly improving its ductility and fatigue resistance. The surface layer of a Mg–Gd–Y–Zn–Zr alloy processed under optimized conditions exhibits a gradient heterogeneous structure with a thickness of up to 2.4 mm, accompanied by the formation of numerous nanoscale precipitates (NM particles). This heterogeneous structure originates from the layer–by–layer breakdown of shear–induced layered cellular phases during rolling, combined with the dissolution and reprecipitation of long–period stacking ordered (LPSO) phases. High shear stresses fracture the LPSO phases along grain boundaries; the fragments are then dynamically recrystallized into dispersed nanoparticles. These nanoparticles effectively pin nanocrystalline grain boundaries through the Zener pinning effect, inhibiting grain growth, further refining the microstructure, and enhancing its structural stability [51].

Asymmetric shearing, which is also used to construct gradient microstructures in Mg alloys, introduces shear stresses through differential speeds between upper and lower rolls or asymmetrical die geometries [66]. This generates significant variations in deformation depth through the material thickness, resulting in a macroscopic gradient in grain size distribution from the surface to the core. This process not only promotes gradient DRX through high shear strain but also reduces the CRSS for <c + a> slip on pyramidal planes through asymmetric stress states. Compared to conventional symmetric deformation, asymmetric shearing reduces the CRSS for pyramidal <c + a> slip by approximately 20–30%, significantly increasing its activation chance. This effect disrupts the orientation concentration of the strong basal texture, inducing a multi–peak texture distribution that enhances the material’s overall plasticity [51]. Experimental data indicate that Mg alloys processed by asymmetric shearing exhibit higher dislocation densities and more complex grain boundary networks. These microstructural features provide additional obstacles for dislocation storage and motion, thereby simultaneously enhancing both the material’s strength and work–hardening capacity.

Advanced shear deformation techniques have distinct microstructural control capabilities: ECAP creates ultrafine–grained structures through multi–pass processing, with Route Bc especially good at texture randomization [54,56]; CSE and its derivatives achieve significant grain refinement while keeping processing continuous, though challenges remain in temperature control and microstructural homogeneity [60,62]; ultrasonic vibration–assisted methods reduce flow stress by 20–30% via acoustic softening to enhance formability [63,64]; and gradient structure fabrication techniques like TRSR successfully create surface–to–core microstructural gradients that simultaneously improve strength and ductility [51,65]. There are significant discrepancies regarding optimal processing parameters across these advanced techniques. The ideal number of ECAP passes is still controversial: some studies report ongoing grain refinement with more passes [54], while others find grain coarsening when exceeding four passes [56]. Similarly, the effectiveness of ultrasonic vibration assistance varies greatly with frequency and amplitude parameters, and there is limited consensus on optimal values [63,64]. Major obstacles stand in the way of industrial application of these processes. This is particularly true for CSE and TRSR, where tooling design and thermal management systems need significant improvement. The economic feasibility of ultrasonic–assisted processes at the industrial scale has not been confirmed, while predictive models for gradient structure control in asymmetric shearing are still underdeveloped. Future research should overcome these limitations by integrating process–structure–performance modeling and conducting pilot–scale demonstrations. This will help close the gap between laboratory innovations and industrial implementation.

5. Effects of Shear Deformation on Mechanical Properties

5.1. Strength Enhancement Mechanisms Through Grain Refinement and Dislocation Hardening

Shear deformation significantly enhances the mechanical properties of Mg alloys through two combined mechanisms: grain refinement strengthening and dislocation hardening. This is achieved via high strain accumulation and localized stress concentration, which together promote grain fragmentation, DRX, and substantial dislocation multiplication [21,24,68]. For AZ31 Mg alloy processed through Mg–Al composite billet co–extrusion, {10–12} extension twins effectively divide the original grains (13.5–15.9 μm) and provide favorable sites for DRX nucleation. This results in grain refinement to 6.3–7.6 μm, a reduction of approximately 50% in grain size. The resulting heterogeneous microstructure, composed of DRX grains and UDRX regions, introduces strength gradients of 50–80 MPa that promote more uniform stress distribution [68]. This strengthening behavior follows the Hall–Petch relationship. For example, AZ31B alloy subjected to hot compression under a Z of 9 × 10¹⁴ s⁻¹ forms DRX grains as small as 2.6 μm, increasing yield strength from 60 MPa to 225 MPa [24]. Regarding dislocation hardening, WE43 Mg alloy undergoing high–temperature bending exhibits a GND density of 5.24 × 10¹⁴ m⁻² in deformed grains, which decreases to 2.23 × 10¹⁴ m⁻² in DRX grains, reducing stress concentration [21]. Concurrently, in AZ31 alloy, shear stress promotes dislocation multiplication and entanglement. It activates non–basal slip systems (prismatic <a> and pyramidal <c + a> slip) that interact with basal dislocations to form three–dimensional networks, further raising flow stress levels [68].

LPSO phases provide additional strengthening mechanisms during shear deformation of alloys such as Mg–Gd–Y–Zn–Zr [51,69]. These fibrous LPSO phases (15–25 vol%, 0.5–2 μm length, 50–200 nm width) undergo fragmentation and dissolution during deformation. They provide abundant DRX nucleation sites that refine grains from 126.3 μm to 2.03 μm or smaller, while simultaneously enhancing dislocation bypass stress by 80–100 MPa through Orowan strengthening [51]. The coherent interfaces between LPSO phases and the Mg matrix generate significant stress concentrations that activate <c + a> dislocations. Dislocation pile–ups require additional stresses of 120–150 MPa to overcome [51]. In alloys with Gd content ≥ 4%, dense LPSO precipitation in UDRX grains inhibits twinning and restricts crack propagation through kink band formation. After shear deformation, these alloys achieve a maximum hardness of 127.67 HV (a 43.5% increase over the homogenized state) and yield strengths exceeding 320 MPa, reaching 350 MPa in some systems. In contrast, GZ21M alloy achieves 40% elongation due to the combined effects of rare–earth texture modification and grain refinement [69].

5.2. Improvement in Ductility and Formability

At RT, Mg alloys are limited by their HCP crystal structure, which inherently restricts the number of available slip systems. This often leads to poor plasticity and significant anisotropy. Shear deformation has been shown to effectively weaken the basal texture by adjusting the stress field distribution, reducing the activation threshold for non–basal slip systems. It also promotes a combined effect between DRX and twinning, enhancing uniform elongation and fracture toughness while effectively mitigating anisotropy. This section outlines three primary mechanisms for improving the room–temperature plasticity of Mg alloys.

Strong basal texture is a key cause of anisotropy and low ductility in Mg alloys. Conventional extrusion and rolling processes tend to align grains with the basal planes, severely limiting the activation of non–basal slip systems. Introducing directional strain through methods like asymmetric angular rolling (AAR) and differential temperature rolling (DTR) disrupts this textural symmetry, promotes grain reorientation, and weakens the basal texture. For instance, in AZ31 Mg alloy, the AAR process imposes bidirectional asymmetric shear stresses that tilt the basal planes by 30–45° away from the shear direction, changing the grain orientation distribution from concentrated to dispersed. This concurrently activates non–basal slip systems such as prismatic <a> slip, raising the material’s uniform elongation to 17.9–18.5% and significantly reducing the anisotropy of yield strength, tensile strength, and elongation [70]. Similarly, DTR effectively weakens the basal texture. Under high differential temperature processing (HDTR) conditions, the applied high strain promotes the formation of more twinning systems, including contraction twins (CTWs) and DTWs, greatly increasing the recrystallized fraction. After annealing, the resulting elongation reached ~33.6% (Figure 16) [71], further validating the critical role of texture weakening in enhancing plasticity.

Twinning activation is another core mechanism for enhancing the ductility of Mg alloys, though its efficacy greatly depends on precise control of twin type. {10–12} extension twins (TTWs) have a low CRSS and easily activate under applied stress. These twins not only segment coarse grains by forming twin boundaries but also generate high–energy zones at these interfaces. This provides ample nucleation sites for DRX, leading to a combined effect that integrates grain refinement, texture optimization, and ductility enhancement. For instance, during differential temperature rolling of AZ31B Mg alloy, {10–12} TTWs dominate the deformation process. This enhances material plasticity [63] by adjusting local crystal orientation to weaken the basal texture and coordinating deformation accommodation between grains. The multi–directional impact forging (MDIF) process induces numerous {10–12} TTWs in Mg–Gd–Y alloys through repeated multi–directional shear stresses. This simultaneously achieves grain refinement and promotes DRX, ultimately resulting in a UTS of 354 MPa and significant improvement in isotropy due to randomized grain orientation. However, it is important to note that while {10–11} CTWs and DTWs may assist DRX, the orientation mismatch between the twins and the matrix can induce localized stress concentrations. These create potential fracture initiation sites, which can adversely affect overall mechanical properties [72].

Recrystallization process optimization is equally critical, as it determines the microstructural basis for improved deformation coordination in Mg alloys by regulating grain size and grain boundary characteristics. Processes such as lateral SE and ES utilize high strain rates and localized strain concentrations to create favorable conditions for CDRX and DDRX, promoting the formation of fine, uniform recrystallized grains. For example, in AZ31 Mg alloy, the localized strain concentration induced by the SE process facilitates continuous accumulation of recrystallized grains and increases the proportion of high–angle grain boundaries (HAGBs). This effectively coordinates deformation between grains of different orientations while further weakening the basal texture, ultimately achieving a favorable balance between strength and ductility (Figure 17) [73]. Similarly, in microalloyed Mg–1.5Zn–0.5Zr–0.5Sr alloy, the ES process activates non–basal slip systems through shear deformation, reducing basal texture strength and enhancing deformation coordination. This increases fracture elongation from 13.9% (achieved via DE process) to 21.4%, demonstrating significant improvement in formability [62].

5.3. Mitigation of Tensile–Compressive Strength Asymmetry

Mg alloys have significant tensile–compressive yield asymmetry, where yield strengths differ under tensile and compressive loading. This reduces formability and limits engineering applications under complex stresses [9]. Research shows that shear deformation effectively mitigates this asymmetry by forming and reorienting {10–12} extension twin bands (ETBs). In AZ31 alloy, dense {10–12} ETBs generated during shear deformation transform into soft grains with new orientations, facilitating basal <a> slip activation. Subsequent recrystallization homogenizes the microstructure, reducing the tensile–compressive yield discrepancy [9]. Furthermore, shear band formation and propagation promote dislocation accumulation and release in localized strain regions, reorienting slip–resistant grain orientations to enhance overall deformation compatibility.

Texture evolution is another critical factor controlling tension–compressive asymmetry [9]. Conventional extrusion/rolling typically induces a strong basal texture, with basal planes predominantly parallel to the processing surface. This enhances tensile slip capacity but reduces compressive plastic deformation ability. Shear deformation disrupts this uniform orientation by introducing bimodal texture components or tilting basal planes, making them deviate along the shear direction. This establishes balanced slip system activation conditions under both tensile and compressive stresses. Such texture weakening effectively reduces mechanical response disparities caused by crystallographic orientation dependence.

DRX plays an important role in controlling tension–compression asymmetry during shear deformation. Fine grains formed via DRX not only enhance the material’s overall strength but also reduce orientation sensitivity by providing multiple slip paths. Notably, under low–temperature shear deformation conditions, the combined interaction between CDRX and DDRX mechanisms facilitates the formation of grains with diverse orientations, further improving the consistency of the material’s tensile and compressive yield behaviors [9].

The tension–compression yield asymmetry in AZ31 Mg alloy arises from the asymmetric activation of deformation mechanisms under tensile versus compressive loading. Shear deformation effectively mitigates this asymmetry by controlling the bimodal texture, {10–12} twin–derived structures, DRX, and dislocation activation modes [9]. Specifically, shear deformation induces the formation of a bimodal texture consisting of the typical shear texture T1 and the atypical shear texture T2 (Figure 8). The T1 texture features basal planes inclined at 30–45° to the ED and dominates basal <a> slip, while the T2 texture has basal planes nearly parallel to the ED and is associated with <c + a> slip. The {10–12} extension twins formed in the initial deformation stage develop into banded structures made of low–energy symmetric tilt grain boundaries (STGBs) that inherit the [10,11,12,13,14,15,17,18,19] rotation axis under continuous shear deformation (Figure 7). At a low temperature (250 °C), this banded structure increases the T2 texture intensity, raising the preferential activation of twinning during compression and resulting in a compressive CYS/TYS ratio of only 0.72. As the temperature increases to 340 °C, the DRX fraction rises, the average DRX grain size increases from 3.9 μm to 5.7 μm, the T2 texture intensity decreases, and the dispersion angle of deformed grains rises from 15° to 26°. Concurrently, grains with the T1 texture activate multiple slip systems (including prismatic <a> and pyramidal <a> slip), while grains with the T2 texture predominantly undergo basal <a>/<c + a> slip. The combined effects of low–energy grain boundaries and DRX–refined grains suppress the asymmetric contribution of twinning, increasing the CYS/TYS ratio to 0.87. Ultimately, the combined effects of texture regulation, grain refinement, and dislocation activation modes significantly mitigate the tension–compression yield asymmetry [9].

5.4. Summary and Synthesis

Shear deformation significantly enhances the mechanical properties of Mg alloys by facilitating grain refinement, texture modification, and dislocation hardening. A fundamental trade–off between strength and ductility is commonly observed in conventional Mg alloys during plastic deformation. While grain refinement strengthens alloys via the Hall–Petch relationship, excessive refinement hinders dislocation motion and induces strain localization, leading to a significant loss of ductility. Shear deformation effectively overcomes this limitation by constructing multiscale heterogeneous structures. Processes such as AAR and MDIF introduce directional shear strains that disrupt strong basal textures. For example, they tilt basal planes by 30–45° in AZ31 alloy, thereby reducing the activation threshold for non–basal slip and diminishing mechanical anisotropy [70]. Additionally, {10–12} extension twins with low CRSS segment coarse grains, forming high–energy interfaces that serve as nucleation sites for DRX. This process works synergistically with CDRX and DDRX to generate fine grains with a high fraction of HAGBs [68,73]. The resulting gradients in grain size, dislocation density, and recrystallized fraction–such as the notably reduced GND density in DRX grains of WE43 alloy [21]–promote coordinated deformation across different regions and mitigate local stress concentration. Meanwhile, LPSO phases enhance strength via Orowan strengthening and kink band formation while preserving ductility [51,69]. These integrated mechanisms enable simultaneous improvements in strength and ductility, as demonstrated by the following results: AAR–processed AZ31 sheet exhibits a uniform elongation of 17.9–18.5% [70], MDIF–processed Mg–Gd–Y alloy achieves a UTS of 354 MPa with improved isotropy [72], differential temperature rolling followed by annealing leads to an elongation of ~33.6% [71], and extrusion–shear processed Mg–Zn–Zr–Sr alloy attains a fracture elongation of 21.4% [62].

6. Defect Evolution and Regulation Under Severe Shear Deformation

6.1. Void Nucleation, Growth, and Closure

In metallic materials, internal voids (microscale pores or cavities) are key defects that cause a decline in mechanical properties. These defects significantly weaken the load–bearing capacity of materials and severely restrict their application in load–bearing structures in fields such as aerospace and automotive engineering. First, void nucleation is the initial stage of void evolution, specifically the process where nanoscale to microscale tiny voids first form in internal stress concentration regions (e.g., second–phase particle–matrix interfaces, high–angle grain boundaries, and dislocation accumulation zones) due to the mismatch between local stress distribution and material deformation capacity. The core of this stage is that local stress exceeds the material’s local cohesive strength, inducing the formation of tiny cavities. Subsequently, entering the growth stage, driven by the combined effect of external stress and atomic diffusion, the nucleated voids either undergo independent volume expansion or merge with adjacent small voids. This gradually forms larger–sized voids and further exacerbates imperfections in the material’s internal structure. However, under specific thermomechanical conditions, materials can fill, compress, or even completely eliminate existing voids through mechanisms such as high–temperature atomic diffusion (atoms migrate from the matrix to void interiors to fill cavities) and DRX (grain boundary migration of recrystallized grains squeezes small voids, promoting their shrinkage). This enables the repair of internal defects, thereby restoring or enhancing the material’s density and mechanical properties.

6.2. Shear–Specific Crack Initiation and Propagation

Mg alloys, with limited slip systems at RT, are prone to strain localization during conventional processing and easily form narrow shear bands followed by crack formation. This severely restricts their application in load–bearing components for high–demand fields such as aerospace and automotive engineering. As a core SPD technique, shear deformation exerts dual regulatory functions on crack initiation and propagation, and its effectiveness is related to the compatibility between shear parameters and material microstructural evolution.

The regulatory effect of shear deformation on cracks in light metals depends on this parameter–microstructure compatibility. Improper conditions (e.g., coarse initial grains, low deformation temperature) worsen strain localization and induce crack initiation. For AZ31 magnesium alloy, samples with an initial grain size of 30 μm form narrow, concentrated ASB during high–strain–rate room–temperature deformation, and these ASBs become stress concentration zones due to steep strain gradients, easily inducing microcracks. However, refining initial grains to 3 μm produces dispersed shear bands, significantly reducing crack initiation sensitivity [13]. Optimizing shear parameters (e.g., adjusting deformation passes, controlling temperature) further refines AZ31 grains to 3–10 μm, transforming shear bands from localized to widely dispersed, reducing strain hotspots and crack density, which stands in sharp contrast to the slender microcracks in coarse–grained samples [13].

For crack propagation, shear–induced texture optimization and microstructural adjustment are critical. Seong–Sik Lim et al. [74] noted that in Al–Mg alloys, shear deformation weakens strong rolling textures and promotes Goss texture (which hinders crack propagation, unlike Brass texture that allows crack penetration). It also activates multiple slip systems, relieving crack–tip stress and reducing rapid propagation. EBSD analysis shows that high KAM regions around cracks decrease with shear parameter optimization, confirming better deformation coordination (Figure 18). In Al–Mg–Si alloys, shear–induced precipitate shearing and planar defect formation affect deformation transfer, indirectly inhibiting continuous crack growth [75]. In summary, the core of shear deformation’s impact on cracks lies in optimizing strain transfer and stress distribution through grain refinement, dispersed shear bands, and texture regulation. Poorly controlled shear worsens cracking via strain localization and texture concentration, while optimized shear inhibits both crack initiation and propagation. This is achieved through grain refinement strengthening, activation of multiple slip systems, and formation of favorable textures, ensuring the structural reliability of light metals in load–bearing components.

6.3. Residual Stress Distribution and Regulation

Residual stress refers to the internal stress retained in metallic materials after thermomechanical processing, and it directly determines the dimensional stability, fatigue resistance, and long–term service reliability of components. The distribution of residual stress is typically jointly dominated by the strain gradients and temperature inhomogeneities generated during processing. For high–Mg–content Al–Mg alloys, studies by Seong–Sik Lim et al. [74] have confirmed that conventional rolling processes tend to induce significant stress concentration in the edge and central regions of the material, markedly raising the cracking risk. However, adjusting the rolling mode to cross–rolling and regulating the formation of Goss texture can effectively alleviate local stress accumulation (Figure 18). This phenomenon arises because Goss texture enhances the material’s deformation coordination, fundamentally reducing stress localization issues. For Al–Mg–Si alloys, analysis by Emil Christiansen et al. [75] indicates that shear deformation induces the shearing of β’’ precipitates and the formation of planar defects. This process can improve internal deformation transfer efficiency, thereby indirectly alleviating stress concentration. At the same time, dynamic recrystallization accompanying the shear process generates fine equiaxed grains, and the grain boundary migration process can further squeeze and dissipate local stress, providing auxiliary support for reducing the residual stress level.

6.4. Summary and Synthesis

Three key defects in metallic materials–internal voids, shear–induced cracks, and residual stress–seriously degrade mechanical performance and restrict their use in high–demand load–bearing structures. Of these defects, internal voids form in stress concentration regions caused by mismatches between stress and deformation. They grow under external stress and atomic diffusion, while high–temperature atomic diffusion and dynamic recrystallization can effectively close them. In terms of crack regulation, shear deformation–a core severe plastic deformation technique–acts as a key control tool. Improper parameter matching worsens strain localization and triggers cracking, whereas grain refinement, dispersed shear bands, and texture optimization effectively suppress crack initiation and growth. Residual stress, the third critical defect, arises from strain gradients and temperature inconsistencies during processing. Adjusting rolling methods or utilizing shear–induced microstructural changes–such as precipitate shearing and dynamic recrystallization–can alleviate this stress. In summary, optimizing thermomechanical processes to fully control these three defects is essential for improving the mechanical properties and service reliability of light metals.

7. Summary and Outlook

7.1. Summary

Mg alloys hold significant strategic value for lightweight design and sustainable development [62,76,77,78,79]. This review systematically summarizes recent advances in shear deformation techniques that induce microstructural evolution and enhance the mechanical properties of Mg alloys. Studies have shown that shear deformation provides a powerful means of tailoring Mg alloy microstructures by introducing high strain gradients and complex stress states, with the primary underlying mechanisms outlined as follows:

(1)

Shear deformation effectively promotes processes such as CDRX, DDRX, and TDRX. This results in marked refinement of coarse grains to the micron or even sub–micron scale, providing a key route to grain refinement strengthening.

(2)

Shear deformation effectively mitigates the undesirable strong basal texture in Mg alloys. Through activating {10–12} tensile twinning and non–basal slip, as well as fostering random orientations in recrystallized grains, it induces a bimodal texture or tilts the basal plane. In turn, this notably diminishes mechanical anisotropy and enhances the material’s formability.

(3)

Intense shear strain fragments, refines, and uniformly distributes second–phase particles. This not only enhances strength via dispersion strengthening but also leverages the Zener pinning effect to stabilize the fine–grained structure. Furthermore, these particles serve as nucleation sites for recrystallization, thereby further refining the microstructure.

(4)

Shear deformation processes enable the construction of surface–to–core gradient structures or recrystallized/UDRX heterostructures. These engineered microstructures facilitate the simultaneous improvement of strength and ductility via their synergistic actions.

(5)

Optimizing thermomechanical processes regulates metallic materials’ internal voids, shear–induced cracks, and residual stress, and these integrated defect control measures facilitate the synchronous enhancement of mechanical properties and service reliability of magnesium alloys via synergistic regulation.

Synergistic evolution of these microstructures markedly improves mechanical properties: grain refinement and high dislocation density–induced dislocation strengthening underpin enhanced strength, while texture weakening, increased slip system activation, and twinning–recrystallization coordination notably boost ductility, formability, and tension–compression yield symmetry. Advanced techniques (e.g., ECAP, CSE and its hybrid derivatives, ultrasonic vibration–assisted shear deformation) have opened new processing windows for precise, efficient control of Mg alloys’ microstructure and properties.

7.2. Outlook

Despite significant advances in shear deformation techniques for Mg alloy processing, sustained efforts in the following areas remain crucial to expand their industrial applicability and elevate performance to unprecedented levels (Figure 19):

(1)

Future research should prioritize investigating shear deformation processes coupled with multiple physical fields, including thermodynamics, mechanics, acoustics, and electromagnetism. Leveraging the synergistic effects of acoustic field–induced texture refinement, thermal field–driven microstructural homogenization, and electric field–enhanced regulation, texture asymmetry and stress concentration during shear processing can be effectively mitigated.

(2)

Furthermore, in situ characterization techniques enable real–time tracking of dynamic interactions between dislocations, twinning, and recrystallization during shear deformation. Coupled with multi–scale simulations–including crystal plasticity finite element (CPFE) and phase–field methods–this integrated approach quantitatively uncovers the intrinsic physical mechanisms underlying shear deformation–induced microstructural evolution and performance enhancement, spanning atomic, dislocation, and macroscopic scales.

(3)

For novel high–strength–high–ductility and multi–functional Mg alloys, customized shear deformation paths and parameters are critical to tapping their full potential. Notably, precisely controlling the interactions among LPSO phases, NM particles, and the shear strain field facilitates significant performance breakthroughs.

(4)

Current research primarily centers on static mechanical properties; future efforts should prioritize investigations into the service performance of shear–processed Mg alloys, including dynamic mechanical properties, fatigue resistance, corrosion resistance, and wear resistance. Efforts should also advance the scaling of processes from laboratory to pilot to industrial demonstration, establish a comprehensive “process–microstructure–properties–service behavior” database and corresponding control criteria, and deliver reliable material solutions for critical high–end equipment components.

(5)

The future development of Mg–based high–entropy alloys should center on developing shear deformation processes tailored for multi–principal element systems. By regulating shear parameters, processing inhomogeneity stemming from atomic size mismatch in these alloys can be mitigated, thus further improving the stability of the amorphous phase.

Notably, future shear deformation processes will achieve the coordinated optimization of environmental performance and energy efficiency: they enable indirect emission reduction through lightweighting and energy efficiency improvements via precise processing. With only a handful of potential environmental and energy–related challenges–all addressable via technical measures–these processes align with the long–term trend of low–carbon, high–efficiency manufacturing.