The as-cast Zn–0.8Li–0.5Mn (wt%) alloy ingots were supplied by the Hunan Institute of Rare Earth Metal Materials. Chemical composition was confirmed by inductively coupled plasma atomic emission spectroscopy (ICP-AES, iCAP6300, United States), showing actual Li and Mn contents of 0.81 wt% and 0.49 wt%, respectively. The phase components of the as-cast Zn-Li-Mn alloy was characterized by X-ray diffraction (XRD, D8 Discover, Bruker, Germen) at a scanning rate of 2°/min and scanning range from 30°–80°. Rectangular billets (10 × 10 × 90 mm) were cut from the ingots and processed through equal-channel angular pressing (ECAP) according to established procedures (Bednarczyk et al., 2019; Bednarczyk et al., 2018). ECAP with a Bc route was conducted at 200 °C, with specimens subjected to 8, 12, or 16 passes, including a 5 min holding step after every fourth pass. Following ECAP, plates (10 × 10 × 2 mm) were extracted for systematic investigations, including microstructural observations, electrochemical and immersion studies, cellular assays, osteogenic evaluation, and antibacterial testing.

To elucidate the structural features of the processed alloys, multiple advanced techniques were applied. Optical microscope (OM, BX53M, Olympus, Japan), Field-emission scanning electron microscopy (FE-SEM, Helios 5 CX, Thermo Fisher Scientific, United States) and electron backscatter diffraction (EBSD, e-Flash, Bruker, Germany) were used for microstructural mapping. Elemental distributions were analyzed using energy-dispersive X-ray spectroscopy (EDX, X-Flash 7, Bruker, Germany) attached to the SEM. For metallographic preparation, samples were ground, polished, and etched with an ethanol-based solution containing 4 wt% nitric acid. The polished samples are necessary to electrically polish before EBSD. The electrolytes are composed of 50% phosphoric acid and 50% ethanol. The electrically polish process was conducted at −30 °C and 15V for 15–20 s. Step size of EBSD patterns was set as 0.07 μm. Thin foils (1 mm) were mechanically reduced to ∼50 μm using silicon carbide abrasive papers, followed by twin-jet polishing. The solution for twin-jet polishing is composed of 10% perchloric acid and 90% ethanol. TEM samples were prepared at liquid nitrogen environment. Final thinning for transmission electron microscopy (TEM, Talos F200X, Thermo Fisher Scientific, United States) was achieved by low-temperature ion milling (Fischione, United States).

Tensile behavior was evaluated on a universal testing system (Autograph AGS-X, Shimadzu, China) at a constant strain rate of 10⁻³ s^-1. Following ASTM-E8 guidelines, miniature dog-bone specimens with dimensions of 2 × 3 × 10 mm were prepared and tested at room temperature. Fractured surfaces were subsequently analyzed using scanning electron microscopy (SEM) to assess failure morphologies.

Electrochemical behavior was assessed using a standard three-electrode setup connected to an electrochemical workstation. Open-circuit potential was monitored for 3,600 s, followed by potentiodynamic polarization (PDP) scans from −1.4 V to −0.6 V at 0.1 mV s^-1. Both electrochemical measurements and immersion experiments were performed in Hank’s solution for up to 30 days, maintaining a solution-to-sample surface ratio of 20 mL/cm². After immersion, specimens were rinsed with deionized water, air-dried, and corrosion products were removed using a 200 g/L chromium trioxide (CrO₃) solution prior to further analysis. The corrosion rates of ECAPed Zn alloys were calculated from the equation: v c o r r = Δ m ρ A T . Where Δ m was weight loss after removing corrosion products, ρ was the density of the studied Zn alloys, A was the exposed area, and T was immersion time.

Cytotoxicity and osteogenic assays were performed with murine pre-osteoblast MC3T3-E1 cells (ATCC). Cells were cultured in α-MEM supplemented with 10% fetal bovine serum (FBS) and 1% penicillin–streptomycin under standard conditions (37 °C, 5% CO₂, 95% humidity). Medium was refreshed every 2 days. Prior to testing, alloy disks were ultrasonically cleaned and UV-sterilized for ≥4 h per side. Extracts were prepared by incubating sterile samples in complete medium at a ratio of 1.25 mL/cm² for 24 h, centrifuged to remove particulates, and stored at 4 °C. Zn²⁺ and Mg²⁺ ion concentrations in the extracts were quantified by ICP-AES (Thermo Scientific M Series, United States).

For viability analysis, cells were harvested by trypsinization, resuspended, and exposed to undiluted or 50% diluted extracts for 3 and 5 days. Cell proliferation was measured using the CCK-8 assay (Dojindo, Japan). Live/dead staining with Calcein-AM/PI was performed, and cellular morphology was visualized with an inverted fluorescence microscope (Zeiss Axio Vert.A1, Germany).

Osteogenic differentiation was examined via alkaline phosphatase (ALP) activity and extracellular matrix mineralization. Cells were cultured in osteogenic medium mixed 1:1 with extracts, with medium renewal every 48 h. After 7 days, ALP staining was conducted; after 14 days, mineral deposition was assessed by Alizarin Red S (ARS) staining following fixation with 3.7% formaldehyde. ARS-bound calcium was visualized and quantified spectrophotometrically. In addition, total RNA was extracted with RNAiso Plus (TaKaRa, Japan), reverse-transcribed into cDNA, and analyzed by qPCR to evaluate osteogenic gene expression.

All experimental data was repeated for at least three times independently, and presented in form of mean ± standard deviation (SD). Besides, the one-way analysis of variance (One-way ANOVA) and the following Tukey’s post hoc test were conducted to manifest the differences between groups.

Figure 1 illustrates the phase component and microstructure of the as-cast Zn alloys. According to the Zn–Li phase diagram, eutectic and eutectoid transformations occur at 401 °C and 65 °C, producing β-LiZn₄ and α-LiZn₄ phases (Liang et al., 2008; Okamoto, 2012). The Zn–Mn phase diagram indicates the formation of MnZn₁₃ phases (Shi et al., 2018). Thus, the studied alloys primarily contain LiZn₄ and MnZn₁₃. XRD patterns (Figure 1a) confirmed these phases, showing prominent diffraction peaks of Zn at (0002), (01 1 ¯ 0), and (01 1 ¯ 1) planes (36.5° and 39.1°), alongside peaks attributed to LiZn₄ and weaker signals assigned to MnZn₁₃. OM (Figure 1b) and SEM images (Figures 1c,d) revealed lamellar eutectoid structures, where lath Zn alternates with LiZn₄ (Cao et al., 2023). The lath thickness and spacing ranged between 300 and 500 nm (Figure 1d).

OM images (Figure 2a) revealed that after ECAP, the Zn matrix and second phases were elongated along the extrusion direction (ED), forming alternating bands. Low-magnification SEM images (Figure 2b) showed a mixture of dynamically recrystallized (DRX) grains and elongated second phases, appearing as granular and strip-like features aligned with ED. With increasing passes, regions consisting of DRX grains became progressively thinner, and secondary phases were fully refined. Figure 2c highlights second-phase evolution: in 8p- and 12p-ECAP samples, fine particles coexisted with lath structures which were almost absent from the microstructures of 16p-ECAP, indicating complete refinement at higher deformation.

EBSD analysis (Figure 3) further demonstrated microstructural changes. Inverse pole figure (IPF) maps displayed equiaxed ultrafine grains with varied orientations: red for (0001), blue for (01 1 ¯ 0), and green for ( 1 ¯ 2 1 ¯ 0). DRX was evident in all samples. It can be measured from Figure 3a that the average grain sizes were 1.41 μm for 8p-ECAP sample, 1.36 μm for 12p-ECAP sample, and 0.69 μm for 16p-ECAP sample, respectively. Moreover, the fraction of fine grains (<5 μm) reached 64% for 8p-ECAP, 38% for 12p-ECAP, and 100% for 16p-ECAP (Figure 3b). Kernel average misorientation (KAM) maps (Figures 3c,d) showed average values of 0.75°, 0.90°, and 0.59° for 8p-, 12p-, and 16p-ECAP, respectively.

Due to Zn’s low melting point, grain refinement primarily results from dynamic recrystallization (DRX). DRX grains contained fewer dislocations, reflected by low grain orientation spread (GOS <1°). GOS mapping (Figures 4a,b) indicated DRX fractions of 42.9%, 24.2%, and 80.2% for the 8p-, 12p-, and 16p-ECAP samples, respectively, suggesting greater DRX activity at higher passes. Pole figure (PF) analysis (Figure 4c) revealed a pronounced basal texture in ECAP-processed alloys. Texture intensity first increased from 26.21 (8p) to 74.96 (12p), then dropped sharply to 11.05 (16p), consistent with higher DRX grain fractions and reduced dislocation density.

TEM observations confirmed abundant fine phases in ECAP-processed alloys. In the 12p-ECAP sample, selected area electron diffraction (SAED) identified LiZn₄ as the second phase (Figure 5a). High-resolution TEM and SAED also detected MnZn₁₃ precipitates (Figure 5b), frequently located at DRX grain boundaries (Figure 5c). EDS confirmed Mn enrichment in these precipitates relative to the Zn matrix. ImageJ measurements indicated an average precipitate size of 98 ± 29 nm, significantly smaller than the surrounding DRX grains.

Tensile testing was conducted to evaluate the mechanical performance of ECAP-processed Zn alloys. The representative stress–strain curves (Figure 6a) exhibited clear elastic and plastic stages. In the elastic region, stress increased linearly with strain, followed by a sharp rise to the ultimate tensile strength (UTS), and then a decline. Comparisons of UTS and elongation are presented in Figures 6b,c. The 8p-ECAP sample showed a UTS of 416 MPa, which increased to 433 MPa for the 12p-ECAP and further to 450 MPa for the 16p-ECAP. In contrast, ductility decreased with additional passes, dropping from 76.6% (8p) to 75.2% (12p) and 59.8% (16p). Notably, the 12p-ECAP alloy achieved the most favorable balance between strength and ductility. Fractography (Figure 6d) revealed abundant dimples on the fracture surface, typical of ductile failure, confirming that all ECAP-treated alloys fractured via plastic deformation mechanisms.

Electrochemical measurements are summarized in Figure 7. Potentiodynamic polarization (Figure 7a) and Nyquist plots (Figure 7b) revealed that corrosion potential (E_corr) shifted positively and corrosion current density (i_corr) decreased with increasing ECAP passes. Tafel analysis (Figures 7c,d) showed that the 8p-ECAP alloy had the lowest E_corr (−1.194 V_SCE) and the highest i_corr (10.93 μA/cm²). After 12 passes, Ecorr shifted to −1.188 VSCE with an icorr of 10.53 μA/cm², while the 16p-ECAP exhibited the most positive E_corr (−1.149 V_SCE) and the lowest icorr (8.17 μA/cm²). Nyquist plots confirmed larger semicircle diameters for the 16p sample, indicative of higher charge-transfer resistance. Equivalent circuit fitting (inset of Figure 7b) included solution resistance (R_s), corrosion film resistance and capacitance (R_f, CPE_f), and double-layer resistance and capacitance (R_ct, CPE_dl). The data (Table 1) showed increasing R_ct with ECAP passes, suggesting improved corrosion protection, despite a reduction in R_f values.

Immersion experiments (Figure 7e) further supported these findings. Corrosion rates declined from 34.60 μm/year in 8p-ECAP to 32.83 μm/year in 12p-ECAP, and to 26.66 μm/year in 16p-ECAP, confirming enhanced corrosion resistance with additional passes.

Surface morphologies after immersion in Hank’s solution are presented in Figure 8. After 30 days, all alloys developed corrosion films with uniformly distributed cracks (Figure 8a). EDS analysis (Figure 8b) identified corrosion products containing C, O, P, Ca, and Zn. Removal of surface films (Figure 8c) revealed numerous pits characteristic of localized corrosion. The number density of corrosion pits decreased with more ECAP passes. Moreover, the high-magnification SEM images show that the sizes of corrosion pits are a few micrometers and the grain boundary corrosion cracks are observed. In contrast, as-extruded alloys displayed large, irregular pits indicative of severe localized attack. These results collectively demonstrate that ECAP effectively reduces corrosion susceptibility and promotes the formation of more protective surface layers.

The biocompatibility of ECAP-processed Zn alloys was evaluated using MC3T3-E1 pre-osteoblasts, with titanium (Ti) serving as the control. Cell viability after 3 days of culture in alloy extracts is shown in Figure 9a. At full-strength extracts (100%), Zn alloys initially exhibited lower activity compared to Ti. However, when diluted to 50%, cell viability improved significantly, exceeding that observed in undiluted conditions. After 5 days (Figure 9b), viability in the diluted extracts stabilized around 100%, confirming good cytocompatibility. Live/dead staining (Figure 9c) further supported these findings: cells cultured in 50% extracts for 3 and 5 days showed green fluorescence indicative of viable cells. Although cell density in Zn alloy groups was slightly lower than that of Ti, cells exhibited healthy morphology and normal spreading behavior, demonstrating that the alloys provided a supportive environment for proliferation.

The osteogenic potential of Zn alloys was examined through alkaline phosphatase (ALP) activity at day 7 and Alizarin Red S (ARS) staining at day 14. As shown in Figures 10a–d, both ALP activity and mineral deposition were higher in Zn alloy groups compared with Ti, indicating enhanced early osteogenic differentiation. This effect is attributed to the gradual release of metallic ions during alloy degradation, which are known to stimulate bone-forming pathways. Gene expression analysis further confirmed these observations (Figure 10e). In MC3T3-E1 cells exposed to 12p-ECAP extracts, osteogenesis-related markers—including ALP, collagen I (COL-I), osteopontin (OPN), and Runx-2—were consistently upregulated relative to Ti. Collectively, these results demonstrate that ECAP-processed Zn alloys significantly promote osteogenic differentiation, underscoring their promise as next-generation biodegradable orthopedic implant materials.

Previous reports on Zn alloys processed by ECAP or rolling have demonstrated that yield strength ( σ y ) is mainly governed by grain boundary strengthening ( σ G B ), dislocation strengthening ( σ ρ ), and second-phase strengthening ( σ S ) (Ren et al., 2021; Huang et al., 2020; Huang et al., 2021; Zhuo et al., 2023). The same mechanisms operate in this work. Severe plastic deformation during ECAP promotes extensive dynamic recrystallization (DRX), resulting in significant grain refinement. This contribution can be estimated through the Hall–Petch relationship (Figueiredo et al., 2023): σ G B = σ 0 + k d − 1 2 . Where the σ 0 is the intrinsic strength (32.2 MPa for pure Zn), k (80 MPa μm^-0.5) is the Hall–Petch relation slope, and d is the average grain size (Wang et al., 2021; Diaa et al., 2025). As shown in Figures 11a,b, the DRX grain sizes were 1.41 μm (8p-ECAP), 1.36 μm (12p-ECAP), and 0.69 μm (16p-ECAP), corresponding to σGB contributions of 63.7, 38.3, and 120 MPa, respectively.

Plastic deformation also increases dislocation density, which can be estimated from KAM values using (Bednarczyk et al., 2023): ρ = 2 θ ¯ μ b . Where b is Burger’s vector (0.266 nm) and μ is the tangent line step unit (0.08 μm). Figure 11c shows that non-DRX regions (GOS >1°) had dislocation densities of 17.46 × 10¹⁴ m^-2 (8p-ECAP), 18.50 × 10¹⁴ m^-2 (12p-ECAP), and 11.39 × 10¹⁴ m^-2 (16p-ECAP), as shown in Figure 11c. The strengthening contribution is obtained from (Ma et al., 2014): σ ρ = α M b G ρ . Where α = 0.2, M = 2.76 (Taylor factor for Zn), and G = 42 GPa.

The calculated σ ρ values are 147.1 MPa (8p-ECAP), 201.1 MPa (12p-ECAP), and 41.2 MPa (16p-ECAP). Thus, the strength contributions from grain boundary strengthening are smaller than those from dislocation strengthening in the 8p and 12p Zn alloys. For the 16p alloy, the grain boundary strengthening contributes a higher strength value compared with dislocation strengthening. As shown in Table. 2, the measured tensile yield strengths (TYSs) are higher than the sum of strength contributions from grain boundaries and dislocations. This is attributed to the formation of secondary phases, including high fractions of LiZn₄ phases and nanosized MnZn₁₃ phases (Figure 5b). It can be estimated that the strength contributions from secondary phases are 101.2 MPa (8p-ECAP), 34.6 MPa (12p-ECAP), and 168.8 MPa (16p-ECAP), respectively. Consequently, it can be referred that the ECAP passes play a critical role in the transitions of strengthening mechanism. Dislocation and secondary phase strengthening are dominated in 8p- and 12p-ECAP Zn alloys, while grain boundary and secondary phase strengthening are dominated in 16p-ECAP Zn alloys.

The large number of second-phase particles, including LiZn₄ and MnZn₁₃, also plays a crucial role. These particles can pin dislocations, act as nucleation sites for DRX through the particle-stimulated nucleation (PSN) mechanism, and hinder grain boundary migration, all of which promote ultrafine-grained microstructures (Huang et al., 2018). The superior strength of the 16p-ECAP alloy is therefore linked to this additional refinement. Moreover, nano-sized MnZn₁₃ precipitates at grain boundaries strengthen the matrix through a load transfer effect (Ding et al., 2021). Their deformation incompatibility with the Zn matrix restricts dislocation motion (Azizian et al., 2024) while pre-existing dislocation walls serve as barriers to newly generated dislocations. Together, these mechanisms reinforce the strengthening effect, explaining the high yield strength achieved after 16 passes of ECAP. Table 3 shows the mechanical properties of as-cast and thermomechanical process Zn-Li-Mn alloys (Shen et al., 2025; Yang et al., 2024; Yang et al., 2023; Zh et al., 2024; Yang et al., 2020). The tensile strengths of the majority of Zn-Li-Mn alloys are located in the range of 400–500 MPa after thermomechanical process. After ECAP, tensile strengths of Zn-Li-Mn are below 450 MPa, while the elongations (76.6% for 8p-ECAP sample) are higher than that of most of Zn-Li-Mn alloys.

Weight-loss measurements (Figure 7e) revealed that corrosion rates decreased with increasing ECAP passes. Corrosion in Zn alloys is strongly influenced by grain boundary density, second-phase characteristics, and dislocation density. Grain refinement after ECAP increases boundary density, which can act as preferential sites for corrosion initiation, generally accelerating dissolution (Ralston et al., 2010). However, when dissolution rates fall below ∼10 μA/cm², finer grains instead enhance passivation and reduce corrosion. Electrochemical testing confirmed that the 16p-ECAP sample exhibited an i_corr of 8.17 μA/cm², consistent with this transition. Thus, grain refinement at high ECAP passes inhibits corrosion, explaining the improved resistance observed in the 16p-ECAP alloy compared with 8p- and 12p-ECAP samples. Moreover, dislocation density has a significant effect on corrosion behaviors of Zn alloys. With increasing dislocation densities, the surface energy is elevated, rendering the alloy more susceptible to activation upon contact with corrosive media (Feng et al., 2024). At low ECAP passes, a large number of dislocations are present in 8p-ECAP and 12p-ECAP alloys, which will reduce the corrosion performance. The 16p-ECAP alloys have finer grains and lower dislocation density, which results in superior corrosion resistance.

In addition to grain size effects, phase evolution also plays a role. With increasing ECAP passes, the originally continuous LiZn₄ phases were broken down into dispersed particles, while MnZn₁₃ precipitates were promoted. Previous studies demonstrated that the localized corrosion regions were observed in the regions of secondary phase of Zn-Li based alloys after ECAP (Huang et al., 2024). This is attributed to the occurrence of galvanic corrosion between the LiZn₄ phase and Zn. Although second-phase particles can increase galvanic coupling and accelerate localized attack, their dispersion reduces corrosion sensitivity and facilitates the formation of a more uniform and protective surface film. This film integrity contributes to the lower corrosion rates of highly deformed alloys.

Grain uniformity further influences degradation behavior. As shown in Figure 3, the 16p-ECAP alloy exhibited more homogeneous grain structures, which promoted uniform corrosion rather than localized pitting. Meanwhile, dislocation density decreased significantly with higher ECAP passes (Figure 3c), reducing micro-galvanic activity around dislocation tangles. Collectively, these changes suggest that the beneficial effects of refined, uniform grains and reduced dislocation density outweigh the detrimental influence of second-phase increase. As a result, 16-pass ECAP alloys demonstrated markedly improved corrosion resistance relative to alloys with fewer passes.

For decades, titanium (Ti) alloys, stainless steel, and cobalt–chromium (Co–Cr) alloys have served as the standard implants in load-bearing sites such as long bone fractures, femoral fixation, and joint replacements (Bandyopadhyay et al., 2023; Cui et al., 2024). These materials share common advantages, including high mechanical strength and reliable biocompatibility. Their tensile strength typically exceeds 400 MPa, ensuring sufficient structural support for bone repair, while their clinical biosafety is well established. Nevertheless, these inert implants present two major drawbacks. First, long-term implantation can trigger complications such as chronic inflammation or persistent pain, often necessitating secondary removal surgery (Reith et al., 2015). Second, their high elastic modulus creates a mismatch with the native bone, leading to stress shielding and progressive bone resorption, ultimately increasing the risk of revision surgery (Shayesteh Moghaddam et al., 2016).

In response, increasing attention has shifted toward biodegradable metallic materials with tailored mechanical and biological performance (Zheng et al., 2014). Magnesium (Mg)-based implants have progressed from laboratory studies to clinical translation after extensive research (Windhagen et al., 2013; Lee et al., 2016). Their in vivo degradation products actively stimulate bone regeneration and accelerate healing (Zhao et al., 2016), providing unique advantages in orthopedic repair. However, the relatively low strength of Mg alloys limits their use in high load-bearing environments.

The Zn–Li–Mn alloy system developed in this study demonstrated ultimate tensile strengths exceeding 400 MPa and elongations above 60%, comparable to Ti while maintaining sufficient ductility. The superior mechanical properties arise from combined contributions of grain boundary strengthening (via fine and ultrafine grains), dislocation strengthening (high dislocation density), and second-phase strengthening. Beyond mechanics, in vitro osteogenesis assays provided further validation. Both ALP and ARS staining confirmed stronger osteogenic induction in Zn alloys than in bioinert Ti, findings supported by the upregulation of osteogenic genes. This enhancement may be linked to the activation of the PI3K–AKT signaling pathway (Yang et al., 2021), a key regulator of bone formation.

Taken together, Zn–Li–Mn alloys integrate the mechanical robustness characteristic of traditional inert metals with the biological activity observed in Mg-based systems. This dual advantage highlights their potential as next-generation biodegradable implants for orthopedic applications, capable of addressing both structural and biological demands in bone repair.