Effect of High-Pressure Torsion on Microstructure, Mechanical and Operational Properties of Zn-1%Mg-0.1%Ca Alloy

Abstract

A study of the effect of high-pressure torsion (HPT) on the structure, phase composition, corrosion resistance, mechanical properties and bioactivity in vitro of Zn-1%Mg-0.1%Ca alloy was performed. It was shown that HPT leads to refinement of the alloy microstructure with the formation of recrystallized α-Zn grains with an average size of 750 ± 30 nm, and grains of a mixture of different phases with a size of 38 ± 7 nm. In addition, precipitation of Ca-enriched particles ~20 nm in size was observed. X-ray phase analysis showed that the Zn-1%Mg-0.1%Ca alloy consists of five phases (Zn, Mg₂Zn₁₁, MgZn₂, CaZn₁₁ and CaZn₁₃), whose volume fraction does not change after HPT. It was found that HPT does not lead to a deterioration in the corrosion resistance of the alloy. At the same time, HPT leads to an increase in the yield stress of the alloy from 135 ± 13 to 356 ± 15 MPa, the ultimate tensile strength from 154 ± 5 to 416 ± 31 MPa, and the ductility from 0.4 ± 0.1 to 5.5 ± 2.8%. No significant increase in hemolytic activity, bactericidal activity, and the ability to colonize the surface of the alloy by cells was revealed during the conducted studies. Additionally, there was no significant difference in these parameters in comparison with the control. However, HPT contributes to a decrease in the cytotoxicity of the alloy by an average of 10% compared to the annealed alloy. The conducted studies allow us to conclude that the Zn-1%Mg-0.1%Ca alloy is promising material for the development of biodegradable orthopedic medical implants.

1. Introduction

Bone injuries, including comminuted and impression fractures, are one of the most common types of injuries. In the vast majority of cases, such injuries lead to a decrease in the bone strength and a violation of its supporting function. In order not to reduce the quality of life of patients with such injuries, temporary reconstruction of bone structures using screws, plates, staples, etc., is often required. At present, bioinert titanium and titanium-based alloys, as well as corrosion-resistant steel, are most often used for these purposes [1,2,3]. The abovementioned materials have high strength characteristics and can successfully perform the function of a temporary bone frame. However, the major drawback is that their strength significantly exceeds the strength of the cortical bone, which can be damaged by the implant due to the so-called “stress-shielding effects” [4]. This significantly increases the risk of re-fracture of the bone in the area of implant attachment. In addition, bioinert implants must be removed after the healing of the fracture, which leads to a re-traumatization of the bone and surrounding tissues and an increase in the patient’s rehabilitation time. Therefore, for the last couple of decades, the interest of researchers and physicians has been directed to the development of new generation materials for osteosynthesis—biodegradable implants [5,6]. Such implants can gradually dissolve in the patient’s body, while transferring the load from the implant to the bone. In addition, due to the gradual degradation of the implants, there is no need for a second operation, which contributes to the speedy recovery of the patient. For a long time, the main candidates for the role of materials for the creation of metal biodegradable implants were magnesium and its alloys [7,8]. They have a high level of biocompatibility and strength characteristics similar to bone tissue properties. However, the disadvantage of magnesium alloys is the rather high rate of their degradation, which overtakes the rate of bone healing. In addition, the degradation of magnesium alloys is accompanied by the release of hydrogen, which can lead to the formation of gas bubbles in tissues and impair cell adhesion to the surface of the implant. Hänzi et al. [9] showed that implantation of Mg–Y–Zn alloy samples into the liver and lesser omentum of mini-pigs resulted in limited gas formation, while implantation into the rectus abdominis muscle resulted in a significant accumulation of hydrogen in the area of implantation. In [10], active outgassing with the formation of gas pockets was also observed during the degradation of Mg-6%Ag and Mg-10%Gd alloys in vivo. Therefore, despite the advantages, magnesium alloys may not always be applicable in osteosynthesis. Unlike rapidly decomposing magnesium alloys, zinc and its alloys show a more suitable degradation rate without the release of excess hydrogen [11,12,13]. Currently, much attention is paid to studying the effect of the composition of zinc alloys on their strength and corrosion properties, as well as biocompatibility. It has previously been shown that the implantation of pure zinc pins into the rectum of mice does not cause serious side effects. Pure zinc has also been found to inhibit acute inflammation through increased expression of ENA-78 (peptide) and F4/80 (glycoprotein) [14]. A study of the biocompatibility in vivo of Zn-0.05%Mg-x%Ag alloys (implantation into the drilled cavity of the bilateral distal rabbit femur) also showed no obvious toxicity during the 24-week implantation period [15]. It is also interesting that pure Zn and Zn-0.05%Mg alloy showed strong antibacterial activity against Escherichia coli and Staphylococcus aureus [16]. However, it was shown in [17] that Zn accumulated in tissues adjacent to the bone after implantation in vivo of the Zn–Mg–Fe alloy (98 wt.% ≤ Zn ≤ 99.5 wt.%, 0.01 wt.% ≤ Mg ≤ 0.5 wt.%, and 0.01 wt.% ≤ Fe ≤ 0.5 wt.%), even having good biocompatibility. Despite the above results, the accumulated statistics on the study of the biocompatibility of zinc-based alloys is still rather scarce and contradictory. Therefore, the study of the effect of zinc and its alloys on various aspects of biocompatibility is currently an urgent task.

Despite the advantages of zinc as a medical material, its main disadvantage is its low strength. In the as-cast state, the ultimate tensile strength of pure zinc is about 50 MPa [18]. Therefore, a vital task is the improvement of the mechanical characteristics of pure zinc, for example, by alloying. For example, the Zn–Mg system has become widespread, since magnesium not only increases the strength of pure zinc due to the formation of Zn₂Mg and Zn₁₁Mg₂ phases, but is also non-toxic. Thus, it was previously shown that an increase in the Mg content from 0.1 to 1.5 wt.% increases the ultimate tensile strength (UTS) from 104 MPa to 265 MPa, respectively [19]. However, a further increase in magnesium leads to a strong embrittlement of alloys due to an increase in the fraction of brittle eutectic phase [20]. Therefore, it is necessary to combine the alloying with deformation processing. It was shown in [21] that hot rolling leads to an increase in the strength and ductility of pure zinc up to 118 MPa and ε ~ 26.8%, respectively. Similarly, there is an increase in the strength and ductility of the Zn-0.4 wt.% Mg (UTS ~230 MPa and El ~23%) and Zn-0.8 wt.% Mg (UTS ~ 268 MPa and ε ~ 7.2%) [21]. It has also been shown that high-pressure torsion (HPT) leads to significant strengthening of pure Zn [22], Zn-0.5%Cu [23] and Zn-0.8%Ag [24] alloys. It was shown in [25] that grain refinement and the formation of a favorable texture in the Zn–1%Cu alloy after equal-channel angular pressing (ECAP) lead not only to an increase in its strength from 110 MPa to 231 MPa, but also to an increase in ductility from 7.3% to 94.2%. At the same time, it should not be forgotten that it is necessary to investigate not only the effect of alloying and deformation processing on the mechanical characteristics of alloys, but also on their main operational property—corrosion resistance. It was shown in [26] that hot extrusion of the Zn-1.2%Mg alloy leads to a small increase in degradation rate from 0.12 to 0.19 mm/y. At the same time, hot rolling of the Zn–1Cu–0.1Ti alloy leads to an increase in its degradation rate from 0.315 to 1.628 mm/g [27]. However, in the case of the Zn–1%Al alloy after hot extrusion, a decrease in the degradation rate from 0.166 to 0.145 mm/g was observed [28]. Therefore, the study of the effect of various types of deformation treatment not only on the strength and ductility, but also on the corrosion resistance and biocompatibility of zinc alloys remains an important task.

Considering the above information, the purpose of this research was to study the microstructure, phase composition, mechanical properties, corrosion resistance and bioactivity in vitro of Zn-1%Mg-0.1%Ca alloys after high-pressure torsion (HPT). The use of HPT makes it possible to significantly increase the strength of zinc alloys due to the strong refinement of their microstructure. It is the next step towards the development of the final medical product. Moreover, it is expected that the doping of Ca will make it possible to achieve the formation of a more uniform microstructure, which in turn will affect the corrosion resistance of the alloy [29].

2. Materials and Methods

The research material was a potential medical alloy with a nominal composition of Zn-1 wt.%Mg-0.1 wt.%Ca in this work. The alloy was obtained by melting in an induction furnace using the following materials: 99.995 wt.% Zn, 99.95 wt.% Mg and master alloys Mg-14 wt.% Ca. Melting was carried out in a graphite–chamotte crucible in air without the use of protective fluxes or atmospheres. The resulting melt was cast into a steel mold with a diameter of 35 mm and a height of 150 mm. Quantitative analysis of the elemental composition of the alloy was carried out on a BRUKER S8 Tiger sequential X-ray fluorescence wave-dispersive spectrometer (series 2; Bruker, Karlsruhe, Germany) in vacuum according to the standard procedure using the QUANT-EXPRESS^TM software (Bruker, Karlsruhe, Germany). The chemical composition of the studied alloy presents in

Table 1. The chemical composition of the alloy.

Alloy	Zn, wt.%	Mg, wt.%	Ca, wt.%
Zn-Mg-Ca	Balance	1.06 ± 0.02	0.12 ± 0.01

.

The resulting cast alloy was subjected to homogenizing annealing at 340 °C for 20 h and farther quenching by water. The disks for HPT (20 mm in diameter and 1.5 mm thick) were cut from the ingot. HPT was carried out at room temperature under a pressure of 4 GPa using the Bridgman anvils (for details see [30]). Deformation was carried out at a rate of 1 revolution per minute in a hole 0.9 mm in deep. The total number of HPT turns were 10. A true strain on the half-radius of the disks is ε = 5.7. The initial microstructure was studied using a JSM-7001F (JEOL; Tokyo, Japan) scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS). The microstructure of the alloy after HPT was studied by transmission electron microscopy (TEM) on a JEM-2100 (JEOL; Tokyo, Japan) transmission electron microscope operated at 200 kV. The phase compositions of the alloy were determined by X-ray diffraction (XRD). The studies were performed on a Bruker D8 Advance diffractometer (CuK_α radiation, λ = 1.54 Å; Bruker, Karlsruhe, Germany). X-ray diffraction patterns were processed by the Rietveld method using a Bruker DIFFRAC.EVA^TM (Bruker, Karlsruhe, Germany), DIFFRAC.TOPAS^TM (Bruker, Karlsruhe, Germany) and ICDD PDF-2 2020 (ICDD, Newtown Square, PA, USA) software.

The corrosion resistance was evaluated by the electrochemical method on an SP-300 potentiostat (Bio-Logic SAS, Seyssinet-Pariset, France). The studies were carried out in 0.9% NaCl solution (pH = 7) at room temperature using a flat PAR cell (Ametek Instruments, Oak Ridge, TN, USA) with a “three-electrode configuration” (working electrode (sample), Ag/AgCl reference electrode and counter electrode from Pt grid). The working surface area was 0.8 cm². The scanning was carried out at a rate of 1 mV/s. The scan range was from −150 mV below open circuit potential to +500 mV above open circuit potential. The time for determining the open circuit potential was 10 min. Six scans were performed for each test sample. Corrosion potential, corrosion current density, and degradation rate (DR) were calculated using the EC-Lab program (BioLogic, Seyssinet-Pariset, France). The calculation of DR was carried out based on the values of the corrosion current density (1) [31]:

$$\mathrm{DR}=\frac{3.27·{10}^{-3}·{\mathrm{i}}_{corr}· \mathrm{EW} }{\mathsf{\rho }· \mathrm{S}}$$

(1)

where DR is the degradation rate (mm/y), i_corr is the corrosion current density, μA/cm²; EW—equivalent weight, g/eq; ρ is the density of the alloy, g/cm³, S—the corroded surface of the alloys, cm².

Uniaxial tensile tests were performed using an Instron 3382 testing machine (Instron, High Wycombe, UK). The tests were carried out at room temperature with an deformation rate of 1 mm/min. Flat samples (n = 3 per state) with a cross-section of 2 mm × 1 mm and a gauge length of 5.75 mm were used for mechanical properties evaluation.

Hemolysis and cytotoxicity were evaluated as described earlier in [32,33]. Evaluation of bioactive properties in vitro was carried out on samples in the form of a 1/8 disk. A diameter of this disk was 20 mm and a thickness was ~ 1 mm. All alloy samples were pretreated by immersion in 70% ethanol for 4 h for studies in vitro. In other words, red blood cells and mononuclear leucocytes (MLs) were separated from mouse whole blood stabilized with 60 IU/m heparin, washed with sterile phosphate-buffered saline (PBS) (PanEco, Moscow, Russia) and suspended in PBS or the complete growth medium based on DMEM, respectively. An amount of 2 mL cell suspension was added to the alloy samples (alloy-treated cells). Cells incubated without alloys were used as control. Induced hemolysis was assessed 4 h later, and cytotoxicity 24 h after the start of co-incubation of alloys and cells at 37 °C in an atmosphere of 5% carbon dioxide. To assess viability, MLs lactate dehydrogenase (LDH) activity was studied using Pierce LDH Cytotoxicity Assay Kits (Thermo Scientific, Waltham, MA, USA) in accordance with the manufacturer’s method, measuring adsorption at 450 nm against 620 nm with the plate reader (Spark, Tecan, San Jose, CA, USA). Hemolysis was evaluated considering the adsorption of the supernatant at 540 nm. The result of the studies was presented as the ratio of the measured adsorption to the control (% vs. control).

The cell colonization on the surface of the alloy was investigated with multipotent mesenchymal stromal cells (MMSCs). The MMSCs were generated from C57Bl/6 mouse mice femur bone marrow as described earlier and suspended in the complete growth medium based on DMEM [34]. A total of 20 μL of the cell suspension was applied to the surface of each alloy sample. In control, the cells were seeded on the bottom of the empty wells at the same volume. After 30 min of incubation, the complete growth medium (2 mL) was added to each well and further incubated for 8 days at 37 °C in an atmosphere with 5% carbon dioxide. The medium in the wells was replaced with a fresh portion every 2 days. Then, the LDH activity of cells on the alloy surface was examined by measuring adsorption at 450 nm against 620 nm (A 450–A620).

The antibacterial activity of alloys was tested with Escherichia coli (E. coli) from collection of the N.N. Blokhin NMRCO. For our experiments we used a 18 h culture of the bacteria in Mueller–Hinton broth (Pronadisa, Madrid, Spain) with an absorption corresponded to 0.420 ± 12 c.u. at 440 nm (100 μL suspension volume). Alloy samples (n = 3) were placed in the wells of a 24-well Nuclon plate (Thermo Scientific, Waltham, MA, USA). Then, 20 μL of the bacterium suspension was applied to the surface of each samples or in empty wells (control). The plate was preincubated for 30 min, and then 1 mL Muller–Hinton broth was carefully added in each well. The plate with controls and alloys was incubated for 1 day at 37 °C in an atmosphere with 5% carbon dioxide. For assessment of the results we used the alamar blue reagent (Invitrogen, Thermo Scientific, Waltham, MA, USA) for registration of fluorescence at Ex530/Em590 with a Spark reader (Tecan, Waltham, MA, USA) in accordance with the manufacturer’s instructions. The protocol of the experiments with animals and cells was evaluated and approved by the Local Ethics Committee of the N.N. Blokhin NMRCO (project#660, Agreement 075-15-2021-965, 24 September 2021).

The statistical analyses were carried out using Statistica 6.0 (StatSoft, Tulsa, OK, USA). The experimental results were presented as a mean ± standard deviation (M ± SD). All measurements were made at least in triplets. One-way ANOVA was used for intergroup comparisons, and Student’s t-test was used for pairwise comparisons with controls. Difference was considered significant at p < 0.05.

3. Results and Discussion

Figure 1 shows the results of studying the microstructure of the alloy in the annealed state (a,b) and after HPT (c–f). The structure of the alloy in the annealed state consists of irregularly shaped α-Zn dendritic cells with an average size of 20.3 ± 1.1 μm and an eutectic phase surrounding them. The lamellae of the eutectic phase completely dissolve after homogenization and it becomes more homogeneous (Figure 1b). A significant refinement of these α-Zn dendritic cells with the formation of an ultrafine-grained (UFG) microstructure occurs after HPT. The alloy structure after HPT consists of regularly shaped recrystallized α-Zn grains with an average size of 750 ± 30 nm (Figure 1c) and a refined eutectic phase surrounding these grains (Figure 1d). It should be noted that the presence of extinction contours is observed along the grain boundaries of α-Zn, which indicates the processes of relaxation of the accumulated stresses along the boundaries. In addition, the grains were found to have well-defined high-angle boundaries. The phase located along the boundaries was also refined. The phase boundary zone consisted of fine grains of the refined phase with an average size of 38 ± 7 nm (Figure 1d and Figure 2a). It should be noted that the finely dispersed rounded particles ~20 nm in size were also found in the structure after HPT. Previously, it was shown in [35] that natural aging can occur in the Zn–0.08%Mg alloy at room temperature. It is likely that natural aging also occurred in our case, which caused the particles’ precipitation. EDS analysis showed that the particle is composed of Zn and Ca (Figure 1f).

Figure 2 shows the TEM-EDS results for different regions of the alloy. According to the results, the structure of the alloy after HPT consists of grains of pure zinc (spectrum 1) and a mixture of phases based on Zn, Mg and Ca (spectra 2–6). An X-ray phase analysis of the alloy was carried out before and after HPT to determine the stoichiometric composition of these phases.

Figure 3 and

Table 2. Results of XRD analysis (Ritveld method) of the Zn-1%Mg-0.1%Ca alloy in the initial state and after HPT processing.

State	Phase	Space Group	Mass Fraction, wt.%	a (Å)	c (Å)
Annealing	Zn	P63/mmc (194)	72.9 ± 9.3	2.666 ± 0.001	4.950 ± 0.001
	Mg2Zn11	Pm$\overline{3}$(200)	5.9 ± 1.3	8.543 ± 0.001	8.543 ± 0.001
	MgZn2	P63/mmc (194)	8.9 ± 1.4	5.201 ± 0.005	8.886 ± 0.010
	CaZn13	Fm$\overline{3}$c (226)	6.0 ± 2.0	12.183 ± 0.002	12.183 ± 0.002
	CaZn11	I41/amd:2 (141_O2)	6.3 ± 1.9	10.945 ± 0.156	6.916 ± 0.088
HTP	Zn	P63/mmc (194)	68.9 ± 12.6	2.666 ± 0.001	4.949 ± 0.001
	Mg2Zn11	Pm$\overline{3}$ (200)	11.3 ± 4.9	8.547 ± 0.003	8.547 ± 0.003
	MgZn2	P63/mmc (194)	6.7 ± 2.5	5.048 ± 0.004	8.891 ± 0.012
	CaZn13	Fm$\overline{3}$c (226)	5.2 ± 2.6	12.253 ± 0.002	12.253 ± 0.002
	CaZn11	I41/amd:2 (141_O2)	7.6 ± 3.1	10.812 ± 0.008	6.734 ± 0.013

show the results of qualitative and quantitative XRD analysis of the alloy. The studies have shown that there are four types of phases (Mg₂Zn₁₁, MgZn₂, CaZn₁₃ and CaZn₁₁), except for pure zinc, present in the Zn-1%Mg-0.1%Ca alloy in both states. The volume fraction of these phases is presented in Table 2. It should be noted that the volume fraction of each phase after HPT generally remains unchanged (within the experimental error). In addition, the studies have shown that the basal texture component is intensified in the alloy after HPT. This is indicated by the change in the height of the intensity of the lines (002) and (004) on the X-ray pattern of HPT-treated alloy compared to the annealed state (Figure 3).

The study of the corrosion resistance of the alloy is presented in Figure 4. The study showed that HPT does not lead to a deterioration in the corrosion resistance of the Zn-1%Mg-0.1%Ca alloy. Thus, the corrosion potential of the annealed alloy was −1009 ± 5 mV, while that of the alloy after HPT was −1023 ± 13 mV. At the same time, the corrosion current density was 13.4 ± 5.5 and 12.4 ± 4.1 µA/cm², and the degradation rate (DR) was 0.20 ± 0.08 and 0.19 ± 0.06 mm/year for the alloy before and after HPT, respectively. It was noted that after HPT, there is no change in the phase composition of the alloy, as well as in the volume fraction of these phases. This could be the reason that deformation does not lead to a deterioration in corrosion resistance. In this case, a strong refinement of the microstructure is compensated by the formation of grains with high-angle boundaries and the crushing of a mixture of phases located along the boundaries of the α-Zn zones.

It is known that the distribution of the deformation over the cross section of the disk is not uniform: the farther from the center of the disk a point is, the greater the degree of deformation [30]. This leads to the inhomogeneous structure along the sample. In order to increase the uniformity of the structure after HPT, the number of revolutions is usually increased. However, the formation of a larger microstructure is usually observed in the center of the sample after HPT, which becomes smaller with increasing distance from the center. To assess the homogeneity of the microstructure, the change in microhardness along the sample diameter was studied, and the microhardness matrices of the sample surface in contact with the fixed deforming rod were also measured. The results of the study are presented in Figure 5. The study showed that the microhardness of the alloy after HPT has a lower value in the center of the sample. However, when moving away from the center by more than 2 mm and up to the edge of the sample, the microhardness becomes constant within the measurement error. It should also be noted that the microhardness of the alloy after HPT increased from 838 ± 35 to 1125 ± 25 MPa.

Table 3. Mechanical properties of the Zn-1%Mg-0.1%Ca alloy in the initial state and after HPT processing.

State	YS, MPa	UTS, MPa	El, %
Annealing	135 ± 13	154 ± 5	0.4 ± 0.1
HPT	356 ± 15	416 ± 31	5.5 ± 2.8

and Figure 6 present the results of a study of the mechanical characteristics of the alloy before and after HPT. The samples for testing the alloy after HPT were cut out from the zone corresponding to the half of the disk radius. The studies have shown that microstructure refinement during HPT leads to a significant increase in the strength of the Zn-1%Mg-0.1%Ca alloy. Therefore, the yield stress (YS) increases from 135 ± 13 to 356 ± 15 MPa, and the ultimate tensile strength (UTS) rises from 154 ± 5 to 416 ± 31 MPa. An increase in ductility (El) from 0.4 ± 0.1 to 5.5 ± 2.8% was also observed. The increase in the ductility of the alloy can be associated with several factors. First, this is the fragmentation of large clusters of the eutectic phase with the formation of a more uniformly distributed phase. Secondly, there are textural changes in the alloy during deformation. It was previously shown that the conventional extrusion of the Zn-1%Mg alloy leads to an increase in its ductility from 2 ± 0.5 to 32 ± 2% due to the weakening of the texture, the activation of non-basal slip systems, high values of the Schmid factors (SF) of prismatic and pyramidal slip systems, and a bimodal structure formation [36]. It was possible to achieve ductility of more than 45% due to grain refinement in the Zn-0.1%Mg alloy after equal-channel angular pressing, because of grain boundary sliding, and weakening of the texture intensity and increase in SF for pyramidal slip [37].

To assess the effect of HPT on the bioactivity of the alloy, an increase in hemolytic activity and a decrease in the viability of white blood cells after co-incubation were evaluated in comparison with the control. Statistical analysis of the obtained data presented in Figure 7a showed that the alloy in both states does not demonstrate essential hemolytic activity, since there is no significant difference from the control (p > 0.05). Therefore, it can be concluded that HPT of the Zn-1%Mg-0.1%Ca alloy does not lead to a significant change in the activity (p > 0.05).

In order to investigate the possible impact of HPT on other parameters of biocompatibility, their effect on the viability of MLs after in vitro incubation was studied. The results showed that the incubation of cells with the Zn-1%Mg-0.1%Ca alloy in the annealed state and after HPT leads to a significant decrease in cell viability: p = 0.01 and p = 0.03 compared with the control, respectively. However, HPT contributed to a decrease in cytotoxicity by an average of 10% in comparison with the initial alloy (p = 0.048). In general, it can be stated that the conducted studies revealed signs of a significant improvement in the biocompatibility of the Zn-1%Mg-0.1% Ca alloy after HPT (Figure 7b).

For a more in-depth assessment of the effect of HPT on the biological properties of the alloy, their ability to stimulate the colonization of sample surfaces by MMSCs with osteogenic potential, as well as to exert an antibacterial effect, were studied.

Zinc is known to have antibacterial properties at high concentrations [16]. It could be expected that the release of zinc ions into the incubation medium would be able to inhibit the growth of the bacterial culture or even would have a bactericidal effect. To test this hypothesis, the studies were conducted on a test culture of E. coli, Gram-negative microorganisms that often mediate the development of infectious processes in bone tissues [38]. According to the results of the bacteriological studies, there was no slowdown in the growth of the model culture of microorganisms incubated in the presence of alloy samples in comparison with the control, where the test culture of microorganisms was cultivated at the bottom of the well of the plate without samples (p > 0.05) (Figure 7c). This allows us to conclude that the Zn-1%Mg-0.1%Ca alloy has no pronounced antibacterial properties. It should be noted that there was no significant change in the activity of samples after HPT in comparison with the annealed state (p > 0.05). Probably, the degradation rate of the alloy in this case is not able to ensure a sufficient amount of Zn²⁺ ions to provide an antibacterial effect.

The study of the stimulation of cell colonization on sample’s surface of the alloy is presented in Figure 7d. MMSC (an adhesive cell culture with osteogenic potential) was used as a cell model. Statistical analysis of the obtained data showed that HPT does not lead to a significant inhibition of cell colonization of the surface of the samples of the alloy in comparison with the annealed state (p > 0.05).

4. Conclusions

The conducted studies allow us to draw the following:

• HPT leads to the formation of an UFG structure with an average α-Zn grain size of 750 ± 30 nm. In addition, fragmentation of the interlayer of phases located along the initial dendrites of α-Zn is observed with the formation of grains with a size of 38 ± 7 nm. The precipitation of Ca-enriched particles with a size of ~20 nm also occurs.
• The presence of five phases (α-Zn, Mg₂Zn₁₁, MgZn₂, CaZn₁₁ and CaZn₁₃) in the Zn-1%Mg-0.1%Ca alloy, the volume fraction of which does not change after HPT, is determined by XRD analysis.
• HPT does not lead to a deterioration in the corrosion resistance of the Zn-1%Mg-0.1%Ca alloy.
• HPT leads to an increase in strength (YS increases from 135 ± 13 to 356 ± 15 MPa, and UTS rises from 154 ± 5 to 416 ± 31 MPa) and ductility (from 0.4 ± 0.1 to 5.5 ± 2.8%) of the Zn-1%Mg-0.1%Ca alloy.
• HPT does not worsen the bioactivity of the Zn-1%Mg-0.1%Ca alloy in vitro, in particular, reducing its cytotoxicity compared to the annealed state.

The obtained results can serve as a rationale for choosing the Zn-1%Mg-0.1%Ca alloy after HPT as the basis for submersible biodegradable implants, metal structures and fixators. They will be in demand, in particular, in the process of treatment and rehabilitation of cancer patients, involving osteoreconstructive operations.