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Article

Effect of Cross- or Unidirectional Rolling on the Microstructure, Corrosion Rate, and Hemolysis of Ternary Magnesium–Zinc–Gallium Alloys

by
Anabel Azucena Hernández-Cortés
1,2,
José C. Escobedo-Bocardo
1,
José Manuel Almanza-Robles
1 and
Dora Alicia Cortés-Hernández
1,*
1
Cinvestav-Unidad-Saltillo, Ave. Industria Metalúrgica 1062, Parque Industrial Saltillo-Ramos Arizpe, Ramos Arizpe C.P. 25900, Mexico
2
Materials Engineering Department, Tecnológico Nacional de México, Campus Saltillo, Blvd. Venustiano Carranza #2400, Col. Tecnológico, Saltillo C.P. 25280, Mexico
*
Author to whom correspondence should be addressed.
Metals 2025, 15(11), 1165; https://doi.org/10.3390/met15111165
Submission received: 10 September 2025 / Revised: 9 October 2025 / Accepted: 17 October 2025 / Published: 22 October 2025
Browse Figures
Versions Notes

Abstract

The effect of cross- or unidirectional rolling on the microstructure, corrosion rate, texture, and hemolysis of the Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga alloys was evaluated. After both rolling processes, the microstructure of the as-cast alloys was considerably refined due to the recrystallization process, obtaining higher grain refinement after cross-rolling. The Mg-1.5Zn-0.375Ga alloy showed a finer microstructure than the Mg-0.5Zn-0.25Mg alloy due to the effect of both the severe plastic deformation obtained after cross-rolling and the higher amount of alloying elements, which act as grain refiners. After unidirectional rolling, the texture intensity of the basal plane increases, while the cross-rolled alloys show lower texture intensity due to the activation of the pyramidal and/or prismatic slip systems. The cross-rolled alloys showed a higher corrosion rate than the unidirectionally rolled alloys due to the basal texture developed. The Mg-1.5Zn-0.375Ga alloy showed a higher corrosion rate than the Mg-0.5Zn-0.25Ga alloy since the voids formed during heat treating were not fully eliminated during rolling. The Mg-0.5Zn-0.25Ga alloy after unidirectional rolling was not hemolytic (4.7%) and showed the lowest corrosion rate (0.8 mm/y). Thus, this alloy may be an excellent candidate for use in the fabrication of biodegradable implants.

Graphical Abstract

1. Introduction

Biomaterials are used to help damaged tissues recover their function [1]. Metallic materials are mainly used in applications where the assisted tissue requires mechanical strength. These materials, in addition to meeting certain specific requirements depending on their application, must have corrosion resistance and biocompatibility [2] due to their contact with body fluids and their interaction with biological systems. In recent years, a field of metallic biomaterials, known as biodegradable metals, has become more important for biomedical applications. In particular, alloys based on zinc, magnesium, and iron are the most studied since these metals have inherent biodegradability, biocompatibility, and adequate strength for mechanically stressed implants. Nevertheless, magnesium alloys show promising applications for revolutionizing tissue engineering since they have certain inherent advantages, such as density and mechanical properties close to those of bone [3]. Additionally, it has been discovered that Mg2+, a by-product released during the corrosion of magnesium in an aqueous environment, benefits the activation of new bone tissue [4].
Recent studies on magnesium and its alloys have focused on optimizing mechanical properties [5], ensuring their non-toxicity [3], and mainly on regulating magnesium degradation. A balance between the corrosion rate and the kinetics of the healing process may be achieved [6]. When a magnesium implant is in contact with body fluid, it is oxidized and metal cations are released due to the anodic reaction, while H+ and OH− are released by cathodic reactions [7]. The hydroxide anions generate a pH increase in the medium surrounding the implant [8], while the released hydrogen may generate gas embolism [9]. Moreover, as the production of hydrogen gas bubbles may prevent appropriate osteocyte connectivity, it interferes with the initial healing process of cortical bone, resulting in the formation of calluses and cortical defects [10]. These reactions occur at the surface due to the formation of a galvanic coupling formed as a result of the potential difference between the metal matrix and intermetallic impurities in the magnesium matrix or grain boundaries.
The convection generated by the evolution of H2 from the surface to the medium, apart from the processes of migration and diffusion, controls the mass transfer between metallic surfaces and fluids. Simultaneously, proteins, lipids, and amino acids, among other organic molecules, are adsorbed onto the metallic surface, which influences the dissolution of magnesium. The high concentration of Cl− in physiological environments decomposes the magnesium hydroxide protective layer, leading to pitting corrosion [11]. The formation of a MgCl2 layer on the surface of magnesium decreases its corrosion resistance since this compound is highly soluble. In addition, the high concentration of hydroxyl ions in the medium increases alkalinity, and the saturation of calcium and phosphate ions leads to the deposition of a calcium phosphate layer on the undissolved Mg(OH)2 layer. These conditions provide magnesium-based implants with two of their main advantages over traditional implants: biodegradability and bioactivity [2].
The biodegradability of these materials implies that the degradation rate of the implant must not exceed the healing rate of the assisted tissue; hence, the corrosion rate must be carefully designed since degradation depends on the corrosion mechanism that dominates this process, and this, in turn, depends mainly on the alloy and its processing route [2]. Nonetheless, the intervention of bioactive species in the degradation process remains a topic of study since it has been observed that the alkalization of the immediate surrounding environment to the implant promotes the passivation of the magnesium surface under alkaline conditions due to certain phosphate species becoming less soluble [12]. These changes in alkalinity of the surrounding medium as a result of implant degradation may also interfere with some physiological functions and compromise tissues. It has been shown that changes in pH may compromise the erythrocyte membrane [13] and cause hemolysis in the vicinity of the implant. For this reason, various processing routes have been proposed with the aim of influencing microstructure and, indirectly, the degradation rate of the alloy [3,14].
The addition of alloying elements to magnesium may improve corrosion resistance and mechanical properties since these elements modify the structure and the surface morphology of the alloys. Gallium and zinc have been used as alloying elements to improve the corrosion resistance of the magnesium alloys. A Mg-4Ga-4Zn alloy was processed by equal-channel angular pressing, considerably improving mechanical properties and decreasing biocorrosion rates [15]. In one of our previous works, Mg-Ga binary alloys were developed; after heat treatment, the corrosion rate of the alloys with less than 1 wt.% of decreased Ga [16]. Wu et al. [17] used liquid Ga for coating a magnesium alloy, and a Ga-Mg layer was formed on the surface, improving corrosion resistance. Among the strategies used for decreasing the corrosion rate of magnesium alloys, the addition of Zn, combined with surface coating and other fabrication techniques, is recommended [18]. He et al. [19] used friction stir processing (FSP) to refine the microstructure of Mg-2Zn-xGa (x = 3, 5, 7 wt.%) alloys. With increasing Ga contents, the FSP-processed Mg-2Zn-xGa alloys show finer grains but a faster corrosion rate due to the increase in Mg5Ga2 precipitates and weaker (0001) basal texture. The corrosion layer presented a double-layer structure, with an outer layer dominated by Ca3(PO4)2 and an inner layer dominated by Mg(OH)2 [19]. Sahu and Yamamoto [20] reported that the mechanical properties and corrosion resistance of Mg-Zn alloys improve by increasing the quantity of Zn up to 4 wt.%. On the other hand, Zn-Mg alloys have also been developed for biomedical applications since Zn shows a moderate degradation rate [21]. Liu et al. performed studies with Mg-Zn alloys for applications in neurosurgery. In vitro and in vivo evaluations revealed that the heterostructured Mg-1.25Zn with ultra-fine grains and twins exhibits biocompatibility, immunomodulatory properties, and a dual capacity to enhance osseous regeneration and meningeal functional restoration [22].
Grain refinement is a technique that has been demonstrated to be effective for improving the mechanical properties and corrosion resistance of magnesium alloys [23,24,25]. When a magnesium-based alloy is rolled at high temperatures, its structure is modified (microstructure, macrostructure, texture, etc.) and grains are refined. This modification of the structure is different depending on whether the high-temperature rolling process is either crossed or unidirectional.
In our previous work [26], the influence of Ga and Zn (Ga from 0.375 to 1.5 wt% and Zn from 1.5 to 6 wt%) on the corrosion behavior of as-cast Mg–Ga–Zn alloys was studied. Microgalvanic corrosion was observed on the surface of the alloys due to the difference in potential between Mg and the second phases of (Mg, Ga)7Zn3 and (Mg, Zn)5Ga2. It was observed that the corrosion rate increases as the amount of second-phase materials is increased. However, the corrosion rate remains high in alloys with a higher amount of alloying elements. In addition, alloys with a lower amount of alloying elements may require additional treatments, such as heat treatment and high-temperature rolling, in order to modulate corrosion behavior. Taking into account that the addition of the alloying elements is not enough for modulating the corrosion rate of these Mg-Zn-Ga alloys, in this work, the effect of either unidirectional or cross-rolling on the microstructure, corrosion rate, bioactivity, texture, and hemolysis of the Mg-0.5Zn-0.25Ga and the Mg-1.5Zn-0.375Ga alloys is studied.

2. Materials and Methods

Ternary Mg-Zn-Ga alloys were prepared using the following high-purity metals: Mg (Stanford Advanced Materials, Lake Forest, CA, USA, 99.99 wt.%), Zn (Jalmek, San Nicolás de los Garza, México, 99.5 wt.%), and Ga (Sigma Aldrich, St. Louis, MO, USA, 99.99 wt.%). The alloys were obtained using a graphite crucible. The appropriate amounts of magnesium, zinc, and gallium were heated and then kept at 750 °C for 15 min under a controlled atmosphere (Ar-1%SF6). During this period of time, each melt was stirred for homogenization. Inductively coupled plasma atomic emission spectroscopy (ICP-OES Perkin Elmer model Optima 8300, Perkin Elmer, Waltham, MA, USA) was used to determine the chemical composition of pure magnesium and alloys (Table 1).
Specimens of 10 mm × 10 mm × 6 mm of pure magnesium and the alloys were machined and then heat-treated (350 °C for 12 h, T4 solution heat treatment). After this period of time, specimens were quenched in water at 25 °C. Then, specimens were rolled at high temperatures (400 °C) using two different routes: unidirectional or cross-directional rolling according to our previous work [26]. Three replicas were made for each sample. For unidirectional rolling, the rolling direction was alternated by rotating the specimens 180° on the normal rolling axis after each rolling pass. For cross-rolling, the rolling direction was alternated by rotating the specimen 90° on the normal rolling axis after each rolling pass. The heat-treated specimens were heated up to 400 °C and then rolled (International Rolling Mill model 4060, International Rolling Mill, Pawtucket, RI, USA). Ten rolling passes were performed using a roller speed of 0.06 m/s, reheating the specimens at 400 °C before each rolling pass. The final thickness was 0.3 cm; therefore, the total thickness reduction was 50% for all specimens [26]. The alloy texture was measured using an X-ray diffractometer (Bruker, Billerica, MA, USA, D8 Advance) equipped with a texture goniometer. Pole figures of the basal (0001), pyramidal (10 1 ¯ 1), and prismatic (10 1 ¯ 0) planes were obtained. Microstructural analyses, in vitro bioactivity assessments using simulated body fluid (SBF), and the evaluation of the corrosion rate using the weight loss method were performed according to the literature [27]. Instead of electrochemical testing, the immersion method was selected according to Bazhenov et al. [15], who developed a similar ternary alloy (Mg-4Zn-4Ga). They found that the corrosion rate obtained via the electrochemical test was 6–10 times higher than that obtained by the immersion test. The long-time immersion corrosion test was more precise than rapid electrochemical corrosion measurements. Hemolysis was determined according to the ASTM F756-17 standard [28]. In this test, standards were prepared using the hemoglobin standard (Sigma-Aldrich, St. Louis, MO, USA) and Drabkin’s reagent (cyanmethemoglobin [CMH], Sigma-Aldrich, St. Louis, MO, USA). Blood was donated by three healthy nonsmoking adults, and it was collected in Vacutainer test tubes with heparin to prevent coagulation. Three replicas were made for each sample. Quadrangular prism-shaped specimens (0.5 cm × 1 cm × 0.3 cm) were obtained. Specimens were ground using SiC papers (500–2400 grit) and then cleaned with ethanol in an ultrasonic bath for 20 min. The hemolysis percentage was calculated using Equation (1):
H e m o l y s i s   % = S − B T 8 − B   ×   100
where S is the hemoglobin concentration of sample tube supernatant, B is the hemoglobin concentration without a sample (blank), and T is the diluted blood hemoglobin concentration. In addition, corroded specimens were analyzed by scanning electron microscopy (SEM), energy-dispersive spectroscopy (EDS), X-ray diffraction (XRD), and Fourier transform infrared spectroscopy (FTIR).
The appropriate hypothesis tests (tests for difference of means) to support the statistical significance of the differences found for the response variables (grain size, strength, elongation, corrosion rate, and hemolysis percentage) were performed using Minitab 19 statistical software. All hypothesis tests were performed at a significance level of 0.05 (α = 0.05).

3. Results and Discussion

3.1. Microstructure

The matrix of the alloys is α-magnesium, and the formed intermetallic compounds are located at the interdendritic regions and grain boundaries. During the T4 solution heat treatment, large intermetallics located at the interdendritic regions and grain boundaries are dissolved into the matrix. During pre-rolling heating (aging) and the rolling process, the intermetallics precipitate again and are homogeneously distributed throughout the alloy, presenting smaller sizes. Figure 1 shows SEM images of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga before and after heat treatment. As observed, after heat treatment, intermetallics (white particles, Figure 1a) were dissolved, leaving some small voids (black pores, Figure 1b). A larger amount of intermetallics is observed in the Mg-1.5Zn-0.375Ga alloy (Figure 1c). The selected area at higher magnifications shows an intermetallic in the as-cast alloy. After heat treatment (Figure 1d), almost all intermetallics were dissolved.
The as-cast ternary alloys showed both equiaxed and columnar grains (Figure 2a,d). The microstructure of the alloys after rolling was considerably refined, as it is shown in Figure 2. The transformation of large dendritic grains into refined equiaxed grains is observed. Table 2 shows the grain size of the alloys at different processing stages.
A finer microstructure is obtained after cross-rolling, which can be associated with the deformation mode applied during this type of rolling. During unidirectional rolling, the basal slip system is mainly activated. Cross-rolling induces severe plastic deformation, which causes the activation of pyramidal and/or prismatic slip systems, generating higher heterogeneity in the microstructure as a result of a rearrangement of dislocations, deformed grains, subgrains, and high-energy grain boundaries [29]. The recrystallization process that occurs during hot rolling, caused by the microstructure-induced deformation and the atomic mobility promoted by temperature, allowed the refinement and transformation of the microstructure to equiaxed grains [30,31].
Both zinc [23,32] and gallium [15] are grain refiners for the magnesium microstructure. This occurs as a result of the incorporation of the atoms of these elements into the magnesium network, which causes a network distortion due to the contraction of the c parameter [31]. The high degree of heterogeneity of the created areas promotes the dynamic recrystallization process. This may be due to the fact that these areas act as nucleation sites for the formation of recrystallized grains, resulting in a finer microstructure [30]. The higher content of Zn and Ga in the Mg-1.5Zn-0.375Ga alloy compared to that of the Mg-0.5Zn-0.25Ga alloy causes higher grain refinement (Table 2).
The microstructure of the cross-rolled ternary alloys exhibited a high density of twinning and double twinning (Figure 3) as a result of the activation of the twinning slip to satisfy the Von Mises criterion for deformation. Additionally, when the materials show a microstructure of large grains, as observed in these alloys, a subdivision of these grains may occur during rolling, giving rise to deformation bands. The presence of these deformation bands, grain boundary migration, and twinning also contributes to grain nucleation and growth during the hot cross-rolling process [33].

3.2. Texture

The texture of Mg-0.5Zn-0.25Ga after different processing stages is shown in Figure 4. Grain texture with a basal orientation is observed for the as-cast alloy. After unidirectional rolling, a strong basal texture is obtained, as indicated by the high value of its texture intensity. Basal slip in the grains is first activated since its critical resolved shear stress is the lowest. This texture is mainly formed due to the high activity of both basal slip and extension twinning. It is important to note that the strong basal texture leads to anisotropic properties in the alloy [34].
The cross-rolled alloy presents a lower texture intensity than that of the unidirectionally rolled alloy. During hot deformations, a change in the deformation type occurs, increasing the activity of the pyramidal-2 slip, resulting in a decrease in texture intensity [35]. The anisotropy of the cross-rolled alloy is lower than that of the unidirectionally rolled alloy.
The behavior of the Mg-1.5Zn-0.375Ga texture after the different processing stages is similar to that of the Mg-0.5Zn-0.25Ga alloy.

3.3. Corrosion

Figure 5 shows the corrosion rate of pure magnesium and that corresponding to the ternary alloys after different processing stages and after 28 days of immersion in SBF. The as-cast alloys show a lower corrosion rate than that of the as-cast pure magnesium, which is attributed to the contribution of the alloying elements. The respective oxides resulting from the reaction of gallium and zinc with SBF strengthen the Mg(OH)2 reaction layer and decrease the contact of unreacted magnesium with SBF. After heat treatment, no significant change in corrosion rate was observed for Mg-0.5Zn-0.25Ga. However, the corrosion of Mg-1.5Zn-0.375Ga increased significantly. Occasionally, when the quantity of the alloying elements is high and, therefore, the intermetallics are large, voids appear after heat treatment, and they may promote corrosion. These voids may be eliminated during rolling operations so that the corrosion rate decreases considerably after unidirectional and cross-rolling. The cross-rolled ternary alloys show a higher corrosion rate than that of the unidirectionally rolled alloys. This behavior has previously been documented [36,37] for magnesium-based alloys, and it is related to the basal texture developed by unidirectional rolling. When the alloy is unidirectionally rolled, grains are parallelly oriented to the rolling surface (Figure 4), and since this orientation is more compact, atomic detachment is made more difficult since the energy needed to break their bonds is higher. Thus, corrosion resistance is increased [38]. The corrosion rate of the heat-treated Mg-1.5Zn-0.375Ga alloy decreased dramatically after rolling; however, it is higher than that of Mg-0.5Zn-0.25Ga since voids were not completely eliminated during rolling. The corrosion rate values corresponding to the Mg-0.5Zn-0.25Ga are considerably lower than those corresponding to Mg-1.5Zn-0.375Ga and pure Mg. The corrosion rate for the unidirectionally rolled Mg-0.5Zn-0.25Ga alloy is 0.8 mm/year, which is similar to that obtained by He et al. [19], who tested a friction-stir-processed Mg-2Zn-xGa alloy, obtaining values within the range of 0.8–0.95 mm/year. On the other hand, Bashenov et al. [15] tested an ECAP Mg-4Ga-4Zn alloy, obtaining values between 0.16 and 0.22 mm/year.
According to the EDS results and considering the Mg-Zn and Mg-Ga binary-phase diagrams, two different intermetallic compounds were detected: (Mg, Ga)7Zn3 and (Mg, Zn)5Ga2. Figure 6 shows the corroded surface of the as-cast Mg-0.5Zn-0.25Ga alloy after 28 days of immersion in SBF. As observed, corrosion occurs in the α-Mg matrix (Xb) that surrounds the Mg-Zn-Ga intermetallic particle (Xa). The EDS spectra indicate that, in this case, the intermetallic particle corresponds to (Mg, Ga)7Zn3, which is embedded into the α-Mg matrix. According to the SEM image, micro-galvanic corrosion has occurred. This type of corrosion is due to the difference in potential between Mg and the intermetallic compound, since Mg has a lower standard reduction potential (–2.38 V) than those corresponding to Zn and Ga.
The corroded surfaces of the Mg-0.5Zn-0.25Ga and the Mg-1.5Zn-0.375Ga alloys after different processing stages are shown in Figure 7 and Figure 8. These SEM images are in agreement with the corresponding calculated corrosion rates (Figure 5). In the as-cast and rolled alloys, surfaces show a uniform type of corrosion with no pitting, while the heat-treated alloy shows more severe corrosion and pitting. Corrosion is more severe when the intermetallics are larger and located in specific areas, causing pitting corrosion. After heat treatment and high-temperature rolling, intermetallics were dissolved and reprecipitated as small, homogeneously distributed precipitates, which promote more uniform corrosion. It is observed that the most uniform corrosion occurs in the unidirectional rolled alloys (Figure 8a,b).

3.4. In Vitro Bioactivity Assessment

The XRD patterns corresponding to pure magnesium and alloys after 28 days of immersion in SBF are shown in Figure 9. The compounds identified were magnesium oxide (MgO) and magnesium hydroxide (Mg(OH)2), which are products of the reaction between magnesium and SBF. Magnesium from the substrate was also detected. It was not possible, using this technique, to detect the presence of compounds formed between Mg, Ga, and Zn (intermetallics); it was also not possible to identify Ca- and P-rich compounds due to the detection sensitivity of the technique used for the analysis.
Figure 10 and Figure 11 show SEM images of the rolled alloys after 7 days of immersion in SBF. As observed in the SEM images of the Mg-0.5Zn-0.25Ga alloy (Figure 10), the layer of corrosion products (mainly Mg(OH)2) is covered by a second layer of Ca- and P-rich compounds which have two different morphologies: (i) the typical petal-like morphology [39] that resembles that of the apatite formed in the existing bioactive systems (Ca/P ratio of 1.30 to 1.67 [40] and) (ii) irregular white precipitates (Ca/P ratio of 0.62). These Ca- and P-rich compounds may be calcium phosphates or magnesium/calcium phosphates, which are formed due to the reaction of Ca2+, PO3−, Mg2+, OH−, Zn2+, Ga3+, and HPO42−, among other ions present in SBF. On the other hand, the SEM images of the Mg-1.5Zn-0.375Ga alloy (Figure 11) show a layer of a Ca- and P-rich compound with spheroidal morphology.
The presence of magnesium ions resulting from substrate dissolution causes the precipitation of different species. Magnesium ions affect the characteristics of the formed calcium phosphate species in two ways: by inhibiting the development of the species’ crystallinity and by forming magnesium-containing compounds such as Mg3(PO4)2·22H2O [40]. As a result of this modification, the stoichiometry and morphology of the species are affected. A change in the morphology of the phosphate species (brushite) was observed with the substitution of calcium ions by magnesium ions [41]; the species presented a spheroidal morphology after such substitutions occurred, while the unsubstituted species presented an acicular morphology.
Figure 12 shows the pH behavior of the SBF as a function of immersion time for pure magnesium, Mg-0.5Zn-0.25Ga, and Mg-1.5Zn-0.375Ga after different processing stages. As observed, pH increases rapidly during the first seven days due to the cations (Ca, Mg, and Zn). After seven days, pH values remain almost constant. The formation and stability of the apatite species formed on the reaction layer depend on pH and the ions available in the SBF [39,41]. The size and quantity of the species formed on the reaction layer may decrease due to the presence of zinc in the alloy. This phenomenon can be related to two factors. Firstly, in aqueous solutions with pH values between 7 and 10 (Figure 12), Zn2+ competes with Mg2+ to bind to OH− and form Mg(OH)2 and Zn(OH)2, resulting in a decrease in pH, thereby reducing the alkalinity of the medium and increasing the solubility of the apatite species formed so that their quantity on the reaction surface is reduced. The second factor is the rapid dissolution of the alloy as a result of the increase in the quantity of second phases that promote galvanic corrosion, which causes the reaction layer to crack and weaken, so the capacity to retain Mg(OH)2 and (Ca1-XMgX)10(PO4)6(OH)2 species on the surface decreases.
In rolled alloys, the amount of calcium phosphate species formed or retained on the substrate is higher because the alloys present a lower corrosion rate and, therefore, a lower dissolution/detachment of the substrate on which these species precipitate. Alloys with unidirectional rolling are those that present the highest retention of calcium phosphate species on the substrate. This is attributed to the fact that they present a lower corrosion rate due to the basal texture developed during this type of rolling. In this case, the grains of the alloys are oriented parallel to the rolling surface, as determined in the texture evaluation presented in the previous section. This basal orientation, being more compact, makes atomic detachment difficult since the energy required to break their bonds is higher [37], which increases corrosion resistance. Additionally, the reaction products between the magnesium hydroxide layer formed on the alloys and the free calcium and phosphorus ions in the SBF react to form these products, which eventually contribute to a decrease in the corrosion of the substrate [42].
The FTIR spectra corresponding to the ternary rolled alloys after 7 days of immersion in SBF are shown in Figure 13. For comparison purposes, the spectrum corresponding to hydroxyapatite is also included. For both the alloys and the synthetic hydroxyapatite, the absorption bands for the PO43− group were observed at 560, 600, 630, and 1020 cm−1, confirming the formation of calcium/magnesium phosphates on the alloy after immersion in SBF. The broadening of the absorption bands observed indicates the partial substitution of Ca ions by Mg ions into the phosphate crystalline structure, which leads to a decrease in crystallinity [43]. An absorption band at 1452 cm−1 was also observed, which corresponds to the carbonate group (CO32−) found in calcium orthophosphates, calcium diphosphates, and even amorphous calcium phosphates [44]. The presence of this group in the compounds formed is related to a change in morphology, since the substitution of CO32− by PO43− may have occurred and also causes a low resolution of the absorption bands corresponding to the phosphates in the FTIR spectra [44].

3.5. Hemolysis

The hemolytic behavior of the ternary alloys is presented in Figure 14. The ASTM F756-17 standard indicates that materials exceeding a hemolysis value of 5% are considered hemolytic, which implies that their interaction with blood cells may damage the erythrocyte membrane. According to the observations in the figure, the as-cast magnesium and the as-cast ternary alloys are hemolytic. After the processing of pure magnesium (T4 solution heat treatment and rolling), hemolysis increases slightly. On the other hand, after the processing of the alloys, hemolysis decreases markedly. The hemolysis percentage achieved by the Mg-0.5Zn-0.25Ga alloy after unidirectional rolling was 4.7%. This result indicates that this alloy is not hemolytic, while after cross-rolling, this alloy shows a hemolysis value of 6.0%. The decrease in hemolysis obtained after rolling is related to the texture developed by the alloy after this processing stage. Furthermore, in the case of the unidirectionally rolled alloy, the developed basal texture led to a smoother and more compact surface, decreasing the lysis of erythrocytes. Additionally, the decrease in erythrocyte lysis is related to a decrease in the corrosion rate of the rolled specimens, as it reduces the release of ions into the medium [14]. Several authors [14,45] have found a direct relationship between the increase in pH in the immersion medium, generated by the rapid degradation of magnesium, and an increase in the hemolysis caused by the material. Zhang et al. [46] reported that when the pH of the medium is lower than a critical value of 10.3, the hemolysis rate is lower than the recommended value of 5%. Therefore, by improving the corrosion resistance of ternary alloys through alloying with gallium and zinc and applying subsequent rolling processes, the hemolysis percentages were reduced. The other factor contributing to a reduction in hemolysis percentages in rolled alloys is attributed to the retention of apatite species on the surface of the alloys as a result of a lower corrosion rate. Ye et al. [47] demonstrated that the presence of a calcium phosphate film considerably reduces hemolysis since this film keeps pH values stable, and therefore, the damage to erythrocytes is reduced.
Table 3 summarizes the values of grain size, corrosion rate, and hemolysis percentage of the Mg-Zn-Ga rolled alloys. It can be clearly observed that, for both alloys, cross-rolling causes greater grain refinement, and unidirectional rolling causes a lower corrosion rate and a lower percentage of hemolysis.

4. Conclusions

  • The feasibility of obtaining rolled biodegradable magnesium alloys (Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga) for biomedical applications has been demonstrated.
  • After heat treating, the size of the (Mg, Ga)7Zn3 and (Mg, Zn)5Ga2 intermetallics decreased and were homogeneously distributed. Both ternary alloys were either unidirectionally or cross-rolled.
  • After rolling, the alloys’ microstructure was considerably refined. A higher refinement was achieved after cross-rolling. This is due to the fact that cross-rolling induces severe plastic deformations, which cause the activation of pyramidal and/or prismatic slip systems, generating higher heterogeneity in the microstructure.
  • The Mg-1.5Zn-0.375Ga alloy showed a finer microstructure than the Mg-0.5Zn-0.25Ga alloy due to both the severe plastic deformation that resulted after cross-rolling and the higher amount of alloying elements, which act as grain refiners.
  • Texture intensity of the basal plane increases after unidirectional rolling, while lower texture intensity is observed for the cross-rolled alloy due to the activation of additional slip systems, such as the pyramidal and/or the prismatic systems.
  • The lower corrosion rate was observed for the unidirectionally rolled alloys due to the developed basal texture. The Mg-1.5Zn-0.375Ga alloy showed a higher corrosion rate than the Mg-0.5Zn-0.25Ga alloy since the voids formed during heat treating were not fully eliminated during rolling.
  • Both rolled alloys have been shown to be bioactive due to the formation of a Ca- and P-rich layer with a similar morphology to that of the apatite formed on the existing bioactive systems after immersion in SBF.
  • The unidirectionally rolled Mg-0.5Zn-0.25Ga alloy demonstrated no hemolytic properties (4.7%).
  • According to the results obtained in this work, this unidirectionally rolled alloy may be a potential biodegradable material for bone replacement and regeneration.

Author Contributions

Conceptualization, A.A.H.-C., J.C.E.-B. and D.A.C.-H.; methodology, A.A.H.-C.; formal analysis, A.A.H.-C., J.C.E.-B., J.M.A.-R. and D.A.C.-H.; investigation, A.A.H.-C. and J.C.E.-B.; data curation, D.A.C.-H.; writing—original draft preparation, A.A.H.-C., J.C.E.-B., J.M.A.-R. and D.A.C.-H.; writing—review and editing, A.A.H.-C., J.C.E.-B., J.M.A.-R. and D.A.C.-H.; supervision, J.C.E.-B. and D.A.C.-H. All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Acknowledgments

The authors acknowledge the doctoral scholarship for A.A. Hernández-Cortés provided by CONAHCYT.

Conflicts of Interest

The authors declare no conflicts of interest.

References

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Figure 1. SEM images of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga: as-cast ((a,c) white intermetallics indicated by white arrows in (a)) after T4 solution heat treatment ((b,d) black pores).
Figure 1. SEM images of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga: as-cast ((a,c) white intermetallics indicated by white arrows in (a)) after T4 solution heat treatment ((b,d) black pores).
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Figure 2. Micrographs of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga: as-cast (a,b), unidirectional rolled (c,d), and cross-rolled (e,f).
Figure 2. Micrographs of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga: as-cast (a,b), unidirectional rolled (c,d), and cross-rolled (e,f).
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Metals 15 01165 g003
Figure 3. Micrographs of the cross-rolled Mg-0.5Zn-0.25Ga (a) and Mg-1.5Zn-0.375Ga (b) alloys showing a high density of twinning and double twinning (white arrows).
Figure 3. Micrographs of the cross-rolled Mg-0.5Zn-0.25Ga (a) and Mg-1.5Zn-0.375Ga (b) alloys showing a high density of twinning and double twinning (white arrows).
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Figure 4. Pole figures of the Mg-0.5Zn-0.25Ga alloy after different processing steps (as-cast, unidirectional, and cross-rolled): (a) basal plane (0001), (b) pyramidal plane (10 1 ¯ 1), and (c) prismatic plane (10 1 ¯ 0). RD: Rolling direction; TD: transverse direction; ND: normal direction.
Figure 4. Pole figures of the Mg-0.5Zn-0.25Ga alloy after different processing steps (as-cast, unidirectional, and cross-rolled): (a) basal plane (0001), (b) pyramidal plane (10 1 ¯ 1), and (c) prismatic plane (10 1 ¯ 0). RD: Rolling direction; TD: transverse direction; ND: normal direction.
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Figure 5. Corrosion rate of pure magnesium and ternary alloys after different processing stages and after 28 days of immersion in SBF.
Figure 5. Corrosion rate of pure magnesium and ternary alloys after different processing stages and after 28 days of immersion in SBF.
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Figure 6. Corroded surface of the as-cast Mg-0.5Zn-0.25Ga ternary alloy after 28 days of immersion in SBF and EDS spectra of the matrix (Xb) and intermetallics (Xa).
Figure 6. Corroded surface of the as-cast Mg-0.5Zn-0.25Ga ternary alloy after 28 days of immersion in SBF and EDS spectra of the matrix (Xb) and intermetallics (Xa).
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Figure 7. Corroded surfaces of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga after 28 days of immersion in SBF; as-cast (a,b) and heat-treated alloys (c,d).
Figure 7. Corroded surfaces of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga after 28 days of immersion in SBF; as-cast (a,b) and heat-treated alloys (c,d).
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Figure 8. Corroded surfaces of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga after 28 days of immersion in SBF; unidirectional rolled (a,b) and cross-rolled alloys (c,d).
Figure 8. Corroded surfaces of Mg-0.5Zn-0.25Ga and Mg-1.5Zn-0.375Ga after 28 days of immersion in SBF; unidirectional rolled (a,b) and cross-rolled alloys (c,d).
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Figure 9. XRD patterns of pure magnesium and magnesium alloys after 28 days of immersion in SBF.
Figure 9. XRD patterns of pure magnesium and magnesium alloys after 28 days of immersion in SBF.
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Figure 10. SEM images of the Mg-0.5Zn-0.3Ga rolled alloy after 7 days of immersion in SBF: (a,b), figure (b) is a close up of the box marked in figure (a). The EDS spectrum of the irregular white precipitates is shown below figures (a,b).
Figure 10. SEM images of the Mg-0.5Zn-0.3Ga rolled alloy after 7 days of immersion in SBF: (a,b), figure (b) is a close up of the box marked in figure (a). The EDS spectrum of the irregular white precipitates is shown below figures (a,b).
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Figure 11. SEM images of the Mg-1.5Zn-0.375Ga rolled alloy after 7 days of immersion in SBF (a,b), figure (b) is a close up of the box marked in figure (a). The and EDS spectrum of the irregular white precipitates is shown below figures (a,b).
Figure 11. SEM images of the Mg-1.5Zn-0.375Ga rolled alloy after 7 days of immersion in SBF (a,b), figure (b) is a close up of the box marked in figure (a). The and EDS spectrum of the irregular white precipitates is shown below figures (a,b).
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Metals 15 01165 g012
Figure 12. pH behavior of the SBF as a function of immersion time for pure magnesium, Mg-0.5Zn-0.25Ga, and Mg-1.5Zn-0.375Ga after different processing stages.
Figure 12. pH behavior of the SBF as a function of immersion time for pure magnesium, Mg-0.5Zn-0.25Ga, and Mg-1.5Zn-0.375Ga after different processing stages.
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Metals 15 01165 g013
Figure 13. FTIR spectra of the synthetic HA and the ternary rolled alloys after 7 days of immersion in SBF.
Figure 13. FTIR spectra of the synthetic HA and the ternary rolled alloys after 7 days of immersion in SBF.
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Metals 15 01165 g014
Figure 14. Hemolysis percentage of pure magnesium, Mg-0.5Zn-0.25Ga, and Mg-1.5Zn-0.375Ga after different processing stages.
Figure 14. Hemolysis percentage of pure magnesium, Mg-0.5Zn-0.25Ga, and Mg-1.5Zn-0.375Ga after different processing stages.
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Table 1. Chemical composition (wt.%) of pure Mg and Mg-Zn-Ga alloys.
Table 1. Chemical composition (wt.%) of pure Mg and Mg-Zn-Ga alloys.
Nominal CompositionMgGaZnFeNiCu
pure Mg99.9900≤0.0002≤0.0006≤0.0002
Mg-0.5Zn-0.25Ga99.210.250.5220.0030.002≤0.0002
Mg-1.5Zn-0.375Ga98.380.351.350.0020.002≤0.0002
Table 2. Average grain size of pure magnesium and magnesium alloys at different processing stages.
Table 2. Average grain size of pure magnesium and magnesium alloys at different processing stages.
Nominal CompositionAverage Grain Size (mm)
As-CastUnidirectional RolledCross-Rolled
Pure Mg530 (±196.48)--
Mg-0.5Zn-0.25Ga534 (±181.5)90 (±22.45)42 (±27.93)
Mg-1.5Zn-0.375Ga477 (±129.14)64 (±17.22)27 (±11.3)
Table 3. Grain size, corrosion rate, and hemolysis percentage of the Mg-Zn-Ga rolled alloys (UR: unidirectional rolling; CR: cross-rolling).
Table 3. Grain size, corrosion rate, and hemolysis percentage of the Mg-Zn-Ga rolled alloys (UR: unidirectional rolling; CR: cross-rolling).
AlloyGrain Size
(μm)		Corrosion Rate
(mm/Year)		Hemolysis
(%)
URCRURCRURCR
Mg-0.5Zn-0.25Ga90420.80.954.75.6
Mg-1.5Zn-0.375Ga64271.41.856.014.2
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Hernández-Cortés, A.A.; Escobedo-Bocardo, J.C.; Almanza-Robles, J.M.; Cortés-Hernández, D.A. Effect of Cross- or Unidirectional Rolling on the Microstructure, Corrosion Rate, and Hemolysis of Ternary Magnesium–Zinc–Gallium Alloys. Metals 2025, 15, 1165. https://doi.org/10.3390/met15111165

AMA Style

Hernández-Cortés AA, Escobedo-Bocardo JC, Almanza-Robles JM, Cortés-Hernández DA. Effect of Cross- or Unidirectional Rolling on the Microstructure, Corrosion Rate, and Hemolysis of Ternary Magnesium–Zinc–Gallium Alloys. Metals. 2025; 15(11):1165. https://doi.org/10.3390/met15111165

Chicago/Turabian Style

Hernández-Cortés, Anabel Azucena, José C. Escobedo-Bocardo, José Manuel Almanza-Robles, and Dora Alicia Cortés-Hernández. 2025. "Effect of Cross- or Unidirectional Rolling on the Microstructure, Corrosion Rate, and Hemolysis of Ternary Magnesium–Zinc–Gallium Alloys" Metals 15, no. 11: 1165. https://doi.org/10.3390/met15111165

APA Style

Hernández-Cortés, A. A., Escobedo-Bocardo, J. C., Almanza-Robles, J. M., & Cortés-Hernández, D. A. (2025). Effect of Cross- or Unidirectional Rolling on the Microstructure, Corrosion Rate, and Hemolysis of Ternary Magnesium–Zinc–Gallium Alloys. Metals, 15(11), 1165. https://doi.org/10.3390/met15111165

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