Tailoring the microstructure of biodegradable Zn-Cu-Li wires via thermo-mechanical processing towards high-performance surgical staples

Abstract

Although biodegradable Zn alloy fine wires are promising for staples, most exhibit inadequate mechanical properties, and current studies remain preliminary. In this work, Zn-2Cu-0.8Li (wt%) alloy fine wires (0.22 mm) for staples with outstanding mechanical properties were fabricated via hot extrusion, multi-pass drawing at room temperature and annealing. The microstructure, property evolutions and application feasibility for staples were systematically studied. The drawing process induces dramatic elongation of the β-LiZn₄ matrix, accompanied by dynamic recovery (DRV) and continuous dynamic recrystallization (CDRX). The low-angle grain boundary fractions significantly increase and the average grain sizes dramatically decrease. The η-Zn distributes as fine equiaxed grain bands due to DRX. The tensile yield strength (TYS) and ultimate tensile strength (UTS) increase from 376 MPa and 423 MPa to 596 MPa and 631 MPa, respectively, due to grain boundary (GB) and texture strengthening. Meanwhile, the deformability remains good with fracture elongation (EL) of 25.4% owing to DRV, CDRX and the presence of η-Zn. The wires annealed at 120 °C for 1 h show optimal mechanical properties (TYS: 492 MPa, UTS: 537 MPa and EL: 44.2%). The wires exhibit uniform degradation mode with a higher degradation rate of 327 μm∙year⁻¹ than as-extruded wires due to increased GB densities. The fabricated staples show an ultimate tensile force of 1.86 N comparable to Ti staples. They can achieve satisfactory anastomosis of beagle gastric tissue, and show appropriate degradation properties in vitro and in vivo. These findings indicate that Zn-2Cu-0.8Li fine wires and staples are promising for clinical applications.

Graphical abstract

Highlights

• •Novel Zn-2Cu-0.8Li staples with strength comparable to Ti were developed.
• •The fine wires with outstanding strength exceeding 600 MPa were fabricated.
• •The microstructure consists of elongated and CDRXed β-LiZn₄ and fine equiaxed η-Zn.
• •The staples successfully achieved anastomosis of beagles, and the wound healed well.

1. Introduction

Surgical staples have been widely utilized in gastrointestinal anastomosis and organ resection surgeries. Compared with conventional manual sutures, the application of staples can effectively shorten the operation time, reduce tissue trauma and lower the risk of postoperative infection [1]. Currently, most clinically used staples are fabricated from inert metals. Once implanted, these inert metals will permanently exist in the human body after surgery, possibly leading to chronic inflammation and other long-term complications [2,3]. Biodegradable metals, including iron (Fe), zinc (Zn) and magnesium (Mg) and their alloys, have attracted more and more attention in recent years. They can gradually degrade after tissue healing, thereby avoiding the complications associated with inert metals. Among them, Zn and its alloys are promising candidates. They exhibit moderate degradation rates due to the intermediate standard electrode potential (−0.76 V) between Mg (−2.37 V) and Fe (−0.44 V). Furthermore, zinc is the second most abundant trace element in the human body, and plays vital roles in cellular activities essential for wound healing, ensuring good biocompatibility [4,5].

Recent preclinical studies have validated the potential of Zn alloy staples. In 2020, Amano et al. [6] developed Zn-(Cu)-(Mn)-Ti alloy staples with an average in vitro degradation rate of 20 μm∙year⁻¹ in HBSS solution. Successful anastomosis in rabbit gastric resection models was achieved without severe inflammatory reaction or complications. In 2021, Guo et al. [7] reported on Zn-Li-Mn alloy staples for gastrointestinal anastomosis of pigs, which exhibited a higher average in vitro degradation rate of 191 μm∙year⁻¹ in Hank's solution. The staples could maintain mechanical integrity within 8 weeks post-implantation, while avoiding anastomotic leakage and severe inflammation.

Apart from the in vivo studies of staples, more research has predominantly focused on optimizing the fabrication of Zn-based fine wires, as their properties crucially influence staple performance. Several Zn alloy system wires have been developed and fabricated, including Zn-Mg [[8], [9], [10], [11]], Zn-Cu [6,12], Zn-Li [7,13], and Zn-Mn [6]. Wang et al. [9] prepared Zn-0.02 Mg (wt%) alloy wires with a diameter of 0.8 mm by indirect extrusion and 54 passes cold drawing with the area reduction of 4-9% per pass. The mechanical properties of wires are tensile yield strength (TYS): 388 MPa, ultimate tensile strength (UTS): 455 MPa and fracture elongation (EL): 5.4%. Bai et al. [10] employed annealing treatment for Zn-Mg alloy wires during both the intermediate and final passes of drawing, improving the elongation of wires from below 10% to 12%-28%, while decreasing the tensile strength. Cheng et al. [12] prepared Zn-0.2Cu-0.1 Mg wires with a minimum diameter of 0.3 mm, which showed good mechanical properties with a TYS of 256 MPa, a UTS of 313 MPa and EL of 46%. Zhao et al. [13] fabricated Zn-0.1Li (wt%) alloy wires with a minimum diameter of 0.25 mm by hot extrusion and drawing. The wires exhibit a TYS of 238 MPa, a UTS of 274 MPa and EL of 17%, as well as good biocompatibility and appropriate degradation rates.

Despite the above progress, current research and efforts remain in the preliminary stage, and the development of zinc alloy fine wires with high performance still faces critical challenges, hindering their clinical application as surgical staples. The mechanical properties of most zinc alloy fine wires remain inadequate to meet the requirements of staple materials. The enough strength is demanded to afford sufficient force for staples to favorably penetrate tissue and closure tissue without leakage. The strength requirements are: UTS >240 MPa according to Chinese Pharmaceutical Industry Standard YY 0875-2013, higher criteria of TYS >200 MPa and UTS >300 MPa recommended by the recent literature [14], and more demanding requirements for applications involving high-density tissues or complex loading conditions. Meanwhile, the adequate deformability is demanded to accommodate both staple manufacturing and anastomosis processes. Furthermore, during wire processing, the detailed characterization and analysis of microstructure, mechanical property evolution and degradation behaviors are also lacking. Therefore, developing fine wires for staples with high quality, high strength and ductility and suitable degradation properties via proper processing techniques, and clarifying the underlying relationships among microstructure, mechanical properties and degradation behaviors, are very meaningful.

Based on our previous studies [15,16], Zn-2Cu-0.8Li (wt%) alloy shows superior strength and ductility as the most significant advantage compared with other Zn alloys, which is due to its unique hierarchical structure of hard micron-sized β-LiZn₄ matrix and soft submicron η-Zn. It also exhibits appropriate degradation properties and good cytocompatibility [[16], [17], [18], [19]]. In this study, Zn-2Cu-0.8Li alloy fine wires for staples with a diameter of 0.22 mm were fabricated by extrusion, multi-pass drawing at room temperature and annealing treatment. The evolutions of microstructure and mechanical properties during drawing were systematically characterized. The degradation behaviors of wires were clarified. The wires were processed into staples, and the in vitro and in vivo studies were performed to evaluate their clinical potential.

2. Materials and methods

2.1. Fine wires and staples processing

The as-cast alloy ingots with the nominal composition of Zn-2Cu-0.8Li (wt%) were prepared by gravity casting. The practical composition was measured as Li: 0.80 wt%, Cu: 1.90 wt% and Zn: Bal. wt% by inductively coupled plasma atomic emission spectroscopy (ICP-AES, iCAP6300, Thermo Fisher, USA). After being homogenized at 350 °C for 24 h, the ingots (a diameter of 60 mm) were extruded into rods (a diameter of 20 mm) at 180 °C. The detailed process procedures are shown in the reference [15].

Subsequently, the as-extruded rods were fabricated into fine wires by extrusion and multi-pass drawing, and then fabricated into staples. The detailed fabrication steps are shown in Fig. 1. The cylinder billets with a diameter of 19 mm and a height of 30 mm were cut from as-extruded rods, and were further extruded into raw wires with a diameter of 1.5 mm at 300 °C with an extrusion ratio of 160:1 and an extrusion speed of 0.6 mm∙s⁻¹. Then the wires with a diameter of 1.5 mm were drawn into fine wires with a diameter of 0.22 mm through 37 passes at room temperature (25 °C) with a drawing speed of 9 cm∙s⁻¹. The drawing strains are calculated by the area reduction ratio, which is 12.5% per pass for the first 20 passes and 6.5% per pass for the remaining 17 passes. The annealing treatment was employed for fine wires at 120 °C, 175 °C and 225 °C for 1 h to lower the internal stress. The fine wires annealed with the optimal parameters were fabricated into U-shape staples by Suzhou Yuezhong Biotechnology Co., Ltd., and then further anastomosed into B-shape staples by the stapler device.

Fig. 1.

The procedures for preparing alloy fine wires and staples.

2.2. Microstructure characterization

The microstructure of wires was characterized using optical microscopy (OM, Zeiss Axio Observer A1, Germany), and scanning electron microscopy (SEM, TESCAN MIRA3, GAIA3, Czechia) equipped with energy dispersive spectrometry (EDS), time of flight-secondary ion mass spectrometry (TOF-SIMS) and electron backscattering diffraction (EBSD). The TOF-SIMS was employed to characterize the Li element distribution, and EBSD was used to analyze the information on grain orientation, grain size, phase identification, and so on. The acceleration voltages for SEM morphology observation, EDS, TOF-SIMS and EBSD analysis are 5 kV, 20 kV, 30 kV and 20 kV, respectively. The wires were cut along the extrusion direction (ED) or drawing direction (DD) and mounted. Then, the samples for OM and SEM observation were ground by 320-grit, 1200-grit, 3000-grit and 7000-grit SiC paper, polished by diamond paste, and finally etched by an acid solution (20 g CrO₃, 1.5 g Na₂SO₄, and 100 mL ultra-pure water). The samples for EBSD observation were ground by SiC paper, and polished by diamond suspension and OPS (oxide polishing suspension), in sequence. The EBSD data were collected with the step sizes of 0.2 μm for 1.5 mm and 0.81 mm wires and 0.03 μm for 0.22 mm wires. They were analyzed by HKL-Channel 5 and AZtecCrystal software. An X-ray diffractometer (XRD, D/MAX 2550, Rigaku Corporation, Japan) with Cu-Kα radiation was employed to analyze the phase constitution of the wires. The scan range of 2θ is 10° to 90°, and the scan rate is 5°∙min⁻¹.

2.3. Mechanical properties

The tensile tests were employed for wires according to ASTM E8/E8M-21. The samples were cut along the ED or DD of wires with a total length of 70 mm and a gauge length of 20 mm. The tests were conducted by a universal material test machine (Zwick/Roell Z20, Germany) at 25 °C with an initial strain rate of 1 × 10⁻³ s⁻¹. At least three parallel specimens were measured for each type of wire.

The tensile tests were employed for the studied alloy and pure Ti staples to evaluate their mechanical properties. The closing B-shape staples with each loop passing through the nylon thread were employed with tension load. The ultimate tensile forces were recorded by a force gauge (ALGOL NK-10, Japan) until the staples were pulled apart. At least five parallel specimens were measured to obtain the average value.

The residual stress distribution of alloy staples during deformation was simulated by finite element analysis using Abaqus software (Abaqus CAE 2022, dynamic explicit). The assembly for 3D simulation consisted of a staple, an anvil groove surface, four rigid plates above the crown for staple push and a rigid surface surrounding the staple for position limitation. They were meshed by 25900 8-nodes linear hexahedral solid elements with reduced integration (C3D8R), 1368 4-nodes bilinear rigid quadrilateral elements (R3D4), 68 R3D4 elements and 5640 R3D4 elements, respectively. The displacement of the staple was completely limited along the X direction and partially limited along the Y direction by the rigid surface to avoid sliding. A displacement loading of 2.45 mm was applied for 4 rigid plates above the crown along the Z direction to press staple down into the anvil groove, thereby completing the deformation.

2.4. In vitro degradation behavior

The in vitro degradation behavior of wires was evaluated by immersion tests according to ASTM G31-21. The tests were employed for as-extruded 20 mm rods, as-extruded 1.5 mm raw wires, as-drawn 0.22 mm wires and the staples. The disc samples for as-extruded rods were cut along the ED with a diameter of 14 mm and a thickness of 3 mm. The samples for 1.5 mm wires were cut along the ED with a length of 27 mm, and those for 0.22 mm wires were cut along the DD with a length of 50 mm. The length of the 0.22 mm wire samples for contrast with staples is 9 mm. These samples were ground, ultrasonically cleaned in alcohol and dried for use. After being weighed, the samples were immersed in conventional simulated body fluid (c-SBF) at 37 ± 0.5 °C. The ratio of sample surface area to solution volume is 1 cm²: 25 mL, and the solution was updated every 48 h. Four parallel specimens were set for every group. After 21 days, the samples were taken out, cleaned by ultrapure water and dried. To measure the weight loss after immersion, the degradation products were removed by an acid solution (200 g CrO₃, 1000 mL H₂O). Then the degradation rate is calculated by:

$$CR=\left(K\times W\right)/\left(A\times T\times D\right)$$ (1)

where $CR$ denotes the degradation rate (μm∙year⁻¹), $K$ is a constant of 8.76 × 10⁷, $W$ denotes the weight loss after immersion (g), $A$ is the immersed sample surface area (cm²), $T$ is the immersion time (h), $D$ is the density of the sample (g∙cm⁻³). The surface and cross-sectional morphologies of samples before and after removing degradation products were observed by SEM (acceleration voltage of 20 kV). The cross-sectional samples were pretreated by gold sputtering. The compositions of degradation products were analyzed by EDS and X-ray photoelectron spectroscopy (XPS, Nexsa, Thermo, US). The XPS data was collected at a depth of 100 nm to avoid surface contamination.

2.5. In vivo property evaluation of staples

The animal experiment involving gastric resection and anastomosis in beagles was approved by the Animal Ethics Committee of Sanitation & Environment Technology Institute, Soochow University, and conducted in strict compliance with relevant guidelines and standards. The staples were prepared and sterilized by Gamma irradiation. Four Beagles (weight 8-15 kg, both sexes) sourced from Shanghai Jiagan Biotechnology Co., Ltd., were prepared. The surgical procedures for staple implantation are as follows. The animal was restrained in a supine position. The abdominal surgical area was shaved, disinfected, and draped with sterile towels. The incision site was exposed. A left upper abdominal incision, approximately 10-15 cm below the xiphoid process, was made. The skin, subcutaneous tissue, muscle, and peritoneum were sequentially incised to expose the abdominal cavity. A gastric area with prominent curvature was resected, and the remaining tissues were anastomosed using the linear cutting stapler. The integrity of the anastomosis was checked to ensure no leakage, bleeding, or other complications. After hemostasis, the abdominal cavity was closed. At postoperative 12 months, X-ray microscope scanning (XRM, radia 520 Versa, Germany) was performed for the gastric tissues anastomosed by staples to investigate the in vivo degradation of the staples. The Avizo software was employed for the three-dimensional reconstruction, and grayscales were utilized to distinguish the corrosion products and residual metals. Simultaneously, the characterizations related to biosafety were also employed, including physiological index detection, metal ion concentration measurement and histological analysis.

2.6. Statistical analysis

The statistical significance of the difference was analyzed by SPSS software. One-way analysis of variance (ANOVA) with Tukey's post hoc test and two-tailed unpaired Student's t-test were performed. The data are expressed as mean value ± standard error (n ≥ 3, independent samples). Differences of ∗p < 0.05 and ∗∗p < 0.01 are considered statistically significant and highly statistically significant, respectively.

3. Results

3.1. Microstructure and mechanical property evolution of wires during drawing

3.1.1. Microstructure evolution

As shown in Fig. 1, the fabricated fine wires show high quality with metallic luster and a length exceeding 100 m. The original as-extruded wires with a diameter of 1.5 mm, intermediate as-drawn wires with a diameter of 0.81 mm, and final as-drawn wires with a diameter of 0.22 mm (abbreviated as 1.5 mm, 0.81 mm, 0.22 mm wires) were chosen to study the microstructure evolution of wires during drawing at room temperature. The results are shown in Fig. 2. After extrusion at 300 °C, the matrix grains of 1.5 mm wires are fully dynamic recrystallized (DRX) and equiaxed with sizes over 10 μm (Fig. 2a). Some needle-like second phases with a length close to grain size and a width of submicron-size (<1 μm) can be observed, most of which distribute along specific directions in the grain interior and terminate at grain boundaries. Fig. 2a1 shows the Li element distribution map by TOF-SIMS. Compared with the needle-like second phase, the matrix is enriched with Li. EDS results show that the element contents of the matrix are Cu: 2.27 ± 0.05 at% (2.21 ± 0.04 wt%) and Zn: 97.73 ± 0.05 at% (97.79 ± 0.05 wt%), and those of the second phase are Cu: 1.76 ± 0.04 at% (1.71 ± 0.03 wt%) and Zn: 98.24 ± 0.04 at% (98.29 ± 0.03 wt%), which are consistent with those reported for as-extruded Zn-2Cu-0.8Li rods [15]. Fig. 2a2 shows the EBSD image of phase distribution. The matrix and needle-like second phase are indexed as β-LiZn₄ and η-Zn, respectively. After multi-pass drawing at room temperature, the microstructure of 0.81 mm wires is shown in Fig. 2b, b1 and b2. The grains are elongated along the DD. Furthermore, the original needle-like second phase in 1.5 mm wires transforms to submicron-sized equiaxed grains of 0.81 mm wires due to DRX, as marked by white arrows. The Li distribution (Fig. 2b1) and phase identification by EBSD (Fig. 2b2) of the β-LiZn₄ matrix and η-Zn of 0.81 mm wires are similar to those of 1.5 mm wires. The microstructure of the final 0.22 mm wires is shown in Fig. 2c. The matrix is further elongated to form banded grains along the DD. Fine DRXed second phase shows a banded distribution along the DD, as marked by white arrows. In addition to elongated grains, some newly-formed fine grains can be observed in the matrix, as marked by red arrows and ellipses in the enlarged image at the top right corner, which indicates DRX occurs in the matrix during the drawing at room temperature. The EDS results of the matrix are Cu: 2.56 ± 0.06 at% (2.49 ± 0.05 wt%) and Zn: 97.44 ± 0.06 at% (97.51 ± 0.05 wt%), and those of the second phase are Cu: 1.85 ± 0.05 at% (1.80 ± 0.04 wt%) and Zn: 98.15 ± 0.05 at% (98.20 ± 0.04 wt%). Fig. 2d shows XRD patterns of 1.5 mm and 0.22 mm wires. The diffraction peaks for both wires are indexed as β-LiZn₄ and η-Zn. The peak intensity of β-LiZn₄ is dramatically stronger, indicating it is the matrix. It is worth noting that for the 1.5 mm and 0.22 mm wires, the β-LiZn₄ matrix and η-Zn both contain a trace amount of Cu. It should exist as the solute in the two phases. Therefore, when the wires are drawn from 1.5 mm to 0.22 mm, the constituent phases of the alloy are always the β-LiZn₄ (solid solution with Cu) matrix and the η-Zn [15]. The β-LiZn₄ micron-sized grains are gradually elongated to form bands along the DD, and DRX occurs in partial bands to form fine grains. The needle-like η-Zn transforms into fine equiaxed grains due to DRX. These fine η-Zn grains are distributed uniformly in the matrix along the DD. Additionally, some dispersive ε-CuZn₄ nanoprecipitates may also form within the two phases [15].

Fig. 2.

Microstructure of Zn-2Cu-0.8Li alloy wires: (a-b) SEM image, (a1-b1) TOF-SIMS maps (Li), and (a2-b2) EBSD phase distribution images of 1.5 mm and 0.81 mm wires, respectively (the black, red and white lines represent β-LiZn₄ HAGB, η-Zn HAGB and phase boundary (PB), respectively); (c) SEM image of 0.22 mm wires; (d) XRD patterns.

The EBSD analysis of three types of wires is shown in Fig. 3. Fig. 3a and b shows the EBSD orientation map and grain orientation spread (GOS) map of as-extruded 1.5 mm wires, respectively. The GOS map exhibits the mean value of misorientations between the orientation of pixels and the average orientation within the grain [20], and reflects the deformation extent of grains [21]. In accordance with the SEM image of Fig. 2a, the grains are fully DRXed and equiaxed after extrusion. Most grains show uniform orientation colors (Fig. 3a), accompanied by low GOS values (Fig. 3b), indicating limited intragranular misorientation and a low degree of deformation within the grains. Fig. 3c and d shows the EBSD images of 0.8 mm wires. After multi-pass drawing at room temperature, the original equiaxed grains are obviously elongated along the DD. Most grains exhibit gradient colors in the EBSD orientation map (Fig. 3c) and higher GOS values (Fig. 3d) compared with Fig. 3a and b, suggesting increased intragranular misorientation and a greater extent of deformation within the grains. Furthermore, in the elongated deformed grains, plenty of newly-formed low-angle grain boundaries (LAGBs), along with the coexistence of LAGBs and high-angle grain boundaries (HAGBs), as well as several fine grains with low GOS value, are observed. As shown in domain 1 of the magnified GOS map on the right, the transformation from LAGBs (marked by gray arrows) to HAGBs (marked by black arrows) is observed in the region circled by white dotted lines. In domain 2, some LAGBs and HAGBs networks are observed, and fine DRXed grains characterized by low GOS values form inside the original deformed grains (marked by red arrows). These features are similar to those of dynamic recovery (DRV) and continuous dynamic recrystallized (CDRX) occurring in as-extruded Zn-2Cu-0.8Li rods during tension [15]. Fig. 3e and f shows the EBSD images of 0.22 mm wires. The grains are further elongated as a whole to form threadlike bands. Although many grains still exhibit gradient colors in the EBSD orientation map (Fig. 3e), the gradient range is narrower and GOS values (Fig. 3f) are lower compared to Fig. 3c and d, indicating a reduction in intragranular misorientation and overall deformation extent. Furthermore, in contrast to the disordered and discontinuous LAGBs and HAGBs in 0.81 mm wires, most newly generated GBs in 0.22 mm wires are organized and interconnected to form subgrains or DRXed grains. The details are shown in enlarged domain 3 on the right, and the subgrains are marked by blue arrows and rectangles. This phenomenon should be due to the enhanced DRV, dislocation rearrangement and CDRX. Accordingly, an increased number of the newly-formed DRXed fine grains with low GOS values are observed compared with those in 0.81 mm wires. In domain 4, the original deformed grains are replaced by newly-formed DRXed grains with low GOS values, and new LAGBs are generated in DRXed grains due to the continued deformation. Therefore, CDRX occurs during drawing, leading to a reduction in intragranular misorientation.

Fig. 3.

EBSD images of Zn-2Cu-0.8Li alloy wires with several diameters during drawing: (a, c, e) EBSD orientation maps and (b, d, f) GOS maps and the corresponding enlarged domains of the wires with diameters of 1.5 mm, 0.81 mm and 0.22 mm, respectively.

The texture and microstructure evolution of wires during drawing are summarized in Fig. 4. The inverse pole figures (IPFs) show the texture of the wires consistently corresponds to the basal fiber texture (Fig. 4a). The crystal directions parallel to ED in 1.5 mm wires mainly distribute between < $\bar{1}$ 2$\bar{1}$ 0> and <01$\bar{1}$ 0>, i.e., <$\bar{1}$ 2$\bar{1}$ 0>−<01$\bar{1}$ 0>//ED. After drawing, the crystal directions parallel to DD shift gradually towards <01$\bar{1}$ 0>, i.e., <01$\bar{1}$ 0>//DD. Simultaneously, the texture intensities multiply from 4.48 multiples of a random distribution (MRD) to 35.57 MRD. Fig. 4b summarizes the microstructure evolution as the drawing deformation extents obtained from EBSD. The newly-formed DRXed grains are counted by the threshold value of GOS <1.5°. When the drawing strain increases from 0% to 70.8% (from 1.5 mm to 0.81 mm wires), the average GOS value and LAGBs fraction dramatically increase from 0.7° and 5.8% to 9.4° and 67.9%, respectively. Meanwhile, the average grain size and volume fraction of newly-formed DRXed grains show a slight decrease and a slight increase (from 11.6 μm to 11.2 μm, and from 0% to 3.5%), respectively. The microstructure is predominantly characterized by deformed grains under this level of drawing strain, with CDRX playing a limited role. When drawing strain further increases to 97.8% (to 0.22 mm wires), the average grain size dramatically decreases from 11.2 μm to 3.0 μm, which is due to the increase in volume fraction of newly-formed fine DRXed grains from 3.5% to 41.2%. On account of the extensive activation of CDRX, the GOS value and LAGBs fraction decrease to 4.3° and 51.2%, respectively. However, they are still higher than those of as-extruded 1.5 mm wires due to the incompletion of DRX.

Fig. 4.

Textures and microstructure evolution of alloy wires during drawing: (a) inverse pole figures along the extrusion/drawing direction; (b) the evolution of average GOS values, LAGBs fractions, newly-formed DRXed grain fractions (GOS value < 1.5°) and average grain sizes as drawing deformation extents.

3.1.2. Mechanical property evolution

The tensile mechanical properties of three types of wires at room temperature with an initial strain rate of 1 × 10⁻³ s⁻¹ are shown in Fig. 5. The TYS, UTS and EL of 1.5 mm wires obtained from engineering stress-strain curves are 376 ± 2 MPa, 423 ± 6 MPa and 88.8 ± 5.0%, respectively. As the drawing pass increases, the strength of wires gradually increases to 471 ± 1 MPa (TYS) and 508 ± 3 MPa (UTS) of 0.81 mm wires, and 596 ± 4 MPa (TYS) and 631 ± 2 MPa (UTS) of 0.22 mm wires. The EL gradually decreases to 62.7 ± 9.7% of 0.81 mm wires and 25.4 ± 2.1% of 0.22 mm wires. Therefore, the strength of wires can be dramatically improved up to over 600 MPa by drawing, while EL is sacrificed.

Fig. 5.

The mechanical properties evolution of Zn-2Cu-0.8Li alloy wires during drawing: (a) representative engineering (solid lines) and true (dotted lines) stress-strain curves; (b) the corresponding tensile yield strength, ultimate tensile strength and fracture elongation.

3.2. Microstructure and mechanical property of wires after annealing treatment

According to the above results, although DRV and CDRX occur in 0.22 mm wires after multi-pass drawing, a large number of LAGBs and deformed grains still exist, resulting in the high internal stress and thermodynamic instability of wires. Furthermore, although the strength of wires is high, the EL is relatively low. Therefore, in order to decrease the internal stress, lower the risk of cracking and stress corrosion, improve thermodynamic stability, and increase the EL of wires, the annealing treatments were employed. They were conducted in an oil bath for 1 h at three different temperatures: 120 °C, 175 °C and 225 °C. The corresponding samples are abbreviated as 0.22 mm-120, 0.22 mm-175 and 0.22 mm-225, respectively. The microstructure and mechanical properties are shown in Fig. 6. The microstructure of 0.22 mm-120 wires (Fig. 6a) is the most similar to that of 0.22 mm wires, which is composed of threadlike deformed grains and submicron-sized DRXed grains, while the deformed grains decrease and fine DRXed grains increase after annealing. As the annealing treatment temperature increases, the grains of 0.22 mm-175 wires are basically DRXed and grow to sizes exceeding 1 μm (Fig. 6b), accompanied by coarsening of the second phase. As the temperature further increases to 225 °C, the equiaxed DRXed grains further grow to about 5 μm (Fig. 6c), and the second phase also further grows. Correspondingly, the tensile engineering and true stress-strain curves and properties are shown in Fig. 6d and e. As the annealing temperature increases, the strength of wires gradually decreases, which is from original 596 ± 4 MPa (TYS) and 631 ± 2 MPa (UTS) to 492 ± 3 MPa and 537 ± 1 MPa of 0.22 mm-120 wires, 460 ± 5 MPa and 495 ± 4 MPa of 0.22 mm-175 wires, and 387 ± 1 MPa and 413 ± 2 MPa of 0.22 mm-225 wires. The EL firstly dramatically increases from 25.4 ± 2.1% to 44.2 ± 4.2%, whereafter, although it further increases, the increment is slight (49.2 ± 4.7% and 50.4 ± 6.2%). Therefore, the mechanical properties of 0.22 mm-120 wires are optimal. Based on the mechanical properties, the annealing technique of 120 °C × 1 h is chosen for further studies.

Fig. 6.

Microstructure and mechanical properties of 0.22 mm wires after annealing treatment: microstructure after annealing for 1h at (a) 120 °C; (b) 175 °C; (c) 225 °C; (d) tensile engineering (solid lines) and true (dotted lines) stress-strain curves; (e) tensile yield strength, ultimate tensile strength and fracture elongation.

EBSD was employed to further study the microstructure of 0.22 mm-120 wires, as shown in Fig. 7. After annealing at a relatively low temperature, more DRXed fine grains with low GOS values form, in accordance with Fig. 6a. Compared with 0.22 mm wires, the grains exhibit more uniform orientation colors (Fig. 7a) and lower GOS values (Fig. 7b). The texture of wires after annealing is still <01 $\bar{1}$ 0>//DD with no significant change in intensity. Fig. 7c summarizes the microstructure information of 0.22 mm-120 wires. Compared with the wires before annealing, the average GOS value and LAGBs fraction decrease from 4.3° to 3.8° and from 51.2% to 45.2%, respectively. The volume fraction of newly-formed recrystallized grains increases from 41.2% to 46.1%. The average grain size slightly increases from 3.0 to 3.3 μm.

Fig. 7.

EBSD images of wires after annealing treatment at 120 °C for 1 h: (a) EBSD orientation map and the corresponding inverse pole figure along the drawing direction; (b) GOS map; (c) average GOS value, LAGBs fraction, newly-formed recrystallized grain fraction (GOS value < 1.5°) and average grain size compared with those of wires before annealing treatment.

3.3. Degradation behaviors of wires

The degradation behaviors of 0.22 mm-120 wires were studied by being immersed in c-SBF at 37 °C for 21 days, with as-extruded 1.5 mm wires and as-extruded 20 mm rods as controls (Fig. 8). As shown in Fig. 8a, the degradation rate of 0.22 mm-120 wires is 327 ± 20 μm∙year⁻¹, which is signally higher than 236 ± 11 μm∙year⁻¹ of 1.5 mm wires and 112 ± 11 μm∙year⁻¹ of 20 mm rods. The accumulated Zn and Li ion concentrations released during immersion were measured, as shown in Fig. 8b and c. Both ion concentrations for 0.22 mm-120 wires are higher than those for 1.5 mm wires and 20 mm rods, corresponding with the higher degradation rate. These differences about ion concentration and degradation rate are possibly related to the microstructure change induced by drawing, which will be discussed in detail in Section 4.3. Furthermore, as the immersion time increases, the curves gradually become flatter, and the increment of ion concentrations is gradually smaller, especially for 0.22 mm-120 and 1.5 mm wires. This may be related to the decrease in degradation rate with the prolongation of immersion time. The accumulated ion release concentrations of three samples after 21 days are Zn²⁺: 5.32 ± 0.01 ppm, 17.26 ± 1.0 ppm and 20.90 ± 0.10 ppm and Li⁺: 1.80 ± 0.05 ppm, 4.96 ± 0.17 ppm and 5.87 ± 0.93 ppm, respectively. The Cu²⁺ was non-detected, possibly due to its too-low concentration below detection limits. This phenomenon is likely attributed to the fact that the Cu element is electrochemically inert compared with Zn and Li, resulting in a significantly slower dissolution rate and too low contents of Cu²⁺, which will be discussed in detail in Section 4.3. Furthermore, once Cu degrade, the released Cu²⁺ tend to preferentially transform into insoluble Cu-containing compounds such as Cu(OH)₂, due to their low solubility (K_sp 2.2 × 10⁻²⁰ (‌mol∙L⁻¹)³ for Cu(OH)₂, in contrast to K_sp 1.2 × 10⁻¹⁷ (‌mol∙L⁻¹)³ for Zn(OH)₂ [22] and the good solubility for LiOH). Notably, the Li⁺ release concentration is markedly higher than that of Zn²⁺ in all groups, as shown in the Li⁺/Zn²⁺ atomic ratios exceeding 1 (Fig. 8d), which is possibly induced by the Li dealloying (discussed in detail in Section 4.3). An inverse correlation between the Li⁺/Zn²⁺ ratio and degradation rate is observed, which suggests that slower degradation may facilitate the preferential release of Li⁺ over Zn²⁺, potentially due to their differing electrochemical activities. Actually, the slower degradation would benefit the Li dealloying, which is usually controlled by the diffusion needing adequate time [23,24].

Fig. 8.

The degradation properties of 0.22 mm wires compared with 1.5 mm wires and 20 mm rods: (a) degradation rates after being immersed in c-SBF for 21 days; The accumulated ion concentration curves of (b) Zn²⁺, (c) Li⁺ and (d) atomic ratio of Li⁺/Zn²⁺.

The surface degradation morphologies of the three samples after being immersed for 21 days are shown in Fig. 9. As shown in Fig. 9a–c, a degradation product layer accompanied by some dispersive white clusters can be observed on the surface. The amounts of degradation products for three samples can be arranged as: 0.22 mm-120 > 1.5 mm > 20 mm, which is consistent with degradation rates and ion release concentrations. The EDS results at the bottom left of images show that the degradation products mainly contain O, P, Ca, Zn, C and trace Cu. Fig. 9d–f shows the surface morphology after removing degradation products. Tiny corrosion pits can be observed on the surface of the three samples. The surface of 20 mm rods is flat with a few pits, while that of 0.22 mm-120 wires is loose with plentiful pits, and the surface of 1.5 mm wires lies between them. It is worth noting that these pits exhibit very small size and depth and uniform distribution, indicating a homogeneous degradation behavior. The regions marked by white arrows in Fig. 9e and f correspond to the needle-like η-Zn of 1.5 mm wires (Fig. 2a) and fine η-Zn along the DD of 0.22 mm-120 wires (Fig. 6a). These regions of η-Zn preferentially degrade compared with the matrix due to the micro-galvanic corrosion between the two phases, which is consistent with the lower potential of η-Zn relative to the β-LiZn₄ (solid solution with Cu) as reported in the previous studies [19]. The β-LiZn₄ (solid solution with Cu) matrix should act as the cathode, and the η-Zn phase should act as the anode. Additionally, the Cu distribution on the corroded surface varies with degradation severity. In intact regions, the Cu content remains stable at the original matrix level (2.1–2.5 at%, 2.0–2.4 wt%, marked by black dotted lines and points), whereas it becomes significantly elevated (up to 7.3 at%, 7.1 wt%, marked by blue points) in severely corroded areas, as seen to be most evident in the 0.22 mm-120 wires.

Fig. 9.

SEM images showing degradation morphology of surface and EDS results of the corresponding points: (a-c) containing degradation products; (d-f) removing products of 20 mm rods, 1.5 mm and 0.22 mm-120 wires, respectively (the black digits and points mark the Cu content remaining at the original matrix level in intact regions, and the blue digits and points mark the elevated Cu content in severely corroded regions).

The cross-section containing degradation products of 1.5 mm and 0.22 mm-120 wires is studied, and the results are shown in Fig. 10. The thickness of degradation products of 0.22 mm-120 wires is generally larger than that of 1.5 mm wires (Fig. 10a and d), corresponding with a higher degradation rate. EDS maps show that O, P and Ca are deposited in degradation products, and TOF-SIMS results show that the degradation products contain Li, though at a lower content than in the alloy matrix (Fig. 10b and e). The element distributions of the degradation products of the two wires are similar. The further local enlarged images are shown in Fig. 10c and f. The needle-like η-Zn and fine η-Zn along the DD marked by white arrows preferentially degrade to form degradation products, and show different contrast from the matrix, corresponding with Fig. 9e and f. Moreover, the Cu contents of the matrix near degradation products by EDS (5.6-7.0 at% and 4.9-7.4 at%) are significantly higher than those of the non-degraded matrix far away from degradation products (2.5-2.6 at% and 2.2-2.8 at%), which is consistent with Fig. 9e and f. Therefore, as degradation progresses, the Cu is enriched in the alloy matrix near degradation products. This Cu enrichment phenomenon is possibly correlated with the preferential dissolution of Li and Zn due to the electrochemical activity difference, which will be discussed in detail in Section 4.3.

Fig. 10.

Cross-sectional degradation morphology: (a, d) macroscopic SEM morphology; (b, e) local enlarged SEM images showing degradation products and EDS (O, P, Ca, Zn, Cu and C) and TOF-SIMS (Li) maps; (c, f) further enlarged SEM images showing degradation behaviors of 1.5 mm and 0.22 mm-120 wires, respectively (the black digits and points mark the EDS Cu content remaining at the original matrix level in intact regions, and the blue digits and points mark the elevated EDS Cu content in severely corroded regions).

XPS analysis was employed to further study the components of degradation products, as shown in Fig. 11. The information at a depth of 100 nm was collected by XPS depth-profiling to avoid surface contamination. Fig. 11a shows the element contents of degradation products and the corresponding XPS spectra with characteristic peaks clearly labeled. The types of elements for the two samples are the same, corresponding with Fig. 9, Fig. 10. The O contents are the highest, and the Li contents are 5.4 at% (1.5 mm) and 2.7 at% (0.22 mm-120), respectively. The high-resolution XPS scans were employed for Zn 2p_3/2, O 1s, C 1s, P 2p, Li 1s and Ca 2p peak regions. The spectra were fitted in reference to the binding energies of possible compounds [[25], [26], [27]] in Fig. 11d. The results show that degradation products for both samples mainly contain Li₂CO₃, LiOH, Zn₅(OH)₆(CO₃)₂, ZnO, Zn(OH)₂, Zn₃(PO₄)₂ and Ca/P compounds. Furthermore, the relative order of the fitted XPS peak areas for the diverse compounds in both samples is identical, indicating the corresponding order in their relative content is identical.

Fig. 11.

XPS depth profile analysis of degradation products (100 nm depth): (a) element contents; XPS spectra and high-resolution spectra of Zn 2p_3/2, O 1s, C 1s, P 2p, Li 1s and Ca 2p regions for (b) 1.5 mm and (c) 0.22 mm-120 samples, respectively; (d) binding energies of possible compounds [[25], [26], [27]].

3.4. The properties of staples

The 0.22 mm-120 wires were further fabricated into U-shape staples, and then were formed into B-shape by the stapler (Fig. 1). In vitro and in vivo studies of the staples were conducted to explore their application feasibility.

3.4.1. The properties of staples in vitro

The mechanical properties of B-shape staples made from the studied Zn alloy and pure Ti were measured. The ultimate tensile force of Zn alloy staples is 1.86 ± 0.10 N, which is approximately 94% that of the pure Ti staples (1.99 ± 0.04 N). The developed Zn-2Cu-0.8Li (wt%) alloy staples exhibit favorable mechanical properties. Furthermore, considering that the residual stress induced during fabrication and B-shape formation could influence the degradation behaviors of staples [7,28], the degradation behavior of B-shape staples was studied by being immersed in c-SBF for 21 days, in comparison with the wires (Fig. 12). Fig. 12a shows the morphologies of B-shape staples. The first-row images present the overall and locally magnified views of the staples before immersion, which demonstrate structural integrity without deformation-induced cracks. The second-row images show the surfaces after immersion for 21 days, where abundant degradation products are observed on the staple feet and leg arc regions, as marked by red and orange rectangles and the corresponding enlarged images, respectively. In contrast, fewer products are found in other regions (blue rectangles). After removing degradation products, as shown in the third row, the staples retain good structural integrity with no fracture or disintegration, meeting the clinical requirements for staples used in soft tissue healing that necessitate maintaining integrity for at least 3 weeks [5]. The degree of local degradation is consistent with the amounts of degradation products. More severe degradation is observed on the feet and the leg arc parts of staples, which exhibit looser, more porous, and rougher surfaces (red and orange rectangles). Other regions (blue rectangles) show relatively milder degradation with smoother surfaces. Although regional variations in degradation exist, the differences are not significant. Fig. 12b shows the morphologies of wires before and after removing degradation products, indicating a generally uniform degradation. Fig. 12c shows that the degradation rate of staples is slightly higher than that of the wires, which is possibly related to the residual stress concentration on local regions of the staples after anastomosis deformation. The residual stress could usually lead to degradation acceleration by causing the decrease of corrosion potential, rupture of the surface layer and activation of more dislocations and defects [29,30]. The feet and leg arc parts of staples with more severe degradation correspond to the regions concentrated with residual stress of the staples by simulation in Fig. 12d and references [7,28]. The difference is not significant, which is possibly attributed to the weak stress concentration degree of staples or the low stress corrosion sensitivity of the alloy [30].

Fig. 12.

Degradation behaviors of staples after being immersed in c-SBF for 21 days: (a) original morphology, morphology with corrosion products and after being removed corrosion products of staples; (b) wires with and after being removed corrosion products; (c) comparison of degradation rates between staples and wires after 21 days; (d) residual stress diagrams of original U-shape and deformed B-shape staples by finite element analysis.

3.4.2. The properties of staples in vivo

According to the above results, the Zn²⁺ release concentration of fine wires and staples are higher than that of rods. In our previous study [17], the Zn²⁺ release concentration in extracts of fine wires was lower than the adverse effect threshold [31], and the endothelial cytocompatibility of the Zn-2Cu-0.8Li alloy fine wires was good, meeting the biomedical application requirements [32]. Furthermore, the Zn-2Cu-0.8Li alloy staples were utilized for gastric anastomosis in beagle models to preliminarily evaluate their in vivo performance. Fig. 13a shows the surgical stapler components loaded with the studied staples. In Fig. 13b, gastric resection and anastomosis are performed, demonstrating successful tissue penetration by the staples and satisfactory anastomosis without significant leakage or hemorrhage. The resected anastomotic tissue is shown in Fig. 13c, confirming a neat and intact stoma. The experimental animals exhibited normal clinical parameters postoperatively, including appearance, behavioral patterns, food consumption, glandular secretion levels, and wound healing progression, with no significant infection or complications. The body weight increased normally from 10-11 kg preoperatively to 13-14 kg postoperatively. At postoperative 12 months, the anastomotic gastric tissues healed well. The morphologies of the staples with and without degradation products analyzed by XRM are shown in Fig. 13d and e, respectively. Some regions on the feet and the leg arc part of staples have completely degraded, which is consistent with the results in vitro (Fig. 12). The volume of the residual metal is 0.184 mm³, which is roughly 32% of the original volume (0.569 mm³) of the staple. According to the reduced volume, the average degradation rate of the staple is calculated as 60 μm∙year⁻¹. The volume of the staple containing degradation products (0.486 mm³) is smaller than that of the original staple, which indicates the absorption of the degradation products in vivo. It can be noted that the degradation rate in vivo is significantly lower than that in vitro in Fig. 12c. This is possibly attributed to the differences in environment and measurement period [33]. In vivo, the Cl⁻ concentrations in blood plasma (103 mM) are lower than those in SBF (147.8 mM) [34], and proteins and amino acids exist to possibly form a protective layer on the surface of alloys in some cases [35]. Furthermore, after implantation in vivo, some host tissue responses will occur, resulting in an initial pH decrease and degradation acceleration to form a protective corrosion layer [33]. The encapsulation of the implant by surrounding tissues may also suppress micro-galvanic corrosion [36]. They all benefit long-term lower degradation rate in vivo. Additionally, the longer measurement period in vivo (12 months) than in vitro (21 days) is also likely one of the reasons.

Fig. 13.

The implantation of staples for gastric anastomosis in beagles: (a) surgical stapler components; (b) gastric tissue anastomosis; (c) the cut anastomosed tissue; XRM analysis of staple degradation (d) containing and (e) after removing degradation products at postoperative 12 months.

Moreover, the staples also show good biosafety, as shown in Fig. S1–2 and Tables S1–2. At postoperative 12 months, the overall structure of the gastric mucosa appears normal, with no signs of erosion, hemorrhage, necrosis, or fibrotic changes (Fig. S1). Only slight hyperplastic muscle fibers with a few lymphocytes, macrophages and new blood vessels exist in the local submucosal layer. No necrosis or calcification is observed in the surrounding tissues. Furthermore, at postoperative 12 months, the Zn²⁺, Cu²⁺, and Li⁺ concentrations in the blood specimens of animals show no significant differences compared to those at preoperative 1 day (Fig. S2). The blood biochemical indicators basically remain normal (Tables S1–2). In general, the biosafety of the staples should be good. In terms of the biosafety of the added element of the alloy staples, Zn and Cu are the essential trace elements of the human body. Li also has good biosafety, which is usually used to prevent and treat some diseases like brain injuries, strokes, Alzheimer's disease, mental disorders, etc. [37,38]. However, the long-term excessive intake of Li will lead to toxicity of organs and various disorders such as cardiac arrhythmia, convulsion, hypothyroidism, and sclerotic glomeruli [38,39]. Actually, roughly fifty staples are implanted into every animal, and the total weight is 0.19 g (3.87 mg per staple). Assuming that the complete degradation period of staples is 18 months based on the XRM results, the average daily release of Zn, Cu, and Li is estimated to be 0.34 mg, 7 μg, and 2.8 μg, respectively. These values are far below the recommended dietary allowance (RDA) for adults (11 mg∙day⁻¹ [40]-3 mg∙day⁻¹ [41] and 1.0 mg∙day⁻¹ [37], respectively), confirming good biosafety. This in vivo experiment is only a preliminary evaluation of the application feasibility of the alloy staples. It has some limitations, including insufficient sample sizes and a lack of a control group. More detailed and in-depth research and preclinical trials on staple performance in vivo will be further carried out for the large animal with large sample sizes in the future.

4. Discussion

4.1. Microstructure evolution of wires

The microstructure of wires mainly consists of the β-LiZn₄ matrix and η-Zn phase. After multi-pass drawing at room temperature, the β-LiZn₄ is gradually elongated along the DD. Simultaneously, DRV and CDRX occur in some deformed β-LiZn₄, generating LAGBs, HAGBs and new fine grains (Fig. 2, Fig. 3), thereby resulting in the increase of LAGBs fractions and decrease of average grain size. The basal fiber texture intensity significantly increases (Fig. 4). The η-Zn transforms into fine equiaxed grains due to DRX, and distributes as bands along the DD in the matrix.

The above microstructure evolution of wires during the drawing is mainly related to the alloy properties and drawing conditions. Our previous study has found that [15] the β-LiZn₄ in Zn-2Cu-0.8Li (wt%) alloy exhibits favorable deformability during the uniaxial tensile deformation at room temperature, characterized by activation of numerous dislocations and occurrence of CDRX. For the drawing process, although the stress state of the alloy (biaxial compression and uniaxial tension) is more complex than that of uniaxial tension, the microstructure evolution is similar. On the one hand, the good deformability of the β-LiZn₄ phase facilitates the generation and proliferation of dislocations during drawing, thereby contributing to the grains’ elongation and affording driving force for DRV and CDRX [15]. On the other hand, the appropriate drawing deformation conditions also benefit the occurrence of DRV and CDRX. Firstly, the large enough deformation strains accumulated by 37 drawing passes afford sufficient driving force for DRV and CDRX [42]. Secondly, the drawing temperature is room temperature (298 K), which is approximately 0.4 times the maximum melting point of β-LiZn₄ (754 K) and belongs to intermediate temperature [[43], [44], [45]]. This temperature can ensure the multiplication of dislocations during the deformation to afford enough driving force, while it can also provide appropriate thermal activation energy for DRV and CDRX [42]. Thirdly, the drawing rate is relatively low (9 cm∙s⁻¹), which affords sufficient time for the transformation from dislocations to LAGBs and then to HAGBs during CDRX [46,47]. In aspects of textures, the β-LiZn₄ exhibits intensified basal fiber texture during drawing primarily because its dominant deformation mode is basal <a> slip [15]. Although the occurred CDRX during drawing has the effect of weakening basal fiber texture, the effect is limited, which will not cause the dramatic orientation change due to the gradual and continuous transformation process from LAGBs to HAGBs [48,49]. In terms of the η-Zn phase, it transforms into fine grain bands through deformation and DRX, possibly due to the good deformability at room temperature [[50], [51], [52]] and a relatively low melting point (693 K).

After drawing, annealing treatment was employed. The microstructure evolution mainly involves static recrystallization (SRX). After annealing at 120 °C, the grains of wires are non-fully recrystallized. Meanwhile, the LAGBs fractions and deformation extent decrease, and SRX fractions increase (Fig. 6, Fig. 7), which is possibly induced by the partial coalescence and migration of subgrains in 0.22 mm wires. As the annealing temperature increases to 175 °C and 225 °C, the grains are basically recrystallized, and the grain sizes significantly increase (Fig. 6), which is due to the increased rates of SRX and grain growth [53,54]. Therefore, for the developed wires, relatively low annealing temperatures and thermal activation energies can achieve SRX and grain growth. This should be attributed to the sufficient deformation storage energy and driving force afforded by amounts of LAGBs and dislocations forming during multi-pass drawing, and the relatively low maximum melting point of β-LiZn₄ (754 K).

4.2. Mechanical property evolution of wires

The data of mechanical properties, degradation rates and biocompatibility of representative Zn alloys and Zn-2Cu-0.8Li alloy are shown in Table 1. Compared with other Zn alloys, the most significant advantage and novelty of Zn-2Cu-0.8Li alloy is its high strength and ductility. In this study, outstanding mechanical properties were also achieved for Zn-2Cu-0.8Li alloy wires in virtue of hot extrusion, multi-pass drawing and annealing treatment. As shown in Fig. 14a, compared with the reported Zn alloy fine wires in Refs. [[6], [7], [8], [9], [10], [11], [12], [13],[55], [56], [57], [58]], Zn-2Cu-0.8Li alloy fine wires developed in this work meet the TYS requirement of staple materials (purple dotted line), and exhibit an obvious advantage in strength and elongation balance (gray dotted line), as marked by purple pentagrams. The related mechanical property evolutions are analyzed as follows.

Table 1.

The mechanical properties, degradation rates and biocompatibility of Zn-2Cu-0.8Li alloy compared with other Zn alloys, the alloys are all as-extruded except for those with labels.

Alloy system	Alloy	TYS (MPa)	UTS (MPa)	EL (%)	Degradation rate (μm∙year−1)	Biocompatibility
Zn–Mg based [59,60]	Zn–0.5 Mg	159	297	13.0	81 (Hank's, 14 days)	Mg: essential element RDA 240-420 mg [61] Good biocompatibility
	Zn–1Mg	180	340	6.0	83 (Hank's, 14 days)
	Zn-1Mg-0.1Mn As-rolled	195	299	26.1	110 (Hank's, 30 days)
Zn–Li based [62,63]	Zn-0.8Li As-rolled	262	401	80.8	8 (c-SBF, 30 days)	Li: good biosafety, prevent and treat some diseases [37,38] RDA 1.0 mg [37] Good biocompatibility
	Zn-0.8Li-0.4 Mg	433	647	3.0	−
	Zn-0.8Li-0.4Mn	333.3	480	73	−
Zn–Mn based [64,65]	Zn-0.2Mn	132	220	48	−	Mn: essential element RDA 4 mg [41] Good biocompatibility
	Zn-0.8Mn	127	219	64.2	101 (Electrochemical)
	Zn-0.8Mn-0.4Cu	191	308	38.9	133 (Electrochemical)
Zn–Ag based [52,66]	Zn-7.0Ag	236	287	32	84 (Hanks', 14 days)	Ag: good antibacterial performance RDA 0.4–27 μg [67] Good biocompatibility
	Zn-4Ag-0.6Mn	275	302	35	−
Zn-Cu based [15,68,69]	Zn-2Cu	178	222	57.4	75 (c-SBF, 21 days)	Cu: essential element good antibacterial performance RDA 1-3 mg [41] Good biocompatibility
	Zn-2.2Cu-1.0Mn	267	348	29.9	49 (c-SBF, 20 days)
	Zn–2Cu–0.07Li	362	322	28.0	90 (c-SBF, 21 days)
	Zn–2Cu–0.8Li	426	472	63.7	112 (c-SBF, 21 days)
	Zn–2Cu–0.8Li As-drawn wire	596	631	25.4	−
	Zn–2Cu–0.8Li As-annealed wire	492	537	44.2	327 (c-SBF, 21 days)

Fig. 14.

Mechanical property analysis of alloy wires: (a) TYS-EL values comparing with reported Zn alloy fine wires in Refs. [[6], [7], [8], [9], [10], [11], [12], [13],[55], [56], [57], [58]] (purple dotted line represents the TYS requirements of staple materials, gray dotted line represents product of TYS and EL, the colors of data point represent the alloy types and the shapes represent the diameters of wires); (b) estimation of TYS contributions.

After multi-pass drawing from 1.5 mm to 0.81 mm and then to 0.22 mm wires, the strength of the wires dramatically improves. The correlated strengthening mechanisms for this process are analyzed. It is reported that the effect of the η-Zn phase on the strength of Zn-2Cu-0.8Li alloy is weak due to its low volume fraction [15]. Therefore, the analysis primarily focuses on the contribution of the β-LiZn₄ matrix. Combined with the microstructure evolution of wires during drawing, the tensile yield strength ${\sigma }_y$ can be estimated by the lattice friction stress ${\sigma }_0$, precipitation strengthening ${\sigma }_p$, LAGBs strengthening ${\sigma }_{LAGB}$, HAGBs strengthening ${\sigma }_{HAGB}$ and texture strengthening ${\sigma }_{Tex}$:

$${\sigma }_y={\sigma }_0+{\sigma }_p+{\sigma }_{LAGB}+{\sigma }_{HAGB}+{\sigma }_{Tex}$$ (2)

where ${\sigma }_0$ is 166 MPa [15]. Additionally, nanoprecipitations have been observed in as-extruded Zn-2Cu-0.8Li rods [15], whereas they are not characterized in detail in this study due to the great difficulty of preparing Zn alloy fine wire TEM samples. Assuming the nanoprecipitations in wires are the same as those in as-extruded rods, the value of ${\sigma }_p$ is adopted as 104 MPa [15]. Therefore, for the three types of wires during drawing, the value of ${\sigma }_0+{\sigma }_p$ is a constant of 270 MPa. The other strengthening contributions are discussed as follows.

• (1)LAGBs strengthening

As the drawing, the LAGBs fractions of wires dramatically increase. Abundant LAGBs can improve the yield strength of alloys by hindering the motion of dislocations [70]. The LAGBs mainly consist of dislocations, herein, the corresponding strengthening effect is estimated by the dislocation strengthening mechanism [[71], [72], [73]]:

$${\sigma }_{LAGB}=M\alpha Gb\sqrt{{\rho }_0+{\rho }_{LAGB}}$$ (3)

where $M$ is the Taylor factor adopted by 6.00 [15], $\alpha$ is a constant of 0.20, $G$ is the shear modulus adopted by 36.1 GPa [15], $b$ represents the magnitude of the Burgers vector adopted by the value of primary slip system (basal <a> slips), i.e., 0.271 nm [15], ${\rho }_0$ represents the dislocation density of well-annealed materials adopted by 10¹² m⁻², ${\rho }_{LAGB}$ is the dislocation density stored in LAGBs:

$${\rho }_{LAGB}=\frac{3\left(1-f\right){\bar{\theta }}_{LAGB}}{bL}$$ (4)

where $f$ is the fraction of HAGBs, which can be obtained by subtracting the fraction of LAGBs (Fig. 4) from the total value of 1, and the corresponding values for 1.5 mm, 0.81 mm and 0.22 mm wires are 0.942, 0.321 and 0.488, respectively, ${\bar{\theta }}_{LAGB}$ represents the average misorientation of LAGBs, which is calculated as the sum of the products of misorientation angles within LAGBs range (2°-10°) and the corresponding fractions obtained from EBSD data (Fig. 3, Fig. 4), divided by the total fraction of LAGBs, and the value is 4.34°, 4.09° and 4.12° for the three types of wires, respectively, and is 0.16, 0.11 and 0.12 after being converted to radians, $L$ is the average boundary spacing considering both LAGBs and HAGBs, which is measured from EBSD data (Fig. 3) using the threshold degree of 2°, and is 10.65 μm, 6.75 μm and 2.59 μm, respectively. Substituting the above parameters into equation (4), the values of ${\rho }_{LAGB}$ for the three types of wires are calculated as 0.45 × 10¹³ m⁻², 7.95 × 10¹³ m⁻² and 15.76 × 10¹³ m⁻². It is obvious that the dislocation density stored in LAGBs of wires significantly increases as drawing. Substituting ${\rho }_{LAGB}$ and other parameters into equation (3), the LAGBs strengthening contribution ${\sigma }_{LAGB}$ is estimated as 28 MPa, 105 MPa and 148 MPa, respectively.

• (2)HAGBs strengthening

The strengthening contribution from HAGBs of wires follows the Hall-Petch relationship, and can be estimated by the following equation:

$${\sigma }_{HAGB}=k_y{\left(\frac{L}{f}\right)}^{-1/2}$$ (5)

where $k_y$ is the hall-Petch coefficient, which is approximately adopted by the value of pure Zn [74] (220 MPa∙μm^1/2). The HAGBs strengthening contribution of the three types of wires is estimated as 65 MPa, 48 MPa and 95 MPa, respectively.

• (3)Texture strengthening

During room-temperature drawing, the basal fiber texture intensity of wires significantly increases (Fig. 4). The texture closely correlates with the Schmid factor. The increase of basal texture intensity of alloys will result in the decrease of the Schmid factor of basal <a> slips when loading with the uniaxial stress along the wrought direction, thereby improving activation difficulty of slip systems and increasing tensile yield strength. The texture strengthening contributions can be estimated as [75,76]:

$${\sigma }_{Tex}=\left(\frac{m_R}{m_{Tex}}-1\right){\sigma }_{HAGB}$$ (6)

where $m_R$ is the average Schmid factor of polycrystal with random orientation, on account of the unknown value of the β-LiZn₄, $m_R$ is adopted by the average Schmid factor of basal <a> slips of 1.5 mm wires (0.14) as a reference for the wires after drawing, $m_{Tex}$ is the average Schmid factor of the samples with textures, and the values of basal <a> slips for 0.81 mm and 0.22 mm wires are 0.11 and 0.06, respectively. The corresponding texture strengthening contributions of 0.81 mm and 0.22 mm wires based on 1.5 mm wires are estimated to be 13 MPa and 127 MPa, respectively.

Substituting the above strengthening contributions into equation (2), the tensile yield strength ${\sigma }_y$ of 1.5 mm, 0.81 mm and 0.22 mm wires is estimated as 363 MPa, 436 MPa and 641 MPa, respectively. The estimated strengthening contributions and tensile yield strength measured by experiments are shown in Fig. 14b. It can be observed that the estimated values accord well with experimental results. When the wires are drawn from a diameter of 1.5 mm to 0.81 mm, the LAGBs strengthening contribution is dominant, which is consistent with the significantly increased LAGBs fraction by about 60% and the nearly 18 times higher dislocation density stored in LAGBs. The texture strengthening contribution ranks second. When the wires are further drawn from a diameter of 0.81 mm to 0.22 mm, the texture strengthening contribution is dominant, primarily due to the nearly threefold higher texture intensity and increased HAGBs strengthening according to equation (6). The HAGBs strengthening contribution ranks second and is also notable, which is consistent with the dramatic reduction in average grain size to nearly one-quarter due to CDRX. The LAGBs strengthening contribution is relatively low, which is correlated with the decrease of the LAGBs fraction induced by CDRX. After multi-pass drawing at room temperature, the yield tensile strength of wires dramatically increases by about 200 MPa, and is outstanding.

On the other hand, although the EL gradually decreases as the drawing proceeds, it ultimately maintains 25.4% for 0.22 mm wires. Meanwhile, the wires maintain good ductility and workability without fracture during multi-pass drawing at room temperature, even under a cumulative deformation strain of up to 97.8%. This benefits highly efficient, low-energy wire fabrication and high wire surface quality without the need for intermediate annealing. The above property characteristics are mainly related to the following factors. Firstly, DRV and CDRX occur during drawing, which helps with dislocation elimination and fine recrystallized grain formation, thereby benefiting ductility. Simultaneously, the plenty of LAGBs generated by DRV could afford nucleation sites for mobile dislocations, contributing to relaxation of the internal strain induced by dislocation pile-up during drawing [70,77]. Secondly, the possibly activated non-basal slips in the β-LiZn₄ matrix during the deformation of Zn-2Cu-0.8Li alloy at room temperature [15] also play important roles, which is due to the accommodation of the strain along the c axis of <c+a> dislocations [42,78,79], the benefit to CDRX occurrence by forming three-dimensional dislocation networks [78], and the promotion of the basal texture randomization.

Secondly, the softer fine η-Zn in wires has good deformability and facilitates the weakening of stress concentration at GBs of harder β-LiZn₄ and phase boundaries, benefiting ductility [15].

After annealing treatment, the strength of wires gradually decreases as the annealing temperature increases, while the EL initially dramatically increases and subsequently slightly increases. For 0.22 mm-120 wires, the decrease in strength should be attributed to the decreased deformation extent and LAGBs, as well as the slightly increased average grain size. Simultaneously, the reduction of defects like LAGBs contributes to the significant increase in EL. However, as the annealing temperature further increases to 175 °C and 225 °C, the significantly increased grain sizes become the main factors adversely influencing both the strength and ductility.

4.3. Degradation mechanisms of wires

In terms of Zn-2Cu-0.8Li alloy, the degradation modes of 20 mm rods, 1.5 mm wires and 0.22 mm-120 wires are all uniform without large or deep corrosion pits (Fig. 9). This is closely correlated with the microstructure. Firstly, the distribution of the β-LiZn₄ matrix grains with various sizes is uniform. Secondly, the micron-sized fine η-Zn uniformly distributes in the β-LiZn₄ matrix, thereby the micro-galvanic couples formed between the two phases are dispersive. Thirdly, the electrochemical activity difference between the η-Zn and β-LiZn₄ is small [19]. The above factors contribute to the low susceptibility of local galvanic corrosion [80] and uniform degradation modes. Furthermore, compared with as-extruded 1.5 mm wires, the degradation rates and metallic ion release concentrations of 0.22 mm-120 wires are both dramatically higher (Fig. 8). This should be mainly attributed to the improvement of GBs density by drawing. GBs can be categorized as LAGBs and HAGBs. They are usually corroded preferentially due to their higher free energy induced by the concentration of defects and disarrayed atoms [81,82]. Compared with as-extruded 1.5 mm wires, the grains of 0.22 mm-120 wires are significantly refined due to CDRX during drawing, improving the HAGBs density, thereby increasing the degradation rates [82,83]. Besides, through drawing, the fraction and density of LAGBs for 0.22 mm-120 wires are significantly higher than those for as-extruded 1.5 mm wires (Fig. 3, Fig. 4, Fig. 7). The LAGBs are mainly composed of dislocations, and the improved dislocation density is usually thought to contribute to the increase in degradation rates [[84], [85], [86]]. In addition to GBs, other point and line defects like dislocations also generate and proliferate in wires during drawing [[87], [88], [89]], accelerating degradation. Additionally, the galvanic corrosion between the β-LiZn₄ and η-Zn may also influence degradation rates. Compared with the needle-like η-Zn in 1.5 mm wires, the η-Zn (anode) in 0.22 mm-120 wires is dramatically refined. This results in the increased amounts of η-Zn to form more dispersive micro-galvanic couples with β-LiZn₄ [90,91]. Moreover, the refinement of η-Zn (anode) could possibly increase the galvanic corrosion current density, which is usually thought to have a direct proportion with cathodic area/anodic area [92,93]. They all accelerate the degradation. It has been found by several studies that the degradation rates of wires are also possibly related to their diameters. Smaller wire diameters increase the diffusion flux of O₂ during electrochemical reactions, leading to higher corrosion current density and degradation rate [94,95]. The significantly decreased diameters from 1.5 mm to 0.22 mm also possibly increase the degradation rates.

During the degradation of wires, the release concentrations of Li⁺ are 2-5 times higher than those of Zn⁺, which is likely correlated with dealloying. Dealloying is a phenomenon in which the element atoms with higher electrochemical activity in alloys undergo preferential dissolution [96]. Zn, Cu and Li are the consistent elements of the β-LiZn₄ matrix in wires. Their standard electrode potentials are −0.762 V, 0.342 V and −3.04 V, and the electrochemical activity follows the order: Li > Zn > Cu. Therefore, during corrosion, Li atoms in the β-LiZn₄ matrix are preferentially dissolved to form Li⁺. The dissolution of Zn atoms into Zn²⁺ is subsequent, while that of Cu atoms is last. Cu is relatively inert during the degradation of alloys, reflected by the non-detected Cu²⁺ due to too low contents. After the dissolution of Li and Zn, Cu is enriched in the corroded phase, exhibiting greatly higher contents compared with the non-corroded phase (Fig. 9, Fig. 10). Actually, among the metallic elements, Li has the lowest standard electrode potential, thereby exhibiting the highest electrochemical activity. The dealloying of Li is also observed during the corrosion of Al-Li-Cu [97,98], Li-Sn [99] and Zn-Li [19,100] based alloys.

The degradation mechanisms of Zn-2Cu-0.8Li wires are summarized as follows (Fig. 15): (1) The η-Zn as the anode and β-LiZn₄ as the cathode form micro-galvanic couples, and the η-Zn is preferentially corroded to form Zn²⁺. (2) Meanwhile, Li atoms in Li-rich phase β-LiZn₄ are preferentially dissolved into Li⁺ due to their higher electrochemical activity, and subsequently, partial Zn atoms are also dissolved into Zn²⁺. (3) The dealloying of Li and Zn at local regions of β-LiZn₄ results in the enrichment of Cu, improving the electrochemical inertness. These Cu-rich regions act as the cathode, forming micro-galvanic couples with the surrounding β-LiZn₄ matrix, thereby accelerating the degradation of other β-LiZn₄ regions and driving the overall degradation process. (4) Degradation products containing Zn and Li form, including ZnO, Zn(OH)₂, Zn₅(OH)₆(CO₃)₂, Zn₃(PO₄)₂, LiOH and Li₂CO₃. The Ca/P compounds also deposit on the surface of the alloy substrate. Several compounds containing Cu also form. Among them, LiOH, Li₂CO₃ and ZnO show the Pilling-Bedworth ratio (PBR) close to 1 [69], thereby owing good compactness to protect the surface from corrosion. Consequently, the degradation rates gradually decrease in the later stage of degradation.

Fig. 15.

Schematic diagram of microstructure and properties of Zn-2Cu-0.8Li wires and staples.

4.4. Prospects of alloy wires for staples

In summary, Zn-2Cu-0.8Li alloy fine wires with a diameter of 0.22 mm were successfully fabricated via hot extrusion, multi-pass drawing at room temperature and annealing treatment. As shown in Fig. 15, the microstructure of 1.5 mm wires consists of equiaxed β-LiZn₄ grains and needle-like η-Zn. After drawing, the microstructure of 0.22 mm wires consists of elongated deformed β-LiZn₄ grains with LAGBs, CDRXed fine β-LiZn₄ grains and dispersive η-Zn fine grain bands, accompanied by the significantly increased basal fiber texture intensity. The strength dramatically increases due to LAGBs, HAGBs and texture strengthening, while good deformability is ensured due to DRV and CDRX. The wires degrade follow the mechanisms discussed in section 4.1.3, exhibiting a uniform degradation mode and a suitable degradation rate. The staples made from these wires exhibit outstanding mechanical properties comparable to commercial non-degradable pure Ti staples, ensuring excellent operability during tissue penetration and anastomosis deformation, and affording sufficient support for favorable tissue healing without leakage. Simultaneously, they also exhibit an appropriate degradation rate and acceptable biosafety, meeting the clinical requirements. These demonstrate the broad application prospects of Zn-2Cu-0.8Li alloy staples.

In the future, when scaling up the wire production and staple fabrication for clinical use, some challenges also exist. For example, the non-linear translation and re-validation of critical process parameters: the parameters such as drawing passes and heat treatment may not scale linearly with increased billet size and continuous processing demands, and require recalibration; ensuring batch-to-batch consistency through rigorous process control; preventing fine wire fracture through alloy purification by optimizing casting techniques; fabricating staples with optimized structures to meet varied clinical needs … These critical issues merit further investigation and resolution.

5. Conclusions

In this study, Zn-2Cu-0.8Li alloy 0.22 mm fine wires with high strength and ductility for surgical staples were developed. The microstructure, mechanical properties evolution, and degradation behaviors of wires were systematically studied, and the wires were fabricated into staples to verify the application feasibility. The main conclusions are as follows:

• (1)The microstructure of as-extruded 1.5 mm wires consists of β-LiZn₄ equiaxed grains and the dispersive η-Zn needle-like phase. After multi-pass drawing, β-LiZn₄ grains are elongated along the DD to form threadlike bands. Meanwhile, DRV and CDRX occur, generating LAGBs and fine grains, resulting in the increased LAGBs fraction from 5.8% to 51.2% and the decreased average grain size from 11.6 μm to 3.0 μm. The basal fiber texture intensity is improved by nearly 8 times. The η-Zn needle-like phase is deformed along the DD, and transforms into fine grain bands due to DRX.
• (2)After drawing, the TYS and UTS of wires significantly increase from 376 MPa and 423 MPa to 596 MPa and 631 MPa. The main strengthening contributions are LAGBs, HAGBs and texture strengthening. The wires exhibit good deformability with the EL remaining at 25.4%, which is due to the occurrence of DRV and CDRX, as well as the deformation coordination of η-Zn. The wires annealed at 120 °C for 1 h exhibit the optimal mechanical properties with TYS of 492 MPa, UTS of 537 MPa and EL of 44.2%.
• (3)Compared with 1.5 mm raw wires, 0.22 mm-120 wires show a uniform in vitro degradation mode, and an increased degradation rate from 112 μm∙year⁻¹ to 327 μm∙year⁻¹ induced by the increased GBs density. The dispersive η-Zn acting as the anode is preferentially degraded to form Zn²⁺. The β-LiZn₄ matrix acting as the cathode is subsequently degraded with the prior dissolution of Li and Zn atoms and enrichment of Cu atoms.
• (4)The staples fabricated by wires exhibit a satisfying ultimate tensile force of 1.86 N and appropriate degradation behaviors in vitro. They successfully penetrated the gastric tissue of beagles and achieved effective anastomosis without leakage. At postoperative 12 months, the wound healed well, and the staples degraded by 68%.

CRediT authorship contribution statement

Xiyuan Zhang: Writing – original draft, Methodology, Investigation, Conceptualization. Zilong Wang: Visualization, Software, Data curation. Chun Chen: Methodology, Formal analysis. Wenhao Zhou: Methodology, Investigation. Cheng Liu: Resources, Methodology. Jimiao Jiang: Methodology, Formal analysis. Zhiqiang Gao: Methodology, Investigation. Jialin Niu: Writing – review & editing, Methodology, Conceptualization. Hua Huang: Writing – review & editing, Methodology, Conceptualization. Guangyin Yuan: Writing – review & editing, Supervision, Funding acquisition.

Ethics approval and consent to participate

The animal experiment involving gastric resection and anastomosis in beagles was approved by the Animal Ethics Committee of Sanitation & Environment Technology Institute, Soochow University (approval number: SDWH-M202102548-1), and conducted in strict compliance with relevant guidelines and standards.

Declaration of competing interest

The authors declare the following personal relationships which may be considered as potential competing interests: Cheng Liu is currently employed by Suzhou Yuezhong Biotechnology Co., Ltd.

Appendix A. Supplementary data

The following is the Supplementary data to this article: