Microstructural, mechanical, and corrosion characteristics of Mg–Gd–x systems: A review of recent advancements

1 Introduction

Engineers and scientists have developed a keen interest in the optimization of fuel consumption through the reduction of vehicle structure weight in recent years. Steel and aluminum are utilized as structural components in the majority of the aerospace and automotive industries. However, in order to optimize fuel efficiency and reduce the weight of vehicles, lighter materials that do not compromise performance are necessary. Compared to steel (7.86 g/cm³) and aluminum (2.7 g/cm³), magnesium (Mg) is a relatively lightweight material with a density of 1.738 g/cm³. Magnesium is therefore regarded as a viable substitute for traditional metals utilized in aerospace and automotive structural applications. A few properties and applications of the Mg–rare earth (RE) alloy are shown in Figure 1. Automobile interiors utilize magnesium (Mg), including stampings in instrument panels, steering wheels, and steering column components [1,2,3,4,5]. The density and elastic modulus of magnesium fall within the range of values observed in natural human bone, which are between 1.8 and 2.1 g/cm³ for density and 40–57 GPa for modulus [6]. Furthermore, Mg²⁺ ions have the potential to facilitate tissue repair, and Mg can be eliminated from the body through perspiration and excretion [7]. Mg and its alloys are suitable for biodegradable implants due to their excellent mechanical properties, biocompatibility, and biodegradability [8]. The metal hydrides effectively store hydrogen due to their exothermic oxidation with the atmosphere and the generation of hydrogen gas from the reversible reaction of magnesium dihydride (MgH₂) [9]. One possible method for hydrogen storage is a magnesium hydride or a system based on magnesium [10]. Magnesium can serve as the anode in air batteries and oxygen as the cathode. The effectiveness of the magnesium anode is constrained by the self-corrosive reactions and the corrosion product (Mg(OH)₂) [11].

Figure 1

(a) Potential properties of the Mg–Gd alloy. (b) Various potential applications of Mg–Gd alloy [5].

The ductility of the Mg alloy is constrained by a crystal system with hexagonal packing and a reduced number of active slip systems [12]. The aluminum–zinc (AZ), aluminum–manganese (AM), zinc–zirconium (ZK), zinc–rare earth (ZE), yttrium–rare earth (WE), and electron series are just a few examples of commercial alloys that have been produced with various alloying additives, including RE minerals, to enhance their properties. As a result of mechanisms such as dispersion and solid solution strengthening, the Mg–RE alloy displays excellent mechanical properties [13,14]. The mineralogy of RE materials can be broadly classified into two principal categories: light (La–Eu) and heavy (Gd–Lu). The mechanical properties and resistance to creep of magnesium alloys can be improved through the incorporation of RE elements [15]. Several Mg–RE binary systems, such as Mg–La [16,17], Mg–Ce [16,18,19], Mg–Nd [16,20,21,22], Mg–Sm [23,24], Mg–Dy [25], Mg–Y [26,27], Mg–Er [28,29], and Mg–Yb [30], have been recently developed.

It is known that RE elements are included in the majority of high-strength alloys. This review focuses on the Mg–Gd system using different alloying elements among the multiple binary Mg–RE systems investigated over the past decade. In a magnesium matrix, the maximum solubility limit of gadolinium (Gd) (23.49 wt% at 548°C) is at the eutectic temperature and drops off sharply as the temperature drops (5 wt% at 250°C). By adding yttrium (Y) into the Mg–Gd system, the solubility limit of gadolinium in the Mg matrix is decreased. Because Gd and Y have comparable atomic radii, Y replaces the Gd element in the Mg–Gd–Y system [31]. Lattice distortion was generated because gadolinium’s atomic radius (180 ppm) is bigger than Mg’s (150 ppm). Since Gd has a bigger atomic radius than Al and Zn, it produces a stronger solid solution strengthening effect [32]. Thus, the strength of Mg–RE alloy is superior to that of common Mg–Al and Mg–Zn alloys. In Mg–7Gd–xNd–0.5Zr alloys (x = 0, 0.5, 2, and 3.5 wt%), the tensile strength is enhanced at both room and elevated temperatures when Neodymium (Nd) is added at concentrations ranging from 0.5 to 2 wt% [33]. The addition of a large amount of Nd facilitates the precipitate phase’s development and enhances grain refinement [34]. Aluminum, when added in smaller amounts to Mg–Gd alloys, has a similar effect of refining the grains [35,36]. Micro-alloying with Cu increases the tensile and compressive yield strengths, whereas it decreases the Mg–Gd–Y–Zn alloy’s corrosion resistance [37]. Once zinc (Zn) is added to the Mg–RE alloy system, the long period stacking ordered (LPSO) phase is formed, which increases in the alloy’s strength both at room temperature and at higher temperatures [38,39]. Zirconium (Zr) and manganese (Mn) were added to refine the grains Mg–Gd system [40], improving the alloy’s mechanical properties [41]. The addition of Sn to the Mg–Gd–Y system ensures that the β′ precipitates are evenly distributed throughout the Mg-matrix [42]. The Mg–Gd alloy system’s attributes are not only enhanced by alloying additions; the alloy’s microstructure and properties are also significantly affected by processing methods. The objective of this article is to provide an overview of the effects that the various processing procedures that were employed to process the Mg–Gd–X alloy had on the microstructure and properties of the alloy.

2 Processing of Mg–Gd–X alloy

The Mg-based alloy is mostly produced through casting at the moment. The casted billets were subjected to various bulk deformation procedures. Extrusion and rolling, two important bulk deformation methods, were utilized for making the Mg–Gd alloy. In addition, the Mg–Gd alloy can also be processed via forging, pressing, and drawing. The most typical method for processing Mg-based alloys is casting, followed by heat treatment. Castability, molten metal reactivity, and grain structure control are essential considerations when melting magnesium alloys [55]. While heating magnesium alloy to high temperatures in an air atmosphere, it will burn and oxidize. To prevent this, a protective gaseous medium such as air/SF₆, air/CO₂/SF₆, or CO₂/SF₆ can be employed for melting, holding, and casting [2,33,56]. Mg–7.8Gd–2.7Y–2.0Ag–0.4Zr (wt%) alloy was made by conventional metal mold gravity casting, and the mechanical parameters of the fabricated alloy are provided in Table 1 [43]. Increasing the mold temperature will initially increase the fluidity (720–750°C) and then reduce the fluidity (750–780°C) of the molten metal, whereas increasing the pre-heat temperature of the mold will improve the fluidity and coarsening of the grains [57]. The two most important process parameters for squeeze casting are the pouring temperature and the applied pressure. By utilizing the squeeze casting technology, which integrates the casting and forging processes into one, Mg–10Gd–3Y–0.5Zr was produced with pressures ranging from 0.1 to 160 MPa [58]. Squeeze casting and high-pressure die casting (HPDC) both use pressure during the process, but it is important to apply the pressure before the melt solidifies. The yield strength and hardness of the as-cast, solution-treated (T4), and age-hardened (T6) alloy were enhanced by increasing the pressure applied during the squeeze-casting process [59]. The Mg–15Gd–1Zn–0.4Zr alloy’s grain size was decreased from 41 to 28 μm (0–6 MPa) due to increased heat transfer coefficient while applying pressure [60]. Grain sizes ranging from 3 to 10 μm are typical for the Mg–RE alloy that is fabricated using HPDC [61]. The vacuum die cast Mg–8Gd–3Y–0.5Zr also has grains with sizes greater than 20 μm, while some of the grains are less than 10 μm [62].

As a metal-forming process, extrusion involves forcing metal into a die in order to obtain a specific shape. Concurrently, the material is wound onto one or more rolls in a certain order during the rolling process. Driven forward by the frictional force between the material and the rolls, the material undergoes substantial grain refining, leading to the desired improvement in mechanical properties. Hot and cold extrusion/rolling were the two main categories of this procedure according to the temperature at which it was carried out. Both hot and cold extrusion/rolling are carried out at temperatures higher than the recrystallization temperature of the material, respectively.

The microstructure and mechanical characteristics of the Mg alloy are greatly impacted by the extrusion parameters, which include extrusion ratio, temperature, and speed [63,64]. Prior to extrusion at 400 and 450°C at 0.1 mm/s with extrusion ratios of 10 and 20, the cast Mg–8Gd–4Y–1Mn–0.4Sc (wt%) alloy was homogenized at 520°C for 18 h [65]. Rotary forward extrusion (RFE) was used to develop a Mg–13Gd–4Y–2Zn–0.4Zr sheet. RFE combines torsion and forward extrusion, and it successfully reduces grain size and achieves significant grain refinement by increasing the number of RFE revolutions [66]. Better grain refinement from the material’s outer to interior regions is demonstrated by the higher rotating speed during the rotary backward extrusion process [67]. Grain refinement in the microstructure is also greatly affected by rolling parameters, just as it is by extrusion. To obtain a fine-grained microstructure, the homogenized alloy was heated above its recrystallization temperature and passed between two rollers [68]. Molybdenum disulfide lubricant used in nine cycles of multi-directional forging (MDF) at deformation temperatures of 350 and 400°C refines the microstructure of the Mg–8.59Gd–3.85Y–1.14Zn–0.49Zr alloy [69]. Grain sizes in the sub-micron and nanometer range can be achieved by the extreme deformation method known as equal channel angular pressing (ECAP) [70]. An outer arc angle of 0° and a die angle of 90° were used in multi-pass ECAP. For the two consecutive passes, route Bc was utilized. In addition, to guarantee a successful extrusion operation, the samples were heated to 360°C and kept in the die for 15 min before the first pass and after each of the four passes [71]. The bimodal grain is a characteristic of bulk deformation, in contrast to the uniform grain size that results from the casting process. Thus, the grain size of the alloy is a source of influence on the strength of the alloy, and the processing method can be used to control the grain size. The mechanical parameters of the alloys produced with each method are given in Table 1. Casting is the most common technique for processing magnesium alloys and composites. According to our review, the most important factors that affect the microstructure of the cast alloy are the pouring temperature, the mold/die preheat temperature, and the holding pressure. Using mechanical stirring while stir casting allows for the generation of a strong shearing force. When used in large quantities, the HPDC method produces components with good dimensional accuracy and a near-net form. Despite its advantages, the HPDC technique leaves the casted alloy with porosity and hot ripping behavior. Squeeze casting has lower porosity and a few benefits over HPDC, including less shrinkage and a lower hot tearing propensity caused by external pressure. In order to achieve grain refinement and dynamic recrystallization (DRX), the technique employs secondary plastic deformation procedures. The basal slide of the Mg causes edge cracks to appear during the cold rolling process. Metal is rolled through fixed profiles in rolling, but complicated profiles can be created in extrusion by designing the die. To prevent surface cracks produced by oxidation and hot shortness, the billet temperature must be properly managed during the hot extrusion process. The appropriate processing approach can refine the grains and generate a large fraction of stable β′ precipitates, allowing for the development of a high-strength alloy.

3 Microstructure of Mg–Gd–X alloy

The casted alloy comprises a dendritic α-Mg matrix, a eutectic structure at grain boundaries, and an ellipsoid phase within the α-Mg matrix [72]. A eutectic structure at the grain boundary and several block-shaped precipitates inside the grains are depicted by yellow and blue arrows, respectively, along with the energy dispersive spectrometer analysis done at the eutectic phase in the as-cast microstructure of the Mg–8.08Gd–2.41Sm–0.30Zr alloy in Figure 2(a) and (b) [73]. In its as-cast state, the β′ precipitate has a 25 nm diameter and a thickness below 15 nm [74]. The second phase was dispersed into the matrix using solution treatment, which created a super-saturated solid solution [75]. After the solution treatment, the non-equilibrium second phase found at the grain boundary was completely dissolved into the α-Mg matrix [56]. The sequence of precipitate formation for Mg-Gd alloy was Super-Saturated Solid Solution (SSSS) → GP zones (D0₁₉ type) → β′ (c-bco, Mg₇Gd) → β₁ (D0₃, Mg₃Gd) → β (fcc, Mg₅Gd) [76,77,78]. The petal-like β″ phase is formed by these precipitates in the early stages of age hardening in Mg–10 Gd alloy, and they serve as a nucleation site for the β′ phase [79]. The improved tensile strength of the peak-aged Mg–Gd–Y–Zn–Zr–Nd alloy is a result of the oval-shaped β′ phase and the precipitates that form within it [80]. Grain size increased from 46 to 68 µm, which was bigger in the as-cast state, and the alloy mostly comprises fine precipitates in the grain boundary and matrix after the age-hardening process [81]. The β′ precipitate’s 1.11 nm interplanar spacing and excellent coherent contact with the matrix enhance the mechanical characteristics [33]. The alloy’s strength is affected by the size and shape of the precipitates; the hardest precipitates were plate-shaped ones that formed in the prismatic plane [82]. These β′ precipitates greatly limit the mobility of dislocation slip in the various planes [83]. The formation of needle-like and block-precipitated phases of Mg–Gd–Zn–Zr alloy after the T6 treatment is shown in Figure 3 [84].

Figure 2

(a) The optical micrograph of As-Cast Mg–8.08Gd–2.41Sm–0.30Zr alloy and (b) energy dispersive spectroscopy [73].

Figure 3

SEM images of the T6-treated (a) Mg–3Gd–1Zn–0.4Zr, (b) Mg–6Gd–1Zn–0.4Zr, and (c) Mg–9Gd–1Zn–0.4Zr alloy [84].

Dark intermetallic particles of irregular size and shape, known as the LPSO phase, are formed when zinc is added to the Mg–RE system [38]. Furthermore, the Mg₁₂–RE–Zn phase was produced when the Gd element was substituted with Y in the Mg–Gd–Y–Zn alloy [85]. The optical, scanning electron microscopy, and transmission electron micrographs of the multi-directionally forged Mg–6Gd–4Y–0.7Zn–0.4Zr–0.3Ag alloy are displayed in Figure 4 [86]. Upon solidification of the Mg alloy, eutectic, block, and lamellar-shaped LPSO phases were produced [87]. Mg₂₄(Gd, Y, Zn)₅ and LPSO phases at the edges of the α-Mg grains also had lamellar phases, and in the inner portion of the α-Mg grain, some lamellar phases have been observed [88]. Even after the solution treatment, the matrix still retained a significant amount of LPSO. Rather, the LPSO phases were disrupted and scattered randomly throughout the matrix [89]. The solidification process causes the 18R-LPSO to form at the grain boundary, whereas the heat treatment (solution treatment and age hardening) process causes the formation of 14H-LPSO [90]. The grain boundary allowed the 14H-LPSO phase to penetrate into the grain interior. In the final results of high-resolution transmission electron microscopy analysis, the structure of the 14H-LPSO phase is ABABCACACACBAB, which is composed of 14 closely spaced planes that are 3.76 nm apart [91]. Furthermore, certain LPSO phases fractured and became dispersed along the grain boundaries after the extrusion procedure; these fragmented small phases impede the motion of dislocations. Additionally, it establishes a robust interface between the LPSO phase and α-Mg, preventing the formation of any additional voids or cracks during the plastic deformation process [92]. On the contrary, the forged alloy exhibited fragmented block and lamellar LPSO phases within the grains, which served as a source of micro-cracking and stress concentration [93].

Figure 4

(a) Optical, (b) and (c) SEM, and (d) TEM images of multi-directionally forged Mg–6Gd–4Y–0.7Zn–0.4Zr–0.3Ag alloy [86].

Plastic deformation processes, including extrusion, rolling, forging, and ECAP, refine the grains using DRX [85]. An extrusion temperature exceeding 450°C facilitates the full maturation of DRX granules and increases the grain size, thereby enhancing the ductility of the alloy [94]. A substantial reduction in the average particle size and the production of fine DRXed grains are the outcomes of increasing the number of passes [95]. Increasing the percentage of thickness reduction during rolling also results in a higher volume fraction of DRXed grains, as illustrated in Figure 5 [96]. Furthermore, as the quantity of strain hardening increases, the average grain size decreases, and LPSO phases are observed between the deformed grains [97]. The alloy is composed of fine DRXed grains and coarse un-DRXed grains in the extrusion direction following the hot extrusion. The un-DRXed region demonstrates a robust texture, while the DRXed region manifests an attenuated texture as a result of the random orientation of the DRXed region [98]. In the same way, a bimodal structure is composed of DRXed grains that are arbitrarily oriented within the grains and at the grain boundaries/twins; this provides further evidence of the rolling process’s continuous DRX [99]. Furthermore, forged magnesium alloy exhibits a bimodal structure comprising refined grains (49.6%) and irregularly oriented coarse grains (50.4%) [100]. The texture can be strengthened by unDRXed granules, and the increased dislocation count is the cause of the material’s enhanced strength. Concurrently, the DRXed granules are inflexible regarding the activation of the basal slip. Thus, unDRXed regions contribute more to the material’s strength than DRXed regions [101]. Along the extrusion direction, the majority of the second phase was disrupted and destroyed during hot extrusion; some equiaxed grains are visible behind the disrupted grains as a result of DRX [102]. Likewise, the hot extrusion process transforms a negligible number of bulk particles into fine particles [103].

Figure 5

EBSD IPF maps, pole figures, and inverse pole figures of Mg–9Gd–4Y–1Zn–0.8Mn alloy rolled with different reduction percentages: (a and d) 27.5%, (b and e) 40%, and (c and f) 56.3% [96].

After hot extruding Mg–6.0Gd–1.0Y–0.5Zr alloy at 355°C and promptly cooling it with water to facilitate further precipitation and dynamic development of the precipitates by cold rolling with various reduction percentages [104], increasing the reduction percentage increases the number of twins and weakens the texture. After rolling at room temperature, edge cracks are observed [105]. Cold rolling increases the high-density dislocations, and most grains, both DRXed and un-DRXed, have a narrower distribution of orientations [106]. Furthermore, low-angle grain boundaries, lattice deformation, and flaws were created as a result of the cold-rolling [107]. To avoid the creation of edge cracks and increase plasticity after rolling, the Mg alloy was forged or extruded before rolling [108]. As evidence of cold working, certain un-indexed dark areas shear bonds with significant dislocation density were observed [109]. Afterward, the hot extrusion procedure produced a few dynamic precipitates. These precipitates are responsible for the alloy’s improved mechanical characteristics following aging treatment [110]. Refined grains play a vital role in the mechanical and corrosion behavior of the Mg alloy [111]. Formation of the second phase and precipitates will hinder the dislocation movement, improving the alloy’s mechanical properties. The high-strength alloy can be developed by achieving the refined grains and formation of high fraction of stable β′ precipitates. In addition, DRX and twins also play a vital role in the improvement of the mechanical properties of the alloy.

4 Mechanical properties of Mg–Gd–X alloy

An age-hardening procedure, i.e., solid solution strengthening at high temperature followed by hot water quenching and further aging at low temperature, can improve the strength of the Mg–RE alloy [4]. The hardness and tensile strength of the Mg alloy are significantly improved following the age-hardening process [112]. The age-hardened Mg alloy has a much higher Vickers micro-hardness than the as-cast or as-extruded alloy. Increased ageing temperature reduces the time required to obtain peak ageing and peak hardness value [43,113]. The micro-hardness value improves significantly during the hot extrusion process and age hardening of the extruded alloy due to precipitate formation [114]. After the age-hardening process, the characteristics are improved due to grain refinement, the production of nano-scale β′ precipitates, and dynamic precipitates [115]. The strength of the alloy improves as the aging duration increases, while the elongation% decreases [116]. In comparison to the as-cast and T6-treated alloys, the solution-treated alloy had higher ductility [117]. Pyramidal slip activation and nano-scale precipitate action are closely correlated with the alloy’s ductility [118].

At room temperature, activation of the basal slip {0 0 0 1} <1 1 2 0> is easier when compared to the non-basal slip {1 0 1 0} <1 1 2 0> and {1 1 2 2} <1 1 2 3>, because of the value of the critical resolved shear stress [32,119]. Twining is required for grain deformation in Mg alloys due to the less-active slip systems, and these twins are regarded as a barrier to basal slip. As a result, stress concentration may form at the interface between the twins and the matrix, giving nucleation energy for precipitation [120]. Many dislocation twins are observed inside the grains along with the metastable βʺ and β′ phases after the iso-thermal aging process. These precipitates hinder the dislocation movement [121,122]. The extruded alloy has been stretched, resulting in an elongation of 27.4% and a tensile strength of 361 MPa due to the refined grains (4.34 µm) [123]. Because of the high volume proportion of β′ precipitates, the Mg–Gd–Er–Zr alloy has a greater strength of 490 MPa after T5 treatment [124]. These β′ precipitates efficiently stop the basal plane slip and have an efficient dispersion hardening effect on the matrix [125]. Adding alloying elements can enhance the alloy’s strength by improving its aging response and creating β′ precipitates and nanoscale plate-like precipitates [45,126]. Due to the high volume percentage of the LPSO phase, the T6-treated alloy exhibits great strength and ductility, and the age-hardened alloy displays the 14 H LPSO phase and Hʹ precipitates [127]. The age-hardening condition forms the LPSO phase, which restricts plastic deformation and fracture propagation [74]. The alloy’s mechanical properties have been enhanced and the grain growth is restricted during the DRX process due to the pinning effect of the β′ precipitates [128]. According to the Hall-Petch relationship, the alloy’s yield strength was 360 MPa adhering to the fourth MDF pass [129]. The DRXed grains, which are 1–2 µm in size, enhance the yield strength by reinforcing the grain boundaries, while the subgrains in the unDRXed grains limit the migration of dislocations and boost the alloy’s strength in tensile tests [130]. Grain refinement during DRX causes the elongation of the multi-directionally forged alloy to reach 26% after the fourth pass [131]. Directionally solidified alloy shows an excellent elongation (109.55%), which reflects a better work-hardening effect [132]. Due to the strain-strengthening effect, the ultimate tensile strength increased by 30.1% and the yield strength by 52.1% [133]. Figure 6(a)–(c) shows the tensile test results of the as-cast, T4, and T6-treated Mg-15Gd alloy at room and elevated temperature, and Figure 6(d) shows the tensile test results of the as-cast and heat-treated alloy at 260°C [134].

Figure 6

Tensile properties of Mg–15Gd alloy at different temperatures: (a) As-cast, (b) T4 condition, (c) T6 condition, and (d) comparison of tensile properties at 260°C [134].

At higher temperatures, the alloy’s strength decreased but its ductility rose due to dislocation pinning and grain boundary sliding [135]. Segregation near the grain boundary, on the other hand, results in fracture initiation due to the softening of the segregated compound at high temperatures (500°C) [136]. The fraction of low-angle grain boundary grows at high temperatures, and the alloy has undergone slip deformation and twinning [137]. The fracture morphology of the tensile test specimen includes transgranular and intergranular fractures, as well as tear ridges and a cleavage plane, which indicate brittle fracture [138]. At larger stress or strain amplitudes, the alloy’s fatigue life can be reduced, and the alloy will undergo cyclic hardening, cycle softening, and cyclic stability under various testing conditions [139]. Fatigue cracks were started within the grain at the slip band and propagated to neighboring grains by secondary crack coalescence [140]. Mg–12Gd–3Y alloy has higher fatigue strength than AZ80 alloy, and heat treatment improves the alloy’s fatigue strength [119]. The fatigue strength of the extruded and peak-aged Mg–10Gd–3Y–0.5Zr alloy is 165 MPa due to the precipitates present in the plane {2 1 1 0}, which strongly restricts the basal slip [32]. The creep mechanisms transfer from diffusion and grain boundary sliding mechanisms to slip and dislocation climbing as the creep test temperature rises [34]. Mg–Gd–Zr alloy shows that T4- and T6-treated alloys have excellent creep resistance compared to the as-cast alloy. In addition, the T6-treated alloy shows a better creep resistance (below 280°C) due to the stable β′ precipitates with good thermal stability [141]. Grain size and the presence of second-phase particles have altered the mechanical properties of the Mg–Gd–X system.

5 Corrosion behavior of Mg–Gd–X alloy

By utilizing a potential difference (E _cor) between the reference and working electrodes and the corrosion rate (I _cor), potentiodynamic polarization is one of the electro-chemical approaches that may be employed to forecast the material’s corrosion behavior. Corrosion resistance is improved when the I _cor value is minimal and the E _cor value is large; Table 2 lists the I _cor and E _cor values for Mg alloys that were tested with different solutions. Two reactions, the cathodic polarization and the anodic polarization, stand for the evolution of hydrogen (H₂) and the creation of monovalent mg ions (Mg²⁺), respectively [142]. The general corrosion mechanism of Mg alloy has been summarized in equations (1)–(4) [143,144]. Figure 7 shows the surface that has corroded in the 3% NaCl electrolyte solution, which is the result of adding Gd (0–2 wt%), which increases the corrosion rate [145]. Furthermore, when Gd was added to the Mg–2Zn–0.4Mn–0.1Sr alloy, the current density of corrosion increased [146]. Mg–8Gd–3Y–0.4Zr alloy forged with a 12% reduction results in a lower I _cor value and a positive movement in the E _cor value when contrasted with the extruded alloy [147]. Furthermore, when polarization is performed in the 3% KCl solution, the annealing temperature is closely related to the corrosion current [148]. The corrosion current density (I _cor) value and corrosion rate decrease with increasing extrusion temperature in the typical three-electrode corrosion investigation of the simulated body fluid (SBP, up to 360°C) [7].

Figure 7

The local observation corrosion surfaces: (a) Mg–2Y–1Zn, (b) Mg–2Y–1Zn–0.5Gd, and (c) Mg–2Y–1Zn–1Gd₁; (d) Mg–2Y–1Zn–2Gd [145].

The magnesium alloys are characterized by their refined grains, low corrosion resistance caused by their dense grain boundaries, and the passive coating that forms along those boundaries [149]. During corrosion, the second phase at grain boundaries acts as a cathodic site and forms a galvanic couple with the matrix [39]. Surface morphology that has been corroded exhibits deep pits caused by corrosion at the β phase, and these pits grow as the β phase is distributed [150]. Micro-cracks and compact and uniform corrosion products are observed at the grain boundary as evidence of galvanic corrosion [151]. To increase the material’s resistance to corrosion, heat treatment is used to dissolve and disperse the particles of the second phase into the matrix [152]. The corrosion resistance of the alloy is improved when an alloy is subjected to a solution treatment when the eutectic phase and β phase will dissolve into the matrix [143]. The alloy’s corrosion behavior is significantly affected by the second phase of formation, which happens during heat treatment (T6). The galvanic corrosion begins with the precipitates acting as a cathode and the passive layer that prevents the corrosion from spreading and slows it down is created in the second phase [153]. Homogenous and uniform distribution of β′ phase into the matrix will improve the corrosion resistance of the alloy and vice versa [154]. Because fewer second-phase precipitates cause galvanic corrosion, corrosion resistance could be reduced with a shorter aging time and low temperature [155]. Additionally, galvanic corrosion is reduced due to decreased LPSO phase following T6 treatment [156]. On the other hand, the α-Mg grains have allowed the formation of the LPSO phase, which limits the movement of ions into the matrix. In addition to protecting the matrix from galvanic corrosion, these phases served as anodes in conjunction with the (Mg, Zn)₃Gd phase [157]. In the second stage, achieving a balance between grain size and area fraction can improve corrosion performance [158]. The subgrains with a high dislocation density in the DRXed zone function as an anode in the corrosion process [128]. Major corrosion products of the Mg–Gd alloy are Mg(OH)₂ with a small amount of Gd(OH)₃, and the film thickness of the as-cast and T4-treated sample is 2 µm and T6-treated sample is 15 µm [159]. These corrosion product films restrict the corrosive ion penetration and corrosion process [160].

Environmentally assisted cracking in Mg alloy relies on hydrogen, which can cause hydrogen embrittlement if it diffuses into corrosion pits [161]. Additionally, the stress corrosion cracking (SCC) test conducted in 3.5% of NaCl medium results shows that UTS is reduced to 56, 40, and 56 MPa in As-cast, T4, and T6 conditions, respectively [143]. Hydrogen embrittlement causes microcracks to start in the surface’s corrosion pits and spread via transgranular cracking [162]. The alloy’s corrosion resistance is enhanced when its surface is altered using a continuous wave laser, which introduces a dissolving β phase into the molten region and decreases the quantity of a local galvanic couple [163]. To enhance the corrosion resistance of the Mg alloy, surface modification techniques such as anodizing, advanced processing control (such as heat treatments and semisolid processes), and coatings can be employed [164]. Nonetheless, the Mg–Gd alloy’s corrosion resistance has been the subject of a great deal of research. However, additional research is necessary to comprehend how the material acts in real-world atmospheric settings.

6 Conclusions

The microstructure and mechanical properties of the Mg–Gd–X alloy, which have been subjected to different methods of processing by multiple groups of researchers, are reviewed in this article. As a result of the review that was carried out, we were able to draw a few conclusions here.

1. Various casting procedures and deformation techniques, including rolling and extrusion, are used to develop the Mg–Gd–X alloy. The Mg–Gd–X alloy undergoes several heat treatment processes, including T4, T5, T6, and T8. Alloys with high strength and good ductility are created by combining any one of the processing techniques with a suitable heat treatment process.

2. The selection of processing procedures is essential because the microstructure and formation of second-phase precipitates can be influenced by the processing techniques that are used. Processing methods resulting in submicron or nanometer grain sizes can produce high-strength magnesium alloys.

3. The formation of second-phase precipitates is crucial for developing high-strength alloys. The size, shape, distribution, and ratio of volume of these nano-scale precipitates influence the alloy’s mechanical characteristics and corrosion behavior. Plate-shaped, stable β′ precipitates in the prismatic plane can provide high strength.

4. The distribution of nano-level precipitates or LPSO influences the corrosion resistance of the magnesium alloy. The alloy’s corrosion resistance can be improved through heat treatment and the use of surface modification procedures. Although corrosion tests have been undertaken in various mediums, it is requested that the study be performed under real-world atmospheric conditions.

5. Researchers can focus on developing a high-strength Mg–Gd–X alloy with small RE additions, taking advantage of the plethora of resources available. The researchers can focus on fabrication methods such as ECAP, additive manufacturing (wire arc additive manufacturing, selective laser melting), and surface treatment (laser shock peening) where only limited studies are conducted.