One of the metals’ most important microstructural features, crystallographic orientation, shows how corrosion behaves in a particular environment [41,42]. Different atomic coordination densities within different crystallographic planes cause surface energy levels to fluctuate. Surface energy decreases, anodic dissolution is reduced, and water or proton adsorption is enhanced in densely packed planes due to increased binding energies [43]. For instance, in magnesium, the surface energies of the (1 1 2 0), (1 0 1 0) and (0 0 0 1), planes are recorded as 2.156, 1.868, and 1.808 eV/nm², respectively [44], corresponding to 2.99 × 10⁴, 3.04 × 10⁴ and 1.54 × 10⁴ J/mol [45]. Comparatively, the electrochemical degradation rates of alternative crystallographic planes are much greater than those of the base plane (0 0 0 1) [46,47]. Surfaces that are exposed and mostly made of planes with the coordinates (0 0 01) are more resistant to corrosion than those that are mostly made of planes with the coordinates (1 0 1 0) or (1 1 2 0) [43,48]. The formation of oxide coatings in crystalline magnesium alloys is significantly influenced by grain size, as it plays a crucial role in determining the effectiveness of the coating [49]. A study by Ralston et al. [50] showed that the square root of the grain size is inversely associated to the metals corrosion rate. Magnesium oxide layers can also develop within the surface corrosion layers, enhancing the overall protection against further degradation. By increasing the fraction of volume at grain boundaries, grain refining decreases the lattice mismatch between the oxide layers of crystal structures and the Mg matrix. As a result, this promotes a more cohesive interface, increasing corrosion resistance [51,52]. Also, in passivation, surfaces of refined-grain Mg show faster kinetics, especially in chloride-containing electrolytes. Due to enhanced passivation capacity, the rapid production of magnesium oxide coatings allows these surfaces to withstand breakdowns efficiently [53]. Liu et al. [54] studied grain refining, which showed that it may slow the spread of stress corrosion fractures, especially near grain borders grain size modification is crucial for controlling the corrosion behaviour and durability of Mg alloys. Fig. 2 presents the structural configurations of Mg, Al, and Li, highlighting their atomic arrangements and bonding characteristics.

Substantial research was undertaken to understand the mechanisms that enhance the strength and stability of alloying elements in magnesium alloys. Many alloys experience microstructural instability and reduced strength at room temperature over time. However, Mg-Al-Li alloys, known for their relatively stable microstructures, have gained significant attention in the industrial sector. Recent studies by Tang [55] have demonstrated that incorporating aluminum (Al), cadmium (Cd), silver (Ag) and zinc (Zn), into Mg-Li alloys significantly enhances the strengthening effects within the β crystal structure. Tang early research indicated that these ternary alloys could achieve high hardness through water quenching from high temperatures. However, these alloys showed instability below 100°C, leading to microstructural softening. This softening was attributed to the loss of coherency between the θ-phase and the β-matrix due to increased lattice mismatch and the decomposition of the θ-phase into equilibrium phases such as Al-Li and Mg-Li. Reduced strength was linked to coarsening of the θ-phase at grain boundaries, although less so for Li than Al, Zn, and Cd. Recent research by Xin [56] on Mg-Li-Al-based alloys, using modern characterization techniques, has identified a semi-coherent cubic MgAl₃ nanophase as the transition phase. It is uncertain if alloys of Li, Ag, Zn, and Cd can adopt this particular structure. The strength at low temperatures is reduced by the decomposition of strengthening precipitates into equilibrium phases such as Mg-Al and Mg-Li, regardless of the transition phase. Based on early studies, several critical observations can be made regarding the stability and strengthening of the β Mg-Li-Al alloys transition phase. Elements with high supersaturated β-phase solubility, such as Al, Li, and Zn, offer the most substantial age-hardening response [57,58]. This occurs because highly soluble elements improve the coherency between the transition phases and the β-matrix, which in turn slows down dislocation motion. Moreover, higher supersaturated solubility typically results in a greater maximum fraction volume of strengthening precipitates. Conversely, small additions of insoluble elements do not improve the stability and hardening of Mg alloys and provide only limited benefits in Mg alloys [59]. The general elements solubility in the β-solid solution is summarized in Table 1. This knowledge gap suggests an area for potential future research to optimize the mechanical properties of Mg-Li-Al alloys.

Research examined the effects of ternary metals hardening with additions of Ce, La, Nd, and Y in Mg-Li alloys at concentrations of 3–5 wt%. These studies observed a modest increase in strength, approximately 10%–15% higher than that of binary Mg-Li alloys. This increase is notably less than the strengthening achieved with the additions of aluminum (Al), zinc (Zn), silver (Ag), and cadmium (Cd) [60,61]. Although alloys containing rare earth (RE) elements exhibited reduced softening compared to Mg-Li alloys with Zn and Al, the limited strengthening effect rendered RE elements unsuitable as primary alloying agents. Table 1 shows that the alloys failed to harden enough due to eutectic structures formation at the boundaries of grain and the low solubility of the RE components in the solid solution of β-Mg. The strengthening potential of other RE metals with atomic numbers greater than sixty in the β-solid solution remains unexplored [55]. Considering that these elements exhibit high supersaturated solubility in α-Mg [62], it is reasonable that they could also possess high solubility in the β-solid solution, as elements with high solubility in the β phase often exhibit similar behavior in the α phase [57]. Recent research has investigated the alloying of aluminum (Al) or lithium (Li) with magnesium and rare earth (RE) elements [35,39,63]. These studies indicate that such alloying can introduce significant microstructural complexity, including the formation of Mg-AlxREy intermetallic phases. These phases contribute to particle strengthening and may lead to the development of LPSO phases. Chakravorty reported that the Mg11-Li1.5-Al0.14 alloy exhibits stable mechanical properties [64]. Korgiopoulos et al. [65] enhanced the strength of the alloy by increasing the aluminum (Al) content and adding yttrium (Y) to the composition (Mg-10.95, Li-3.29, Al-0.19, Zr-0.59Y). The incorporation of Al₂Y particles refined the microstructure, leading to improved combinations of strength and ductility [66,67]. Additionally, introducing neodymium (Nd) into Mg-Li-Al alloys has enhanced mechanical properties and grain refinement [68]. Conversely, the addition of cerium (Ce) [69] and lanthanum (La) [70] resulted in minimal strengthening effects. LPSO phases have been successfully formed in Mg (α + β)-Al-Li alloys, positively influencing strength and ductility [71]. Table 1 provides the solubility of various elements in magnesium (Mg), highlighting the extent to which each element dissolves within the magnesium matrix.

Alloying represents a widespread method employed to modify the corrosion characteristics of Mg-Li-Al alloys. The inclusion of alloying elements influences the structure and composition of the surface oxide layer formed during corrosion. This is crucial because the hydroxide or oxide generated on the magnesium surface lacks the same passivation capabilities as aluminum oxides. The addition of elements such as Li, Al, rare earth metals, and others is expected to strengthen the stability and homogeneity of the surface layer composed of oxides like Al₂O₃, Li₂O, and Y₂O₃ [101]. Such expansion establishes a sustainable, stable, and restorable protective function. Nevertheless, the incorporation of alloying elements induces changes in the alloy microstructure, primarily affecting its corrosion kinetics. The extent of solid solubility of alloying elements in magnesium dictates whether they will precipitate as intermetallic compound phases, magnesium-based solid solutions, or impurity particles. These variables significantly influence the corrosion behavior of magnesium alloys [88].

Aluminum, zinc, and calcium serve as alloying elements in magnesium-lithium alloys. Incorporating aluminum into magnesium-lithium alloys enhances both the hydrogen and the composition and structure of the oxide coating. In biological environments, the rate of degradation of the LAZ631 alloy, characterized by a high aluminum (Al) content, is observed to be slower than the LAZ611 alloy. This disparity can be primarily attributed to the higher Al concentration in the LAZ631 alloy, resulting in a more stable surface layer that facilitates the passivation process. Furthermore, the addition of aluminum (Al) induces a transformation of dendritic crystals into equiaxial crystals, consequently reducing grain size noticeably [35]. With an increase in Al content, additional phases such as Mg₁₇Al₁₂ and Mg-Li phases are formed, creating micro-galvanic couplings within the matrix, thereby accelerating corrosion. The proportion of the Al-Li phase increases when introducing varying amounts of aluminum (Al) into a magnesium-lithium (Mg-14Li) alloy, forming a network at the grain boundaries. Consequently, the corrosion rate initially decreases, eventually increasing [39]. In NaCl solution of 0.7 mol/L, Kumar et al. [102] observed that the Mg-8Li-3Al-1Zn electrode demonstrates a higher corrosion potential and a lower corrosion current density than the Mg-8Li-3Al-0.5Zn electrode. This observation is attributed to the formation of uniform and fine products on the electrode’s surface. The introduction of calcium (Ca) into magnesium-lithium alloys, such as Mg-4Li [103], Mg-7Li-3Al [104], and Mg-14Li [105], results in a enhances corrosion resistance and reduction in grain size.

The effect of impurity elements on Mg-Li-Al corrosion resistance differs noticeably from observed traditional Mg alloys. Even a minute quantity of iron in Mg alloys can substantially reduce their resistance to corrosion. A particular concentration threshold exists beyond which the corrosion rate accelerates sharply once the iron content surpasses 0.016%. However, in Mg-Li-Al alloys, the rate of corrosion shows a linear increase with rising Fe concentration, devoid of any distinct threshold [106]. Moreover, the corrosion rate of Mg-Li-Al alloys containing minimal levels of Fe is not significantly higher than that of commercial Mg alloys with hcp-structured [107].

Incorporating rare earth (RE) elements into magnesium alloys yield several improvements, including refining the molten metal, enhancing grain structure, reducing potential differences between main and secondary phases, and altering the surface layers structure and composition. The standard electrode potential against SCE remains consistent at −2.372 V for both yttrium (Y) and magnesium. Coating the alloy’s surface with a thick layer of yttrium oxide (Y₂O₃) further enhances its corrosion resistance [108]. When yttrium (Y) is introduced into Mg-8Li-3Al-2Zn alloy, the quantity of AlLi phase decreases, transforming the α-Mg shape from elongated needles to more rounded forms [109]. The formation of an Al₂Y phase at grain boundaries prevents galvanic corrosion, unlike the α-Mg and β-Li phases. A reduction in AlLi content within the phase reduces corrosion within the Ξ-Li phase. For instance, in the LAZ832-1.5Y alloy, corrosion current densities of 83.757 × 10⁻⁶ A/cm² and 21.065 × 10⁻⁶ A/cm² illustrate a significant decrease in corrosion rate upon yttrium (Y) addition [109]. The formation of Al₂O₃ and Y₂O₃, facilitated by the migration of Y and Al atoms within the Al₂Y/LA143 composite, reduces long-term corrosion rates by impeding the ingress of Cl⁻ ions into the matrix [110]. Similarly, tin and yttrium (Y) in Mg-5Li-3Al-2Zn alloy reduce grain size, impede the formation of the AlLi phase, resulting in shorter corrosion pits and cracks, and lower corrosion current densities of 9.304 × 10⁻⁵ A/cm² [111]. Moreover, the presence of cerium (Ce), neodymium (Nd), and gadolinium (Gd) promotes the Al₂(Ce, Nd, Gd) formation with a secondary phase in Mg-Li-Al alloys. RE elements also influence the discharge and corrosion performance of Mg-Li alloys when utilized for batteries. Yttrium and lanthanum additions enhance the alloy’s corrosion resistance and electrochemical activity, improving battery performance [112,113].

Mechanical operations such as rolling, extrusion, and severe plastic deformation can significantly alter the microstructure of Mg-Li-Al alloy, impacting its behavior under corrosion conditions. These processes lead to changes in various microstructure properties of Mg alloys, including grain size, presence of secondary phases, dislocation density, and couple creation, which can drastically affect the corrosion rates exhibited by different Mg alloys. In their study, Cao et al. [114] studied the characteristics of corrosion of hot-rolled Mg-X alloys, where X represents elements such as Ca, Gd, Al, Sn, Mn, Sr, La, Nd, Ce, Si or Zr. They found that after hot rolling, the rate of corrosion decreased for most Mg-X alloys, attributed to the formation of more sophisticated secondary phases with a more uniform distribution. However, the corrosion rate increased for Mg-0.1Zr and Mg-0.3Si alloys due to forming an iron-rich phase on the rolled surface. Similarly, after equal channel extrusion, the corrosion rate of the AZ31 alloy significantly increased, mainly due to the increased dislocation and twin densities within the alloy microstructure [100].

Plastic deformation induces recrystallization, characterized by the development of small grains and redistribution of the second phase, which enhances corrosion resistance by creating a more microstructure (homogeneous). Dynamic recrystallization reduces the grain size of Mg-4Li-Al alloy from 100 to 5 μm during hot rolling [115]. The improved corrosion resistance observed in Mg-9Li-3Al-1.6Y alloy [116] and as-extruded Mg-11Li-3Al-2Zn-1.5Nd-0.2Zr alloy [117] is attributed to the decrease in grain size and consistent intermetallic phases distribution. The grain refinement during deformation in Mg-Li-Al alloys promotes the formation of crystal nuclei, which inhibit the formation of larger grains, thereby reducing dislocation density and energy storage in the material. Microstructural alterations resulting from mechanical deformation influence corrosion behavior. For instance, hot-rolled Mg-Li-Al alloy exhibits pronounced basal roughness parallel to the rolling surface, correlated with a decreased corrosion rate [118]. Ma et al. [119] observed that a 5% reduction in rolling results in more grains oriented perpendicular to the rolling surface, enhancing corrosion resistance. Conversely, a 10% reduction achieves similar results, while a 15% to 20% reduction leads to increased production of extension twins (10/12), promoting the formation of a surface oxide coating that enhances corrosion resistance.

Cross-cold rolling of Mg-Li-Al alloys significantly reduces corrosion rates by forming a uniform and thick layer of corrosion products on the surface [120]. However, plastic deformation, which introduces numerous dislocations and stresses, makes Mg-Li-Al alloys more susceptible to exfoliation corrosion. Exfoliation or pitting corrosion occurs when the oxide layer ruptures in a corrosive environment, leading to delamination along grain boundaries and the formation of tiny fissures. Equiaxed grains may delay fracture initiation, while increased rolling reduction intensifies exfoliation corrosion [121]. Annealing can help mitigate exfoliation corrosion by converting irregularly shaped granules into uniformly sized ones. Wei et al. [122] demonstrated that exfoliation corrosion in extruded Mg-1Li-1Ca alloy results from a combination of intergranular corrosion, filamentous corrosion, pitting corrosion, and stress corrosion, extending from Mg₂Ca particles along grain boundaries. Additionally, the direction of extrusion or rolling ratio affects the extent of exfoliation corrosion, with higher rolling ratios leading to decreased corrosion resistance due to increased microcouplings and area fraction of the β-Li phase [123].

Heat treatment procedures can significantly influence the rate of corrosion. Heat treatment often leads to changes in stress levels, increased atom mobility within the alloy, and the facilitated precipitation or dissolution of subsequent phases. In the case of magnesium alloys, heat treatment plays a crucial role as it alters the quantity and arrangement of secondary phases, thereby affecting corrosion rates [124]. Solution treatment can help establish a continuous corrosion barrier or dissolve secondary phase to delay corrosion. However, heat treatment can also accelerate the formation of iron-rich cathodic phases, potentially leading to faster corrosion rates [11]. In a study by Wang et al. [125], it was found that the β-(Mg, Zn)₃Gd eutectic phase in a Mg-15Gd-2Zn-0.39 Zr alloy resulted from a solid-solution, age-induced transition. This modification reduced micro-galvanic corrosion between the matrix and secondary phase, with the volume fraction and local potential of the LPSO phase significantly smaller than those in the eutectic phase. Various post-deformation processes, including solution treatment, aging, and annealing, can alter the corrosion behavior of Mg-Li-Al alloys. For instance, after treating the Mg-8Li-4Y-2Er-2Zn-0.6Zr alloy for 6 h at 450°C, the fishbone-like secondary phases at grain boundaries and the granular secondary phases within grains disappeared, while block-shaped secondary phases appeared [126].

Solvent treatment of Mg-Li-Al alloys can increase the matrix Al content and dissolve the AlLi phase. Similarly, at higher temperatures and longer aging durations, the quantity of AlLi phase increases. Rapid cooling in an ice-water combination causes the as-cast Mg-9Li alloy to rapidly transform from its bulk phase into tiny needle-like phases equally dispersed throughout the matrix. This transition occurs after heating the alloy to 575°C for 1.5 h. During this transformation, the phase transitions from body-centered cubic (bcc) to hexagonal close-packed (hcp) [99]. This specific microstructure leads to reduced corrosion and the formation of a surface coating on the β-Li phase, covering the entire alloy surface. Additionally, pitting corrosion is less likely to occur in the alloy. The α-Mg and β-Li phases exhibit almost identical elemental concentration and electrochemical activity due to the decreased atom dispersion in the solid solution, enhancing corrosion resistance [99]. A study [127] showed that the corrosion behavior of Mg-9Li-3Al-2.5Sr alloys under various aging circumstances. Alloys with the greatest corrosion resistance were those heated to 150°C for six hours, as the uniform distribution of the Al4Sr phase along the boundary grain helped to prevent further corrosion damage.

Annealing Mg-Li alloy at low temperatures for a brief time after extrusion or rolling is another method to enhance its corrosion resistance. Partially dynamic recrystallization occurs after one hour of annealing at 200°C in cold-rolled LZ91 sheets due to the significant deformation and thermal energy accumulated during cold rolling [128,129]. Corrosion current density and corrosion pits are reduced because the microstructure is homogeneous and the grains are polished. The resistance of corrosion of as-extruded Mg-7.5Li-3Al-1Zn alloy can be further enhanced by treatment of annealing, which speeds up the diffusion of Li atoms and alters the relative composition of the α-Mg and β-Li phases [39,130]. Increasing the volume of β-Li from 29% to 35% effectively decreases the corrosion current, as it reduces the area ratio of the cathode to the anode for corrosion (galvanic) [57]. According to study by Dobkowska et al. [131], changing the grain size of Mg-Li-Al alloy by the treatment of heat which has no effect on its resistance. Therefore, corrosion is mostly determined by the presence of the second phase. To make Mg-Li-Al alloys more corrosion-resistant, it is essential to anneal the alloy to reduce imperfections, age it to produce beneficial secondary phases, and use solid solution treatment to dissolve detrimental secondary phases.

Materials scientists frequently investigate the Mg-Al family of magnesium alloys. Within this system, two distinct mechanisms—strong solution strengthening and precipitation strengthening—contribute to the reinforcement of magnesium alloys [5,62,132]. Aluminum (Al) plays a crucial role in enhancing the overall quality of magnesium alloys by modifying their grain structure [12,34]. Additionally, a homogenous intermetallic phase called Mg17Al12 can disperse along grain boundaries or dissolve inside grains, resulting from the combination of magnesium and aluminum [133,134]. This intermetallic phase significantly enhances the corrosion resistance of magnesium alloys. Comparing AE42 and AE21 magnesium alloys reveals that AE42 alloys exhibit a more consistent element distribution and lower grain sizes, suggesting the rapid formation of stabilizing layers rich in aluminum. Consequently, AE42 alloys demonstrate superior corrosion resistance compared to AE21 alloys [135]. Increasing the aluminum content significantly affects the grain structure, as evidenced by experimental observations in Mg-Al-Ca-Mn alloys. At high aluminum concentrations, increased precipitation of secondary phases makes grain expansion more challenging during dynamic recrystallization. The resulting fine-grained structure substantially reduces corrosion with its increased number of grain boundaries [136]. Table 2 presents the corrosion currents and potentials of materials A04, A08, and A12 in a 3.5 wt% NaCl solution, offering valuable insights into their corrosion behavior.

Research in the field has extensively documented the influence of lithium (Li) concentration on the microstructure and properties of magnesium-lithium alloys (Mg-Li) [32,138]. Varying levels of lithium in Mg-Li alloys lead to the formation of various solid solutions, with the α-Mg phase being the most common at a weight percent of 10.3 lithium [139]. Characteristics of corrosion of Mg-Li alloys are closely linked to their crystal structures, where the concentration of lithium plays a more significant role in corrosion behavior than the type of crystal structure, whether hexagonal close-packed (hcp) or body-centered cubic (bcc) [93]. Table 3 presents corrosion rates for various Mg-Li alloys with different lithium concentrations, revealing that lithium in Mg-Li alloys reduces corrosion current density and inhibits hydrogen formation.

However, corrosion resistance is adversely affected by the β-Li phase, particularly at higher lithium concentrations [57]. Cain et al. [57] have reported that Mg-Li alloys with increased lithium concentrations exhibit superior corrosion resistance. Filiform corrosion is predominantly observed on α-Mg phase surfaces in Mg-4Li and Mg-7.5Li alloys with low lithium concentrations, while Mg-14Li alloy displays less pitting corrosion. Negative difference effect (NDE) experiments indicate that alloys with surface film rupture, such as Mg-4Li and Mg-7.5Li, exhibit increased NDE, whereas alloys with a β-Li phase show decreased NDE due to a persistent surface film [140].

According to Tian et al. [141], the as-extruded LAZ832-0.2Zr alloy demonstrates better corrosion resistance than the LAZ1132-0.2Zr alloy. The α-Mg alloy generates a significant amount of hydrogen gas when Mg-Li alloys with varying lithium contents are utilized as battery anodes, whereas the β-Li alloy notably reduces hydrogen gas emission. According to Wang et al. [142], the Mg-3Al-8Li alloy shows better self-discharge mitigation and more homogeneous dissolving of discharge products compared to Mg-3Al alloys with 12 or 4 wt%. In considering corrosion in Mg-Li alloys with high Li content, it is essential to account for various types of corrosion in the α-Mg and β-Li phases, as well as the impact of oxides, hydroxides, and carbonates on corrosion behavior [143,144]. The lithium concentration determines the crystal structure of Mg-Li alloys. Alloys containing less than 5.7% Li predominantly exhibit a hexagonal close-packed α-Mg phase. With increasing lithium content, alloys transition to α-Mg and β-Li phases. The β-Li phase, characterized by a bc cubic structure, becomes predominant after 11 weight percent [93,140,145]. The addition of lithium to magnesium enhances its strength and ductility compared to pure magnesium, further strengthened by the creation of protective Li₂CO₃/MgCO₃ layers, enhancing anti-corrosion properties [122].

When magnesium alloys are exposed to water and undergo electrochemical reactions, the magnesium matrix gradually dissolves, forming a surface hydroxide film and releasing hydrogen gas [147]. This process, known as general corrosion, is characterized by uniform deterioration across the material’s surface. The resultant film layer displays a hexagonal structure [11], wherein hydroxide and magnesium ions are arranged alternately within the crystal lattice. However, the bonding between these ions is relatively weak, making the film susceptible to cracking and providing limited protection to the underlying substrate. Despite its uniform nature, general corrosion is typically considered less damaging to alloys of magnesium than other corrosion forms [147]. The specific general product corrosion of a magnesium alloy is influenced by both its chemical composition and environmental factors that contribute to corrosion.

Localized corrosion, including pitting and filiform corrosion, presents a more significant risk to magnesium alloy substrates than general corrosion due to the heterogeneous corrosion rates observed across various regions [162,163]. Pitting corrosion arises when the protective surface film is ruptured, exposing the substrate to the corrosive environment and initiating localized corrosion processes [148]. Factors such as grain size, solution chemistry, alloy composition, and the presence of secondary phases significantly influence the occurrence of pitting corrosion. Smaller grain sizes typically correlate with reduced susceptibility to pitting corrosion [149]. Magnesium alloys are notably prone to corrosion in saline environments, particularly those containing chloride ions (Cl⁻), which can infiltrate surface coatings and trigger localized electrochemical corrosion, thereby promoting pitting corrosion [164]. Moreover, the pH of the solution plays a role in the start of pitting corrosion, with alkaline conditions (pH > 10.5) generally impeding its onset. Filiform corrosion, conversely, typically manifests beneath surface oxide layers and protective coatings [150]. This phenomenon is influenced by various factors, including the corrosive environment, microstructure, composition of oxide layers, and defects within the naturally occurring oxide layer on the alloy surface. Precipitates and alloying elements’ distribution within the material can serve as nucleation sites for corrosion, leading to its propagation in diverse directions. The microstructure of the alloy plays a pivotal role in filiform corrosion, where the potential difference between the cathodic tail and the anodic head acts as the driving force for its expansion [165]. A study suggested that when alloys of magnesium are exposed to a NaCl aqueous solution, filiform corrosion originates from second-phase particles accumulated along grain boundaries and propagates along these boundaries [166]. Notably, the microstructure significantly influences the pattern of filiform corrosion, with observations indicating its predominant occurrence and expansion within the α + β phase of layered structures. Grains within the α phase exhibit fewer filiform corrosion, while no filamentary corrosion is observed within the β-phase [167].

Galvanic corrosion, also recognized as bimetallic or contact corrosion, is an electrochemical phenomenon arising from the contact or immersion of two dissimilar metals or alloys with distinct electrode potentials [151,152]. In such scenarios, the metal with a lower electrode potential typically experiences accelerated corrosion, while the one with a higher potential corrodes slower. The degree of galvanic corrosion is primarily contingent upon the variance in electrode potentials between the involved metals [153]. Due to their relatively low corrosion potential, Mg alloys often serve as active anodes in galvanic couples, rendering them susceptible to preferential corrosion within such systems. Studies suggest a direct correlation between the rate of galvanic corrosion and the ratio of cathodic area to anodic area. With an increase in this ratio, the current density escalates, consequently augmenting corrosion rates [152]. The galvanic corrosion relationship between Mg alloys and other metals is influenced by effective contact area, electrode potential, and prevailing environmental conditions. The utilization of magnesium alloys in environments containing active chemicals poses challenges due to their vulnerability to galvanic corrosion. Within magnesium alloys, the presence of secondary phases and impurities introduces varied electrochemical activities, leading to the formation of micro-scale thermocouples. In corrosion, secondary phases or impurities frequently act as cathodes, while the Mg-rich phase acts as the anode due to its lower potential. This setup forms a galvanic couple, resulting in the generation of a corrosion current. The damage caused by galvanic corrosion correlates with the magnitude of the corrosion current flowing between these galvanic couples [168].

This relationship can be quantified using equation: (21)Igalvanic=(Epc−Epa)/(Rm+Rc)where:

I_galvanic represents the galvanic corrosion current.

Enhancing the corrosion resistance of magnesium alloys is a key objective in materials science, particularly through modifications in the corrosion (galvanic) dynamics between the magnesium matrix and the second phase. Recent investigations have highlighted the potential of Mg-RE-TM alloy systems to significantly improve mechanical properties, incorporating transition metals and rare earth elements in optimized proportions. The notable improvement observed in corrosion resistance is primarily attributed to a long-period stacking ordered (LPSO) structure within the second phases of magnesium alloys [169,170]. While the 18R and 14H LPSO structures are the most commonly observed, other phases such as 10H, 12H, 14H, 15R, 18R, 21R, and 24R also exist [170]. Despite significant advancements, a challenge known as micro-galvanic corrosion persists between the cathodic LPSO and anodic α-Mg phases. An in-depth investigation of the 18R-LPSO phase in magnesium alloys ZWM120, ZWM240, and ZWM480 revealed that an increased LPSO volume leads to reduced corrosion resistance under galvanic corrosion conditions [171].

Further research [172,173] suggested that VN3 Mg alloys, which have higher volume fractions of LPSO 18R phases, demonstrate enhanced resistance to corrosion compared to VN1.5 alloys. This improvement is primarily attributed to a continuous network of compacted LPSO phases in VN3 alloys, which inhibits corrosion growth at higher volume fractions. Conversely, alloys containing both 14H and 18R LPSO structures exhibit inferior corrosion resistance, indicating that magnesium-yttrium-zinc (Mg-Y-Zn) alloys featuring a single LPSO structure, such as 14H or 18R, may offer better corrosion resistance [174]. Despite the importance of these findings, further investigation into the corrosion behavior of magnesium alloys in the presence of the LPSO phase is warranted. These studies will contribute to a deeper understanding of the underlying mechanisms and aid in developing strategies to enhance the corrosion resistance of magnesium alloys in practical applications.

SCC arises in Mg alloys when they corrode due to environmental factors, leading to the development, elongation, and eventual spread of fractures under stress [154]. Two primary perspectives on SCC in magnesium alloys are recognized: hydrogen embrittlement and anodic dissolution mechanisms. As per the anodic dissolution theory, stress induces the slipping and rupture of the surface film on magnesium alloys. Consequently, the exposed matrix behaves as an anode compared to the surface film, causing rapid dissolution and continuous corrosion through electrochemical processes [155]. Conversely, the hydrogen embrittlement theory suggests that hydrogen produced during the magnesium alloy corrosion permeates the matrix and congregates at crack tips. When a critical concentration is attained, cracks initiate, elongate, and propagate [156]. Considering the inevitability of stress in practical applications of magnesium alloys, investigations into the AZ91 magnesium alloy corrosion process under varied stress levels have unveiled that while the initial corrosion product forms a protective layer, hindering direct interaction between the corrosive medium and the magnesium alloy matrix, the application of stress facilitates the corrosion pits formation. These pits allow corrosive substances to penetrate the substrate, quickening corrosion [175]. Stress corrosion cracking (SCC) typically occurs due to localized corrosion, and the presence of second-phase particles within the alloy increases its susceptibility to this type of corrosion. The rapid matrix dissolution adjacent to these particles and the detachment of the particles themselves contribute to extensive corrosion cracks. Moreover, grain size and orientation differences govern the crack propagation directions during SCC [137,176].

Determining elastic constants is paramount in evaluating a crystal’s resilience against external stressors and offers valuable insights into the bonding dynamics near equilibrium conditions [177]. Exploring these constants is indispensable for comprehending the mechanical behavior across different phases. Traditionally, alterations in a system’s volume exert a more pronounced effect on its total energy than strain-induced changes. Nevertheless, methods employing volume-conserving strains can mitigate the influence of volume variations on the total energy [178]. For cubic crystal structures such as Mg₁₇Al₁₂ and Mg₂Li, only three independent elastic constants are pertinent: C₁₁, C₁₂, and C₄₄. These constants serve as foundational descriptors for delineating the elastic properties of these phases. Such data stands as a cornerstone for advancing the understanding of the mechanical attributes of these materials and their prospective applications. A comparative analysis has been conducted between the Cauchy pressure (AZ), elastic constants (Cij), and anisotropic properties (AE) for the Mg-Al and Mg-Li phases, compared against values reported in extant literature. For the Mg-Al phase, the study reports C₁₁, C₁₂, and C₄₄ values of 90.5, 26.7, and 29.7 GPa, respectively, with a C₁₂–C₄₄ difference of −3.0 GPa, AZ of 0.93, and AE of 0.07. Another study calculations by [179] (88.38, 23.54, 26.95 GPa, respectively, with a C₁₂–C₄₄ difference of −3.41 GPa). Other studies have also reported similar values [180] (89.86, 27.21, 33.32 GPa, with a C₁₂–C₄₄ of −6.11 GPa), while Reference [181] found values of (97.7, 28.1, 31.4 GPa). Additionally, Reference [182] lists values of (70.58, 25.35, and 30.31 GPa, with a C₁₂–C₄₄ of −4.9 GPa and an AZ of 1.34). Experimental values from [183] indicate slightly different results, with (82.4, 20.8, and 36.6 GPa, showing a more significant C₁₂–C₄₄ difference of −15.8 GPa and an AZ of 1.19). For the Mg₂Li phase, Zhang et al. calculations reported values of (69.8, 25.9, and 31.1 GPa, with a C₁₂–C₄₄ difference of −5.2 GPa). Mg₁₇Al₁₂ and Mg₂Li phases are mechanically stable, satisfying cubic crystal’s corresponding mechanical stability conditions [184].

The precision of elastic constants can be significantly influenced by the reciprocal-space mesh utilized for integrating across the Brillouin zone, with discrepancies often linked to variations in the k-point sampling. In their investigation, De Boor et al. [185] explored the impact of temperature fluctuations on the elastic constants of Mg₂Li employing the quasi-harmonic approximation, observing that C₁₂ experiences the most pronounced alterations in response to temperature variations. Specifically, the elastic constant C₁₁ quantifies the material’s resistance to linear compression along the x-direction. The Cauchy pressure, denoted as C₁₂–C₄₄, elucidates the angular nature of atomic bonding in both compounds and metals [186]. A negative Cauchy pressure indicates directional bonding, a characteristic frequently observed in nonmetallic materials, while a positive value suggests metallic bonding. Notably, research by Wang et al. [187] indicated Cauchy pressures of −3.0 GPa and −4.1 GPa for Mg-Al and Mg-Li phases, respectively, indicative of their intrinsic brittleness akin to many non-metals. Anisotropic behavior in cubic crystals can be quantified by the parameter AE and Zener anisotropic index (AZ) [188]. (22)AZ=2C44/(C11−C12) (23)AE=(C11−C12−2C44)/(C11−C44)

The comprehension of mechanical properties in polycrystalline Mg-Li-Al phases necessitates a thorough understanding of shear modulus (G), isotropic bulk modulus (B), Poisson’s ratio (ν) and Young’s modulus (E). To derive these properties, the Voigt-Reuss-Hill (VRH) methodology is employed. This approach offers a framework to approximate the effective moduli of polycrystalline materials by aggregating upper (Voigt) and lower (Reuss) bounds of elastic constants.

In the context of a cubic crystal system, the limits of shear modulus (G) and bulk modulus (B), known respectively as Voigt and Reuss bounds, can be articulated as follows according to established literature [189]:

Voigt bounds (24)BV=C11+2C123 (25)BV=C11−C12+3C445

Reuss bounds (26)BR=(C11+2C12)(C11−C12)C11+2C12+2(C11−C12) (27)GR=5[4(C11−C12)+3C44]−1

The Hill average is the arithmetic mean of the Reuss and Voigt bounds: (28)BH=CBv+BR2 (29)BH=Bv+GR2

Using the Hill averages, Poisson’s ratio (ν) and Young’s modulus (E) are calculated as follows: (30)E=9BHGH3BH+GH

The elastic moduli serve as pivotal metrics for assessing material hardness. The bulk modulus (B) delineates a material’s capacity to resist volume alterations under pressure, while the shear modulus (G) signifies its ability to withstand alterations in shape [190]. Young’s modulus (E), the ratio of stress to strain, furnishes insights into a material’s rigidity, with higher values indicating greater stiffness [191]. Notably, research by Xue et al. [192] indicated that the Young’s modulus of Mg-Al surpasses that of Mg-Li, suggesting that Mg-Al exhibits greater stiffness.

The Young’s modulus aids as a visual appearance of a material’s anisotropic attributes. For cubic systems, understanding these directional dependencies necessitates familiarity with the theoretical expressions governing Young’s modulus (E). To elucidate the anisotropic characteristics, 3D surface contours of Young’s modulus for Mg-Al and Mg-Li are presented. These expressions for cubic systems are articulated as follows: (31)E(θ,ϕ)=[S11−2(S11−S12−12S44)(cos2⁡θcos2⁡ϕ+cos2⁡θsin2⁡ϕ+sin2⁡θ)]−1where:

S_ij are the compliance coefficients.

Fig. 4 displays the relationship between strain (εx) and various mechanical properties: (a) Bulk modulus (B), (b) Shear modulus (G), (c) Young’s modulus (E), and (d) Poisson’s ratio (µ), illustrating their respective behaviors under strain.

Orhan et al. [193] introduced the pugh ratio concept that the ratio of bulk to shear modulus (B/G) serves as a common metric for assessing the brittleness or ductility of materials. A higher B/G ratio signifies a ductile response, while a lower ratio indicates a brittle behavior. The threshold value distinguishes between ductile and brittle materials at 1.75. Moreover, previous research [194,195] highlighted the significance of Poisson ratio in characterizing a crystal’s resistance to shear deformation, typically ranging from −1 to 0.5. Elevated Poisson’s ratios indicate superior plasticity, with values surpassing 1/3 indicative of ductile characteristics and those falling below 1/3 suggesting brittleness. In another study, Poisson’s ratios below 1/3 further corroborate the brittle attributes of Mg-Li-Al alloys [196]. First-principle calculations of Mg-Al-Li alloys indicate that increased mechanical strength enhances corrosion resistance. Higher strength reduces microcrack formation and propagation, while refining the alloy’s microstructure. This results in fewer grain boundaries and more uniform protective oxide layers, which together help to resist corrosive agents and prevent stress corrosion cracking. Fig. 5 shows the Poisson’s ratio and Pugh ratio of the Mg-Li-Al alloy.

The determination of phase velocities for pure longitudinal and transverse modes within the Mg-Li-Al phases can be achieved by using single-crystal elastic constants in conjunction with Christoffel equation resolution [193]. Prior investigations explained that the determination of sound velocities is depending upon the symmetry of the crystal and the direction of propagation. Specifically, in cubic crystals, the calculation of sound velocities for both transverse and longitudinal waves across three unique crystal orientations—[1 1 0], [1 0 0], and [1 1 1]—was accomplished utilizing the equation provided in Reference [197]: (32)VT1[100]=VT2[100]=C44/ρ,VL[100]=C11/ρ,VT1[10]=C44/ρ,VT2[110]=(C11−C12)/ρ,VL[110]=(C11+C12+2C44)/2ρ,VT1[111]=VT2[111]=(C11−C12+C44)/3ρ,VL[111]=(C11+2C12+4C44)/3ρ.

Phases with lower density and higher elastic constants exhibit greater sound velocities [198,199]. A study by Fu et al. [200] highlighted that longitudinal waves attain their maximum speed along the [1 0 0] direction for Mg-Al and along the [1 1 1] direction for Mg-Li. These variations in sound velocities highlight the elastic anisotropy inherent in both phases. The sound velocity prediction across different crystallographic directions for both materials represents an initial endeavor [201].

The Debye temperature (ΘD) is important in understanding various solid-state physics phenomena, including lattice vibrations, thermal conductivity, and specific heat. Additionally, unlike molecular weight, it provides valuable information about the bonding strength present within a crystal lattice. In determining the Debye temperature for the Mg-Li-Al system, researchers utilize the average sound velocity (Vm), as outlined in the equation [202]: (33)ΘD=hkB[3n4π(NAρM)]13Vm,

The factors here consist of Planck’s constant (h), Boltzmann’s constant [175], Avogadro’s number (N_A), molecular weight (M), mass density (ρ), and the no. of atoms per formula unit (n). By utilizing this equation, one can calculate the avg wave velocity (V_m) of a polycrystalline material [203]: (34)Vm=[13(2Vt3+1Vt3)]−13,

Using Hill’s shear (G) and bulk modulus (B) in Navier’s equation allows us to determine the transverse (V_t), longitudinal (V_l) sound velocities of a material [204]: (35)Vt=Gρ,Vt=(B+43G)/ρ.

Recent research findings have demonstrated a positive correlation between Debye temperature and thermal conductivity, indicating that materials with higher Debye temperatures exhibit greater thermal conductivity. The measured ΘD value for Mg-Li-Al surpasses that of other metallic counterparts, indicating its superior thermal conductivity and enhanced chemical bonding characteristics [55]. Higher Debye temperature and sound velocity in Mg-Al-Li alloys indicate greater atomic bonding strength and elastic stiffness, respectively. These properties enhance corrosion resistance by reducing atomic vibrations and preventing microstructural defects, thereby improving the integrity of protective oxide layers and mitigating pathways for corrosive agents. Fig. 6 depicts the temperature-dependent energy (G), illustrating how energy varies with changes in temperature.

In examining the structural stability and bonding attributes of Mg-Li-Al, an analysis of the partial density of states (PDOS) and the total density of states (TDOS) in proximity to the Fermi level was directed [205,206]. The Fermi level (EF) serves as the reference point, with zero energy corresponding to it. Recent studies [207,208] have revealed the presence of multiple energy states intersecting the Fermi level within the DOS structures, indicative of metallic behavior. The primary peaks in bonding for Mg-Al within the lower energy spectrum are primarily attributed to the Al 3s states, with a minor contribution from Mg 3s and Al 3p states. Al 3p and Mg 2p states predominate close to the Fermi level, suggesting that the metallic behavior is closely linked to the p states of Al or Mg atoms. The observed hybridization between Al and Mg states suggests the presence of bonding similar to covalent interactions in the Mg-Al phase.

Further studies [209,210] stated TDOS for Mg-Li reveals two discernible regions below the Fermi level (0 eV). The first region, spanning from −10 to −5.0 eV, is largely influenced by Li 5s states with some contribution from Mg 3s states, exhibiting strong hybridization between Mg 3s and Li 5s states. In the region spanning from −5.0 eV to the Fermi level, the dominant states are primarily associated with Li 5p states. Additionally, interactions between Mg s-p and Li s-p states are observed above the Fermi energy level. The Mg-Li system exhibits the characteristic small band gap typical of semiconductor materials with a predicted total density of states (TDOS) value of only 0.363 electrons per electron volt per formula unit (eV-1 f.u.-1). Wang et al. [211] have reported a lower number of bonding electrons per atom for Mg-Al compared to Mg-Li, suggesting greater stability for Mg-Li due to a higher number of bonding electrons. Mulliken population analysis by Li et al. [212] indicates electron transfer from Mg atoms to Al and Li atoms in Mg-Li-Al, indicating an ionic nature. This relative electron transfer serves as a qualitative parameter for system characterization. Additionally, bond population analysis, as utilized by another study [213], provides an objective and quantitative measure for assessing bond character. A value of −2 indicates an anti-bonding state, while a value of 1 signifies increasing degrees of covalency. A value of 0 suggests no bonding. The electronic structure analysis reveals that bonding in Mg-Li-Al encompasses a combination of covalent, ionic, and metallic characteristics, with Mg-Li-Al bonds exhibiting a higher degree of covalence than other metallic bonds [214]. The electronic structure of Mg-Al-Li alloys, characterized by metallic behavior and covalent bonding, influences corrosion resistance. Stronger covalent bonds and significant electron transfer enhance the stability of the alloy’s surface, reducing susceptibility to corrosive agents and supporting the formation of durable protective oxide layers.

DFT-based first-principles calculations conventionally operate under absolute zero temperature (T = 0 K), a restriction that also extends to phonon analyses. These computations play a pivotal role in deducing the thermodynamic attributes of materials, encompassing temperature-dependent variables such as entropy (S(T)), free energy (G(T)), enthalpy (E(T)), and constant volume heat capacity (CV(T)). To evaluate these properties for the Mg-Li-Al compound, the PHDOS covering both partial and cumulative aspects is assessed at the theoretical equilibrium volume [205]. Research conducted by Li et al. [215] highlights that the peaks observed in the PHDOS of Mg-Li-Al predominantly arise from the interactions among Mg, Li, and Al atoms. Notably, Mg atoms exhibit slightly greater contributions than Li and Al atoms. Additionally, findings from another study [216] emphasize the prominence of peaks in the low-frequency domain (around 4 THz), primarily attributable to Mg atoms movements. These varied contributions from constituent atoms are influenced by their discrepancies in mass. The outcomes of PHDOS analysis facilitate determining the system’s vibrational free energy. Subsequently, employing the quasi-harmonic approximation (QHA) [217], one can evaluate the thermodynamic parameters vibrational contributions, including enthalpy, entropy, free energy (Helmholtz), and constant volume heat capacity.

The enthalpy in QHA are provided as follows [217]: (36)E(T)=Etot+Ezp+∫ℏωexp⁡(ℏωkBT)F(ω)dω

The equation for E_tot defines the total of electronic energy, while E_zp signifies the zero-point vibrational energy. Moreover, the symbol ħ denotes constant of Plancks (h/2π), and F(ω) represents the PHDOS.

The computation of F(ω) and E_zp can be performed utilizing the expressions provided in [218]: (37)F(ω)=1N∫BZ∑jδ[ω−ωj(q¯)]dq¯Ezp =12∫F(ω)ℏωdω

Then, the free energy is given by [219]: (38)G(T)=Etot +Ezp+kBT∫F(ω)ln⁡[1−exp⁡(−ℏω/kBT)]dω

The entropy can be obtained as follows [220]: (39)S(T)=kB{∫ℏω/kBTexp⁡(ℏω/kBT)−1F(ω)dω−∫F(ω)ln⁡[1−exp⁡(−ℏω/kBT)]dω}

The heat capacity is given by [17]: (40)CV(T)=kB∫(ℏω/kBT)2exp⁡(ℏω/kBT)[exp⁡(ℏω/kBT)−1]2F(ω)dω

According to a study conducted by Guo et al. [205], the computed values of entropy, free energy, enthalpy, and heat capacity for Mg-Li-Al, utilizing the Quasi-Harmonic Approximation (QHA), vary with temperature. These calculations are performed under zero pressure conditions, where the Helmholtz free energy equals the Gibbs energy. The entropy (S) and enthalpy (E) for Mg-Li-Al increase gradually as temperature rises, while the free energy (G) decreases with increasing temperature. Another study by Xin [56] emphasized that the entropy of Mg-Li-Al begins at zero at 0 K and increases rapidly as the temperature exceeds 300 K. The current entropy values are consistent with those presented in a study by Zhang et al. [181], which showed minimal differences between Gibbs energy (G) and calculated entropy with and without contributions of thermal electronic. Another study indicated that the computed Gibbs energy is −6.06 kJ mol⁻¹ at 300 K, which aligns with values reported in [17]. Fig. 7 presents a graph showing the variation of enthalpy, Gibbs free energy, and entropy as a function of temperature in a crystalline solid, providing insight into the thermodynamic behavior of the material.

The capacity of heat is a significant thermodynamic parameter that provides valuable information regarding the materials vibrational characteristics. This study investigates the temperature-dependent variations in heat capacity at constant pressure (CP) and constant volume (CV) for Mg-Li-Al over a temperature range of 0 to 1000 K. The findings reveal that CV experiences a notable increase with temperature up to 300 K, following Debye’s law. Beyond this temperature, CV stabilizes, exhibiting minimal dependence on temperature and converging towards the 3 NR, where R signifies the ideal gas constant and N represents the total number of atoms in the cell and Dulong–Petit limit. The results indicate that CP values, as computed by Zhang et al. [181], demonstrate slight differences with and without contributions of thermal electronics. At lower temperatures, CP closely parallels CV; however, at higher temperatures, CP exceeds CV due to the impact of thermal expansion. Notably, experimental data on the heat capacity of Mg-Li-Al are lacking in many studies. Thus, these predictions serve as valuable guidance for future theoretical and experimental investigations across various temperature ranges. Thermodynamic properties, such as heat capacity, and enthalpy, influence corrosion resistance in Mg-Al-Li alloys. Higher Debye temperatures and stable thermodynamic profiles enhance structural stability and reduce atomic mobility, leading to more robust protective oxide layers and improved resistance to corrosive environments.

Several studies have investigated the relationship between grain size and corrosion properties in alloys, particularly magnesium and magnesium alloys [222]. Grain refinement is crucial for enhancing corrosion resistance as it increases the density of grain boundaries, making it easier to form a protective coating and ensuring better adhesion of the oxide layer to the substrate. However, there is a debate about whether smaller grains accelerate corrosion due to increased active defects at grain boundaries. Moreover, grain refinement can form isolated cathodes along grain boundaries, facilitating micro-galvanic corrosion. Various casting techniques and processing steps are employed to modify the grain size of Mg-Li alloys [223,224]. Wu et al. [225] demonstrated the impact of grain size reduction by using a rapid solidification technique to produce Mg-13Li-X alloys with fine-grained structures. This resulted in a substantial decrease in grain size, from 150 μm in the traditional as-cast alloy to 4.2 μm. The alloy with smaller grains exhibited a Ecorr of −1.354 V and a corrosion current density of 5.830 × 10⁻⁷ A/cm² after immersion in Hank’s Balanced Salt Solution (HBSS) for two hours. Grain refinement also enhances the corrosion resistance of Mg-Li-Al alloys. A delicate spherical shape can be achieved by incorporating ultrasonic vibration into the melt solidification process, transforming the coarse flower-like structure of the α-Mg phase and reducing corrosion current density and corrosion product production [226].

Comparing the Mg-9Li-3Al-1Ca alloys microstructure, cast by centrifugal casting and gravity casting revealed that the latter method significantly reduced the corrosion current density by refining and spheroidizing the α-Mg phase [227]. Research [228] has demonstrated that improving the grain structure effectively prevents corrosion cracks from penetrating the alloy, thereby reducing the corrosion rate and transforming the corrosion pattern from localized to uniform. Friction stir treatment significantly enhanced the surface finish of Mg-9Li-1Zn alloy grains, reducing corrosion current density and forming numerous small, shallow pits instead of large, deep holes. Generally, reducing grain size improves the corrosion resistance of Mg-Li-Al alloys. Adding half a nanometer of nitrogen to the LAZ832 alloy transforms the equiaxed α-Mg phase into a lengthy ribbon structure, enhancing corrosion resistance despite increasing grain size [229]. Aging temperature also affects corrosion resistance, with higher temperatures leading to larger grain size and improved corrosion resistance in Mg-9Li-3Al-2.5Sr alloys [230]. Moreover, the behavior of corrosion changes with increasing grain size due to the rapid growth of passivation layers and segregation of impurities at grain boundaries. Gao et al. [231] observed uneven distribution of grain sizes within the alloy can affect its corrosion performance. Additionally, the finer grain structure in the top layer of as-extruded Mg-1Li-1Ca alloy exhibited greater corrosion resistance [232]. The following equation can be utilized to elucidate the correlation between grain size and corrosion rate. Similar to the Hall-Petch equation, the rate of corrosion (CR) [233] can be expressed as: (41)CR=a+bD−0.5

The parameter denoted as D denotes the grain dimensions, whereas the parameters a and b pertain to additional influencing factors. Typically, diminishing the grain dimensions enhances the corrosion resistance of Mg-Li-Al alloys. Nevertheless, there exists a scarcity of comprehensive data concerning intergranular corrosion within Mg-Li-Al alloys, with sporadic investigations into the association between grain size and corrosion resilience in ternary Mg-Li-Al alloys.

The role of secondary phases significantly influences the corrosion behavior of magnesium alloys, often triggering galvanic corrosion between the matrix and these secondary phases [234,235]. Typically, secondary phases act as cathodes compared to the matrix, accelerating the corrosion rate. Phases like Al-Fe and Al-Mn can drastically increase the corrosion rate of the Mg matrix when serving as cathodes. An intriguing secondary phase known as Mg₁₇Al₁₂, found at grain boundaries of AZ alloys, serves a dual role: acting as a cathode to promote hydrogen evolution and as a barrier to protect the material from corrosion. The impact of the Mg₁₇Al₁₂ phase depends on its volume fraction and distribution within the alloy microstructure. Elements like La, Nd, and Y facilitate the formation of secondary phase anodes, which are more prone to degradation [236,237].

As previously mentioned, Mg-Li-Al alloys feature typical secondary phases in regular Mg alloys and phases containing Li. Galvanic corrosion processes are common in Mg-Li alloys due to the electrochemical mismatch between the α-Mg and β-Li phases. The local potential distribution around the α-Mg and β-Li interface and the internal β-Li phase at grain boundaries indicates micro-galvanic corrosion to some extent [57]. Additionally, the structure and composition of the α-Mg phase significantly influence the corrosion dynamics of Mg-Li-Al alloys. The AlLi phase, present in Mg-Li and Al alloys, due to lithium’s higher electronegativity compared to magnesium and aluminum, increases the susceptibility of the Mg matrix to galvanic corrosion dissolution [39,238]. Filiform corrosion often extends along the AlLi phase in Mg-Li alloys, with the ϋ-Li phase and grain boundaries being common sites for this phenomenon. Moreover, the corrosion properties of Mg-Li-Al alloys are influenced by the quasicrystal phase and the long-period stacking ordered (LPSO) phase. Reference [239] discovered that the I phase at grain boundaries in Mg-6Li-6Zn-1.2Y alloys effectively prevents the formation and development of corrosion pits and filiform corrosion, thereby slowing down the corrosion process. Li et al. [240] studied the microstructure and corrosion behavior of Mg-6.5Li-xY-yZn alloys during extrusion. They found that certain compositions effectively decrease micro-galvanic corrosion and have the least second phase, such as the Mg-6.5Li-0.8Y-0.3Zn alloy. This alloy acts as a barrier between the eutectic phase and the matrix.

The ability of secondary phases to resist corrosion is governed by several key factors, including their quantity, structure, potential gradient relative to the substrate, and surface protection function [241]. The potential differences between particles and matrix affect the corrosion rate in LA83-1.2M alloys containing Al₂Gd/Mg₂Sn/AlCuMg particles. This study comparing the relative potential differences of elements in these alloys revealed significant variations, with AlCuMg showing about 680 mV, AlLi approximately 340 mV, Al₂Gd around 250 mV, β-Li about 135 mV, and Mg₂Sn roughly 60 mV. The addition of more noble AlCuMg particles can reduce mass loss by approximately 70% by forming micro-galvanic couples with Al₂Gd and α-Mg, which have smaller potential differences [242]. Altering the solution treatment time in Mg-3Nd-1Li-0.2Zn alloys affects the size and quantity of the Mg41Nd5 secondary phase, leading to the formation of smaller, shallower pitting pits after six hours of treatment at 535°C [243]. Additionally, in cross-cold rolled Mg-Li-Al alloys, a more uniform and compact corrosion morphology is observed due to the homogeneous distribution of smaller secondary phases throughout the matrix, resulting in the accelerated formation of surface corrosion products [120,244]. Using thin and uniformly distributed second phases can also enhance resistance by creating a protective layer on the surface of alloys, such as the as-extruded Mg-1.21Li-1.12Al-1Y alloy [245]. Purifying Mg-Li-Al alloys to reduce hazardous contaminants and minimize the formation of a strong cathode phase is essential for corrosion resistance. These alloys’ corrosion resistance can be strengthened by reducing the volume fraction of the AlLi phase by adding additional alloying elements. Additionally, generating a network phase similar to the Mg₁₇Al₁₂ phase at grain boundaries can act as a barrier to prevent corrosion, significantly enhancing the resistance of Mg-Li-Al alloys by including small, uniformly distributed second-phase constituents.

The corrosion coatings structure is crucial in determining the corrosion rate of magnesium and its alloys. Effective coatings should be thick and stable to prevent the metal substrate from further exposure to air or ions. When magnesium is exposed to air, it spontaneously forms a thin, compact oxide coating consisting of magnesium oxide (MgO) and magnesium hydroxide (Mg(OH)₂) with an amorphous structure [246]. This coating typically comprises a amorphous multicellular hydration layer as the innermost layer, with an outermost layer usually ranging from 20 to 40 nanometers thick in damp environments [247]. However, when immersed in water, the surface coating formed on magnesium alloys is often loose and porous, providing inadequate protection for the alloy [246]. The surface film structure and composition of Mg-Li-Al alloys exhibit characteristics influenced by active Li and Al elements and different crystal structures. Li et al. [93] studied the properties and effectiveness of surface coatings on pure Mg with a hexagonal close-packed (hcp) structure and Mg-14Li alloy with a body-centered cubic (bcc) structure. They observed Mg(OH)₂ particles developing on the surfaces of pure Mg, while the surface of Mg-14Li showed needle-like particles with tiny holes. Research suggests that excessive anode overvoltage can lead to the disintegration of both surface layers. In Mg-Li-Al alloys, a thicker oxide coating develops with increasing lithium content, and the presence of dual-phase Mg-Li alloys can further increase corrosion rates, forming a thicker oxide coating [248,249]. The presence of Li in Mg-Li alloys introduces Li-containing compounds into the surface film composition, resulting in a unique multilayer structure. The surface of the Mg-8Li alloy exhibits a four-layer structure when exposed to normal air conditions, with layers consisting of compounds like sodium hydroxide, lithium oxide, magnesium oxide, and lithium hydroxide [250].

Iron and lithium components are readily oxidized in Mg-Li-Al base alloys, forming MgO, Al₂O₃, and Li₂O. A partial hydroxide transformation occurs when these alloys come into contact with water molecules. Thermodynamic studies suggest the substantial presence of Mg(OH)₂, Li₂O and Al₂O₃ in the surface coating [251]. Sun et al. [140] observed a four-layer structure in the oxide coating on Mg-Li alloy surfaces after prolonged exposure to air, with the outermost layer primarily composed of lithium compounds like LiOH, Li₂O, and Li₂CO₃. Reference [55] examined the surface and bulk microstructural properties of a new Mg-Li-Al alloy called LATZ9531. He found a six-layer structure of multi-oxide in the surface film, with Li, Mg, and Al being the predominant components. The concentration of Li gradually decreases as one moves inward into the matrix. Alloying elements can significantly influence oxide coatings on metal surfaces, resulting in thinner, denser, and more protective layers [252]. In the LAZ832 alloy, the addition of Yttrium (Y) leads to the growth of numerous tiny, vertically oriented flakes on the surface, enhancing the protective characteristics of the oxide coating and preventing further corrosion of the underlying matrix [109]. The presence of YAl₂ in Mg-Li alloys promotes the Y₂O₃ formation on the surface, improving the adhesive strength between the oxide coating and the substrate material [253]. The protective effectiveness of oxide coatings on metal substrates is often evaluated using the Pilling-Bedworth ratio (PBR) [254]. This ratio, which compares the oxide volume to metal, provides insights into the oxide layer’s stress state, integrity, and compactness. A PBR greater than 1 indicates compressive stress in the oxide layer, while values below 1 suggest tensile stress. Larger deviations from unity indicate higher stress levels and a greater likelihood of surface oxide layer rupture [255].

Cain et al. [57] and Xiang et al. [256] attributed the exceptional resistance of corrosion of the Mg-10.95Li-3.29Al-0.19Zr-0.59Y alloy to the production of a stable and protective Li₂CO₃ oxide coating on its surface. Lu et al. [257] conducted a comprehensive study on the chemical composition of surface coatings formed on Mg-Li(-Al-Y-Zr) alloys in air-and water-based solutions. Their analysis revealed that the Mg-Li(-Al-Y-Zr) alloy exhibits a complex, multi-layered surface coating, mainly Li₂CO₃ and Mg(OH)₂. The outer layer has a porous structure with a consistent distribution of lithium, while the inner layer consists of a thicker layer formed by combining Li-containing compounds with other substances, enhancing corrosion resistance. Surface films crucial for corrosion resistance often include Li₂CO₃ and Al₂(CO₃)₃, which are formed when exposed to carbon dioxide in the air. The dissolution of magnesium, aluminum, and lithium in water generates oxides and hydroxides of these elements. Further studies are needed to fully understand how lithium, aluminum, and rare earth elements affect the stability of these compounds and the corrosion resistance of Mg-Li-Al alloys. Table 5 outlines the factors that influence the corrosion behavior of Mg-Li-Al alloys, detailing the key variables that affect their corrosion resistance.