Severe Plastic Deformation of Mg–Zn–Zr–Ce Alloys: Advancing Corrosion Resistance and Mechanical Strength for Medical Applications

Abstract

Magnesium-based alloys hold potential for medical applications, but face challenges like rapid bioresorption and limited mechanical strength during early bone healing. In our study, we present a novel Mg–Zn–Zr–Ce alloy with low cerium content (up to 0.1 wt.% Ce) processed using two severe plastic deformation (SPD) techniques. Through an innovative combination of multiaxial forging and multipass rolling, we have achieved a fine-grained structure with an average grain size of the primary α-Mg phase of 1.0 μm. This refined microstructure exhibits improved mechanical properties, including a substantial increase in yield strength (σ_YS) from 130 to 240 MPa, while preserving ductility. The alloy’s composition includes α-Mg grains, cerium and zinc hydrides, and intermetallic phases with cerium and zinc elements. Tensile testing of the fine-grained alloy demonstrates an enhancement in yield strength (σ_YS) to 250 MPa, marking a 2.8-fold improvement over the conventional state (σ_YS = 90 MPa), with a modest 2-fold reduction in ductility. Crucially, electrochemical tests conducted in physiological solutions highlight substantial advancements in corrosion resistance. The corrosion current was reduced from 14 to 2 μA/cm², while polarization resistance decreased from 3.1 to 8.1 kΩ∙cm², underlining the alloy’s enhanced resistance to biodegradation. Our results show that the novel Mg–Zn–Zr–Ce alloy, after combined SPD, demonstrates mitigated bioresorption and enhanced mechanical properties. Our findings highlight the fact that the introduction of this innovative alloy and the application of SPD represent significant steps towards addressing the limitations of magnesium-based alloys for medical implants, offering potential improvements in safety and effectiveness.

1. Introduction

One of the most promising directions for the application of magnesium alloys are aerospace engineering [1,2] and medicine [3,4,5]. The advantage of these alloys for use in medicine is their low values of elastic modulus, which correlate well with those of bone tissue, which, in turn, eliminates high stresses in the area of contact of a metal implant with bone. Additionally, magnesium has the ability to biodegrade, making it useful for certain types of orthopedic and vascular surgeries. Magnesium alloys have a low elastic modulus (10–40 GPa), close to the elastic modulus of cortical bone tissue [6]. The disadvantages of magnesium alloys include low corrosion resistance and hydrogen evolution during metal dissolution.

When adjusting the chemical composition of magnesium alloys for medical use, the biocompatibility of various alloying elements is taken into account, some of which are toxic (Cd, Be, Pb, Ba, and Th) or can cause allergic reactions (Al, Co, V, Cr, Ni, Ce, La, and Cu). Elements Ca, Mn, Zn, Sn, Sr, and Ce are considered favorable elements for medical use [7]. In this regard, alloys of the Mg–Zn–Zr system are the most promising.

The addition of Ce to the Mg–Zn–Zr system improves the yield strength, tensile strength, and ductility of the alloy due to grain refinement [8,9]. According to the binary phase diagram of Mg–Ce, the content of the solid solution of the Ce element in the matrix is extremely small; therefore, improvement of the properties of the Mg–Zn–Ce–Zr alloy occurs due to the formation of intermetallics within the grain. In addition, microalloying with a RE element (Ce or Nd) from 0.05 to 0.1 wt.% leads to the formation of smaller and spherical intermetallics, which promote the nucleation of crystallization centers and increase the ductility of alloys [10,11,12,13,14].

In some cases, the achievable level of strength properties of doped magnesium alloys does not always meet the required specifications, which limits their use in medicine. The problem of improving mechanical properties can be solved by refining the grain structure, using various methods of severe plastic deformation (SPD) [15,16], such as high-pressure torsion (HPT) [17,18], equal-channel angular pressing (ECAP) [19,20,21,22,23], multiaxial forging (abc-forging) [24,25], hydrostatic extrusion process (HE) [26], and other types of deformation methods (different types of extrusion, radial-shear rolling, etc.) [27,28,29,30]. This can significantly increase the structural strength of magnesium alloys by 2–2.5 times without alloying with toxic elements for the body by obtaining an ultrafine-grained (UFG) and/or fine-grained (FG) state. The HPT method has been effectively utilized to achieve an UFG structure state and enhance the mechanical properties of biocompatible magnesium alloys, such as Mg–Zn and Mg–Li–Y [31]. A scientific article [23] demonstrated that the application of ECAP treatment to a Mg–Zr alloy, under specific temperature and speed conditions, resulted in the creation of samples that exhibited an FG structure, with an average grain size of 8.6 µm and a remarkable maximum elongation of approximately 380%. This highlights the potential of these processing methods to optimize the structural and mechanical characteristics of magnesium alloys for biomedical applications.

ECAP and MDF are commonly used SPD methods because they are able to create a uniform, fine-grained microstructure in large bulk workpieces, unlike HPT. In previous studies [32,33], SPD of the WE43 alloy of the Mg–Y–Nd–Zr system was conducted by ECAP, abc-forging, and rotational forging to obtain a grain size of less than 1 μm. SPD resulted in an increase in the strength of the WE43 alloy by 40%. Reducing the grain size also had a positive effect on biocompatibility in in vitro experiments.

In addition, grain refinement in magnesium alloys is one of the most effective ways to regulate the corrosion rate due to the homogeneous distribution of nanoscale second phases and the formation of crystallographic textures during deformation. Hence, for the AZ31 alloy of the Mg–Al–Zn system, the best corrosion-resistant characteristics are achieved for the FG state compared to the coarse-grained (CG) alloy [34]. In work [35], it was demonstrated that a Mg–Y–RE alloy in the UFG state has improved polarization resistance and the most positive pitting potential, in contrast to the CG state.

Up until now, there has been an abundance of publications examining the structural and mechanical properties of magnesium alloys after SPD. Nevertheless, there are still unresolved challenges associated with developing deformation techniques that can produce grain sizes less than 1 μm to achieve optimal functional properties in biocompatible magnesium alloys. For instance, in study [31], the authors found that the HPT method fell short of achieving a uniform UFG state throughout the volume of the workpiece in Mg–Zn–Ca and Mg–Y–RE alloys due to localized deformation. A solution to this issue could be the implementation of two-stage deformation processing.

Furthermore, study [36] showed that rolling followed by ECAP can increase the mechanical strength, ductility, and corrosion resistance of pure magnesium. In similar research [37], it has been demonstrated that a combination of ECAP and rotary forging in the magnesium alloy Mg–Zn–Ca (ZX40) led to significant grain refinement and accumulation of dislocations. This, in turn, resulted in increased tensile strength and fatigue limit values of 380 MPa and 115 MPa, respectively.

The Mg–Zn–Zr alloy, with its corrosion-resistant properties and biocompatible alloying elements, is deemed a promising option for medical applications. Therefore, understanding the impact of SPD on the mechanical and corrosion properties of this alloy is crucial from a practical standpoint. While the multiaxial forging method is commonly employed to attain the UFG state in sizable volumetric samples, there is a dearth of studies focused on producing UFG states in Mg–Zn–Zr system alloys with Ce additions through SPD techniques [38,39,40]. In addition, despite the numerous advantages of magnesium alloys in medical applications, it is important to address their limitations, particularly in terms of mechanical strength. Strengthening these alloys is crucial to harness their full potential in medical settings. The Mg–Zn–Zr–Ce alloy emerges as a promising option for medical use, given its composition of biocompatible elements. The inclusion of cerium in small quantities is especially noteworthy, as it gives the alloy antibacterial properties. This feature is especially beneficial in the initial stages of implantation, helping to prevent infections and promote successful recovery. This study addresses this gap by showcasing the potential of producing bulk magnesium alloy workpieces with enhanced strength and corrosion-resistant characteristics using a composite SPD method.

The purpose of the work was to study the influence of the combined SPD method on the microstructure and mechanical and corrosion properties of a magnesium alloy based on the Mg–Zn–Zr–Ce system.

2. Materials and Research Methods

The object of the study was a commercial magnesium alloy of the MA20 grade based on the Mg–Zn–Zr–Ce system, developed by the Federal State Unitary Enterprise VIAM (Moscow, Russia). The alloy billet was cast by VILS (Russia, Moscow) in an argon atmosphere in a crucible furnace, followed by hot rolling at a temperature of 400 °C to a plate thickness of 30 mm. The rolling temperature of the alloy was selected to prevent any inhomogeneity of the alloy structure across the sections of the plate during deformation. The composition of the alloy (

Table 1. Elemental composition.

Alloy	Elemental Composition, wt.%
Mg–Zn–Zr–Ce	Mg—98,0, Zn—1,3, Zr—0,1, Ce—0,1, O—0,5

) was determined using a Niton XL3t XRF portable X-ray fluorescence analyzer (Thermo Scientific, Waltham, MA, USA).

Intensive plastic deformation was conducted by the method of multiple pressing in a press-mold [41] according to the scheme shown in Figure 1.

The temperature of pressing the alloy was 250 °C at a forging speed of 0.5 mm/s, to avoid the effects associated with dynamic recrystallization of the alloy sample and the growth of the average grain size. However, the plasticity of the alloy below the recrystallization temperature (<265 °C) decreases rapidly, which limits the accumulated deformation and eventually leads to the destruction of the sample. In this work, the number of pressings (3abc) was limited to three passes (with a reduction of 30% in height at each), which allows for a significant refinement of the microstructure, compared to the initial state of the alloy. The logarithmic true strain was calculated for each stage in accordance with Equation (1).

$$e=ln\left(\frac{h_1}{h_2}\right),$$

(1)

where:

e—logarithmic deformation;

h₁—initial height;

h₂—height after pressing.

The accumulated logarithmic strain value for all stages of forging is e = 1.1. After each pressing, the sample was removed from the die, rotated 90 degrees, reinserted, and revolved 90° in the same direction around its longitudinal axis. This ensured the uniform structure of the sample.

For further improvement of the properties, the alloy was subjected to multi-pass rolling at a temperature of 250 °C and a speed of 25 mm/s, which increased the accumulated strain value of the sample to e = 1.5. The appearance of the samples at each stage of deformation is shown in Figure 2.

The microstructure and chemical composition were studied using optical microscopy (Altami MET 1 MT microscope; Altami, St. Petersburg, Russia) and transmission electron microscopy (JEOL JEM 2100 microscope; Tokyo Boeki Ltd., Tokyo, Japan). Using the electrical discharge cutting method, plates with a thickness of 0.3 mm were cut from the samples and were mechanically thinned to a thickness of 0.1 mm using abrasive paper with a grit of P2500. Further sample preparation was conducted by ion thinning using the JEOL Ion Slicer EM-09100IS (Tokyo Boeki Ltd., Tokyo, Japan). Miller indices were calculated using the Odpin software. To calculate the average size of structural elements (grains, subgrains, and fragments), the secant method outlined in [42] was used. The determination of the average grain size was carried out according to GOST 21073.2-75, using 100 measurements of three images of each microstructure of the alloy.

X-ray diffraction (XRD) analysis was performed using diffraction patterns obtained with a DRON-8N X-ray diffractometer (Burevestnik, St. Petersburg, Russia) using copper cathode radiation, using a symmetrical shooting scheme in angles from 20 to 40° with an exposure of 120 s with a step of 0.05°. The comparison of interplanar distances was performed using the ICDD PDF-4+ (2019) database.

Mechanical properties were evaluated by Vickers hardness (Duramin-5 microhardness tester; Struers, Ballerup, Denmark), yield stress σ_0.2, tensile strength σ_UTS, and elongation δ determined from tensile and compression tests (UTS-110M-100 testing machine; Test-Systems, Ivanovo, Russia). Flat “dog-bone” shaped samples were used for tensile tests (dimensions of the central part: thickness—3.0 mm, width—2.7 mm, length—9.0 mm), and cylinders with a height-to-diameter ratio of 2:1 were used for compression tests (height—13.0 mm, diameter—6.5 mm). Samples for all tests were cut using electrical discharge cutting followed by abrasion with P2500 grit abrasive paper.

To assess the corrosion resistance of the alloy in different states, an immersion method was used in a PBS solution at a temperature of 37 °C. Prior to the test, each sample was polished with P2500-grit SiO₂ abrasive paper to obtain the same surface quality (Ra roughness value of 0.5 ± 0.1 μm). Five samples were selected for each alloy condition. The mass of the samples was measured every 48 h on Vibra XFR-125E analytical scales (Shinko Denshi, Tokyo, Japan), followed by recording of the surface condition. The solution in the cells was replaced every 96 h, and the total experiment time was 384 h. The corrosion rate was calculated using Equation (2).

$$CR=\frac{\mathsf{\Delta }m·K}{\rho ·A·t},$$

(2)

where:

CR—corrosion rate (mm/year);

Δm—mass difference (g);

K—constant (87600);

ρ—Mg density (g/cm³);

A—Area in contact (cm²);

t—duration of experiment (hours).

The electrochemical measurements were conducted with a P-40X potentiostat–galvanostat coupled with an FRA-24 frequency analyzer module (Electrochemical instruments, Chernogolovka, Russia). All experiments were conducted in an E-7SF three-electrode cell (Electrochemical instruments, Chernogolovka, Russia) at room temperature. SBF electrolyte was used with a graphite counter electrode and a 4.2 M silver chloride reference electrode. The cell featured a fixed sample window, providing a constant area of 1 cm² for electrolyte–sample interactions. A gold-plated textolite served as a current collector. Potentiodynamic curves were drawn at a sweep of 10 mV/s across a range of ±250 mV of the open circuit potential (OCP) for each sample. Prior to the start of the experiment, each sample was kept in the cell for 60 min to ensure stabilization of all processes at the electrolyte–material interface. To determine the corrosion characteristics, Tafel lines with a correlation coefficient of at least 0.95 (±0.05 V from the E_corr value) and an allowable potential error of no more than 1 mV were plotted in the supplied ES8 software (Electrochemical instruments, Chernogolovka, Russia). The polarization resistance (R_p) was computed using the Stern–Giri Equation (3):

$$R_p=\frac{{\beta }_a·{\beta }_k}{2.3j_{corr}·\left({\beta }_a+{\beta }_k\right)},$$

(3)

R_p—polarization resistance, Ohm;

β_a—anode straight slope, V/dec

β_k—cathode straight slope, V/dec

j_corr—calculated value of corrosion current density, A/cm².

3. Results and Discussion

Figure 3a shows an optical image of the microstructure of the Mg–Zn–Zr–Ce alloy in the initial CG state. The microstructure of the alloy is represented by equiaxed grains of the main α-phase based on the magnesium solid solution. The average grain size was 25.0 μm (Figure 3a). Inside the grains, there are a large number of particles, mainly spherical in shape, and areas corresponding to accumulations of particles.

The 3a,b and c pressing deformation of the magnesium alloy resulted in a significant reduction in the average grain size. According to optical metallography data, after 3abc pressing, the alloy formed a fine-grained (FG) state with an average α-phase grain size of 3.0 µm (Figure 3b). Most grains exhibited an irregular shape and were elongated along the deformation direction. Further deformation through rolling contributed to additional refinement of the FG microstructure, reducing the average grain size of the primary phase to 1 µm (Figure 3c).

In Figure 4, SEM images of the alloy are provided for the initial state and after deformation. In the as-received (CG) condition, SEM images reveal fine intermetallic phase precipitates within and along grain boundaries. These precipitates are observed as individual particles and colonies formed by “agglomerated” particles (Figure 4a,b). Smaller individual intermetallic inclusions exhibit a close-to-spherical shape, with sizes less than 1 µm. In localized areas within the microstructure images, larger particles with sizes up to 7–8 µm are present (Figure 4b). According to X-ray spectral analysis, the aforementioned inclusions of various morphologies consist of compounds containing Mg, Zn, and Ce. In certain particles and colonies, the zinc content ranges from 14.2 to 18.2 wt.%, and cerium content ranges from 6.1 to 16.42 wt.%. Isolated particles with high cerium content (up to 97 wt.%) and sizes ranging from 2 to 10 µm are also observed (Figure 4b). It is noteworthy that cerium does not form an independent phase in the Mg–Zn–Zr–Ce system, except for the Ce₂H₅ hydride [43].

Following 3abc pressing and combined deformation (3abc pressing and rolling), SEM images reveal intermetallic phases (Figure 4c–f) that are enriched with Zn and Ce, as in the initial state. Elevated levels of zinc and cerium with varying concentrations may indicate the presence of different types of intermetallic phases. In the microstructure of the alloy subjected to 3abc pressing, several types of inclusions are present: fine spherical particles with an average size of 400 nm, colonies of fine crystalline particles, and individual Ce₂H₅ particles with sizes ranging from 9 to 20 µm. The maximum zinc content in these particles and colonies is 18.2%, while cerium content reaches 16.4 wt.%. In localized areas, inclusions with high Zr content up to 19 wt.% were identified (Figure 4d), which may indicate the presence of zirconium hydrides (ZrH₂). The main factor contributing to the presence of zirconium and cerium hydrides is the purity of the raw materials in the alloy blend, which determines the amount of dissolved hydrogen in the alloy. The morphology of inclusions in the alloy after 3abc pressing and rolling remains unchanged. SEM images show spherical particles with sizes ranging from 0.2 to 4.2 µm and elongated colonies of particles (Figure 4f,g). The spherical particles contain zinc within a concentration range of 11.2–16.3 wt.% and cerium within the range of 6.1–9.8 wt.%. In colonies, the average zinc and cerium content can reach up to 22.6 and 22.2 wt.%, respectively.

Figure 5 displays the diffraction patterns of the alloy captured in various states. In all states, the alloy predominantly consists of the α-Mg phase.

The alloying elements in this alloy (Zn, Zr, Ce) have limited solubility in magnesium, resulting in the formation of several solid solutions of Mg (Zn, Zr, Ce). It was previously mentioned that the alloy contains a certain amount of zirconium hydride (ZrH₂) and cerium hydride (Ce₂H₅), which is confirmed by the X-ray phase analysis. Additionally, the alloy contains a small amount of a triple intermetallic phase of the Zn_xCe_yZn_z type, the exact identification of which is challenging due to its complex lattice structure. Qualitatively, there is no change in the phase composition when the structural state is altered. The lattice parameters of the α-Mg phase change as the degree of deformation of the alloy increases. 3abc deformation reduces the c Mg parameter to 0.5208 nm compared to the initial state of 0.5230 nm. The a Mg parameter changes only slightly. Further refinement of the alloy’s microstructure through rolling reduces the c Mg parameter to 0.5203 nm.

The SEM, XRD, and optical metallography data are in accordance with TEM examinations. In the bright-field microstructure images of the as-received Mg–Zn–Zr–Ce alloy, equiaxed grains of the primary phase based on the α-Mg solid solution (HCP lattice) and inclusions of secondary phases with various morphologies are observed (Figure 6a–c). The secondary phases were further identified through micro-X-ray spectral analysis of the elemental composition. In addition to the Ce₂H₅ hydride phase identified in SEM images during X-ray spectral analysis, individual particles of different morphological structures were observed: individual spherical particles with sizes ranging from 120 to 600 nm, which were present within magnesium α-grains (Figure 6a); particle colonies (Figure 6b); and large irregular fragments (Figure 6c). According to TEM data, the content of zinc in individual spherical particles ranges from 8.5 to 9.8 wt.%, and cerium ranges from 4.7 to 5.1 wt.%. The particle colonies observed in TEM images contain zinc in the range of 23.4-27.0 wt.% and cerium in the range of 15.2–18.2 wt.%. In large fragments, the average zinc concentration is 19.9 wt.%, and cerium is 8.9 wt.%.

After abc pressing, subgrains and irregularly shaped fragments form within the matrix grains. The structure is heterogeneous. A dislocation substructure forms within the fragments. Inside the matrix grains, spherical particles with sizes ranging from 100 to 600 nm and containing up to 26.2 wt.% zinc and 16.3 wt.% cerium are observed (Figure 7a,b), as in the initial state. In regions where particles aggregate, the concentrations of zinc and cerium are elevated. Large, elongated fragments (Figure 7c) with high zinc content (28.82–33.6 wt.%) and cerium content (28.4–28.5 wt.%) are noticeable in the structure. Figure 7d–f depict EDX maps of element distribution obtained from a section of the foil with these large fragments.

After the combined deformation process, a more pronounced grain–subgrain structure is observed, accompanied by a more intense microstructure refinement (Figure 8a). A developed dislocation substructure forms within the subgrains and fragments. Spherical particles and large fragments are also present alongside the matrix subgrains (Figure 8b,c). These fragments have lengths in the range of 1–4 µm and widths of 0.4–0.6 µm (Figure 8c). The sizes of spherical particles remain unchanged after the combined deformation process. In these spherical particles, the maximum zinc content is 38.6 wt.%, and the cerium content is 22.9 wt.%. In the large fragments, the zinc content remains within the range of 30.4–39.3 wt.%, and the cerium content is in the range of 27.1–27.9 wt.%.

Cerium in the alloy primarily exists in the form of a solid solution and triple phases of Mg–Zn–Ce. During cooling, the eutectic initially forms the matrix phase Mg. The Mg–Zn–Ce compound forms in the remaining liquid phase after the solidification of the α-Mg phase. Part of the Zn and Ce forms a magnesium solid solution, while the remaining Zn and Ce exist in the remaining melt. According to the research, elements Mg, Zn, and Ce in the Mg–Zn–Zr–Ce system can form four main intermetallic compounds, including (Mg 1-x Zn x) 12 Ce, (Mg 1-x Zn x) 10 Ce, Mg 7 Zn 12 Ce, and Mg 3 Zn 5 Ce [44,45].

The results of EDX analysis and TEM reveal that the dispersed compounds in the alloy in both the CG and FG (after deformation treatments in two modes) states have distinct morphologies and consist of magnesium, zinc, and cerium.

Figure 9 illustrates the deformation curves from static tensile and compressive tests for CG samples and FG samples subjected to 3abc pressing and 3abc pressing followed by rolling. Tensile tests show that after 3abc pressing, the alloy exhibits the following mechanical properties: a yield strength (σ_YS) of 140 MPa and an ultimate tensile strength (σ_UTS) of 220 MPa. It is evident that the FG alloy resulting from 3abc pressing achieves a yield strength 1.5 times higher and a tensile strength 1.2 times higher compared to the initial as-received (IC) state (σ_YS = 90 MPa, σ_UTS = 190 MPa), albeit with some reduction in plasticity to δ = 12%. After 3abc pressing and rolling, the alloy samples exhibit the highest mechanical properties, namely, σ_YS = 250 MPa and σ_UTS = 270 MPa, which are respectively 2.8 and 1.4 times higher than those of the CG alloy (Figure 9a). The significant strengthening of the alloy is a result of the additional refinement of microstructural elements due to the rolling process. It is noteworthy that the ultimate plastic deformation δ to failure significantly decreased to 3%, attributed to the strengthening effects of grain boundaries and substructures resulting from the two-stage severe plastic deformation (SPD).

It is worth mentioning that in compression tests, the failure of FG samples after 3abc pressing and combined SPD occurred at a deformation of 12%, slightly higher than that of CG samples (10%). At this level of deformation, the strength properties were as follows: σ_YS = 160 MPa and σ_UTS = 380 MPa for 3abc pressing, and σ_YS = 240 MPa and σ_UTS = 470 MPa for combined SPD (Figure 9b). After combined SPD, the alloy samples exhibited the highest mechanical properties in the compression tests, with σ_YS and σ_UTS values respectively 1.8 and 1.5 times higher than the specified values for the CG alloy.

The mechanical properties of the alloy are presented in

Table 2. Mechanical properties of the Mg–Zn–Zr–Ce alloy.

State	Average Grain Size, µm	σYS, MPa	σUTS, MPa	δ, %	Hμ, MPa	E, GPa
Initial, CG	25.0 ± 10.0	90/130 *	190/310 *	13/10 *	605 ± 30	10
3abc pressing, FG	3.0 ± 2.0	140/160 *	220/380 *	9/11 *	740 ± 30	12
3abc pressing and rolling, FG	1 ± 0.7	250/240 *	270/470 *	3/11.5 *	750 ± 40	12

* Compression test.

.

The formation of an FG structure in the alloy through the 3abc pressing and combined severe plastic deformation (SPD) method results in an increase in microhardness values, reaching 740 and 750 MPa, respectively. This represents a 1.2 times improvement compared to the CG state, where the microhardness is 605 MPa. It is noteworthy that the modulus of elasticity in the CG state of the alloy is 10 GPa, whereas in the FG state, it reaches 12 GPa. This modulus is comparable to the elasticity of bone [4].

The corrosion resistance of the Mg–Zn–Zr–Ce alloy was assessed using the immersion method. The results of the experiment were used to create graphs showing the mass loss of the samples (Figure 10). Additionally, an assessment of the surface condition of the samples was conducted every two days (Figure 11). These studies demonstrated that the alloy possesses high corrosion resistance compared to other magnesium alloys [46]. This is evident from the calculated corrosion rate (CR) of 0.09 ± 0.02 mm/year for the original alloy sample. In the case of deformed samples in a fine-grained state, the corrosion rate remains negative throughout the experiment, indicating the redeposition of corrosion products on the alloy’s surface. This phenomenon prevents further corrosion within the alloy volume, although it is temporary.

The evaluation of the alloy samples’ surface conditions reveals that pitting corrosion is observed on the surface of the alloy in the initial state (CG) after 6 h of the experiment (Figure 11a). In contrast, in the FG state alloy samples, this effect only appears on the 14th day of the experiment (Figure 11b,c). Nevertheless, in all states of the alloy, there is an uneven deposition of corrosion products on the surface, resulting in localized areas of increased Mg(OH)₂ concentration.

Further assessment of the corrosion properties through potentiodynamic curves confirms the findings from the immersion method. The alloy in the FG state demonstrates improved corrosion resistance due to the more pronounced passivation effect of magnesium decomposition products, notably Mg(OH)₂. Potentiodynamic curves are presented in Figure 12, and calculated values of corrosion current and polarization resistance are provided in

Table 3. The values of electrochemical characteristics for the Mg–Zn–Zr–Ce alloys.

State	Eocp, V	Ecorr, V	Icorr, µA/cm2	βa, V/dec	βk, V/dec	Rp, KOhm·cm2
Initial, CG	−1.56 ± 0.02	−1.56 ± 0.05	14.4 ± 2.6	0.17	0.26	3.1
3abc pressing, FG	−1.54 ± 0.02	−1.47 ± 0.04	4.3 ± 0.9	0.12	0.30	8.7
3abc pressing and rolling, FG	−1.53 ± 0.03	−1.45 ± 0.05	2.8 ± 0.4	0.11	0.11	8.1

.

Converting the alloy to the FG state through 3abc pressing reduces the corrosion current to 4.3 µA/cm², compared to the initial state of the alloy, which had a corrosion current of 14.4 µA/cm². Additionally, the polarization resistance increases from 3.1 to 8.7 kΩ·cm², indicating an enhancement in the corrosion resistance of the alloy. Further increasing the degree of deformation by rolling the alloy also reduces the corrosion current to 2.8 µA/cm², although the effect is less pronounced than that observed with 3abc pressing.

In the case of 3abc pressing, the significant change in corrosion parameters is attributed to an increase in the fraction of grain boundaries within the material’s volume.

A fine-grained microstructure often results in a more homogeneous distribution of alloying elements and secondary phases [47]. This can reduce the likelihood of galvanic corrosion, which occurs when there are potential differences between different phases or regions in the alloy. It is not only relevant for Mg-based alloys, but also for steels [48]. The refinement of grain structure typically results in improved mechanical properties, such as increased yield strength and hardness. This can enhance the alloy’s resistance to mechanical forms of corrosion, such as erosion and wear (we do not discuss this in this manuscript). A fine-grained microstructure can facilitate the formation of more uniform and protective oxide or hydroxide films on the surface of a magnesium alloy; we believe this is one of the predominant mechanisms that were shown by many authors for other types of Mg-based alloys [49]. These films act as barriers that prevent the ingress of corrosive species, thereby improving corrosion resistance. The grain boundaries can act as sites for the nucleation and growth of these protective films. Also, the grain boundary acts as a physical corrosion barrier. Small grain size creates more grain boundaries; as a consequence, the rate of corrosion in small-grained microstructures is slowed down compared with coarse-grained microstructures [50]. The properties of the hydroxide layer are determined by the combination of the chemical composition, microstructure, crystallographic texture, and grain size of metals, which decides the specific breakdown value (E_corr) of the hydroxide layer [51].

Microgalvanic corrosion can occur in magnesium alloys due to the presence of secondary phases or impurities that have different electrochemical potentials compared to the matrix. A fine-grained microstructure can reduce the size and distribution of these secondary phases or impurities, thereby minimizing the extent of microgalvanic corrosion. Due to the grain refinement, the process of pitting corrosion could be significantly reduced [52]. It could be said that corrosion resistance is a complex phenomenon, as described in [53], and grain refinement does not necessarily result in reduction in corrosion rate for Mg-based alloys. Hence, further studies are required to define the mechanisms of corrosion in the case of Mg–Zn–Zr–Ce alloys after SPD.

Importantly, the qualitative and quantitative composition of the alloy remains unaltered during the transition from the initial state (CG) to the FG state. Moreover, the refinement of the microstructure induces a shift in the corrosion potential to a less electronegative region.

4. Conclusions

In conclusion, the combined severe plastic deformation (SPD) method applied to a Mg–Zn–Zr–Ce alloy contributes significantly to effective grain refinement, resulting in the formation of a finely dispersed fine-grained (FG) structure with an average size of structural elements measuring 1.0 μm. This FG structure is primarily composed of α-grains based on a solid solution, hydrides of Ce₂H₅ and ZrH₂, and inclusions of intermetallic compounds of MgZnCe with various morphologies.

The FG structure obtained through the combined SPD method in the Mg–Zn–Zr–Ce alloy imparts the highest mechanical characteristics to the alloy. In tensile tests, it exhibits a yield strength (σ_YS) of 250 MPa, ultimate tensile strength (σ_UTS) of 270 MPa, and elongation (δ) of 3%. In compression tests, the FG samples exhibit a yield strength (σ_YS) of 240 MPa, ultimate compressive strength (σ_UTS) of 470 MPa, and compression deformation (δ) of 12%. Notably, the FG samples exhibit slightly higher deformation resistance under compression (δ = 12%) compared to the initial condition (IC) samples (δ = 10%).

Furthermore, the Mg–Zn–Zr–Ce alloy exhibits excellent corrosion properties. Additional processing of the alloy through the combined plastic deformation method further enhances its corrosion resistance by promoting the more intensive formation of a passivating layer of Mg(OH)₂ on the material’s surface. Electrochemical tests indicate a significant reduction in the corrosion current of the alloy subjected to combined SPD, decreasing from 14.4 to 2.8 µA/cm², signifying a notable decrease in the corrosion rate.

Consequently, this alloy holds great promise as a bioresorbable material for implants subjected to high mechanical loads.