Investigation on the microstructure, mechanical properties, in vitro degradation behavior and biocompatibility of newly developed Zn-0.8%Li-(Mg, Ag) alloys for guided bone regeneration

Materials Science & Engineering C 99 (2019) 1021–1034

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Investigation on the microstructure, mechanical properties, in vitro degradation behavior and biocompatibility of newly developed Zn-0.8%Li(Mg, Ag) alloys for guided bone regeneration

T

Yu Zhanga,b, Yang Yana, Xuemei Xua, Yujiao Lua,c, Liangjian Chenc, Ding Lid, Yilong Daia, ⁎ ⁎ Yijun Kangd, , Kun Yua,b, a

School of Materials Science and Engineering, Central South University, Changsha 410083, China Department of Materials Science and Engineering, Yantai Nanshan University, Yantai 265713, China Xiangya Third Hospital, Central South University, Changsha 410013, China d The Second XiangYa Hospital, Central South University, Changsha 410011, China b c

A R T I C LE I N FO

A B S T R A C T

Keywords: Zn-Li alloys Microstructure LiZn4 precipitate Mechanical properties Corrosion behavior Biocompatibility

In order to develop a biodegradable guided bone regeneration membrane with the required mechanical properties and high corrosion resistance, Zn-0.8%Li(wt), Zn-0.8%Li-0.2%Mg(wt), and Zn-0.8%Li-0.2%Ag(wt) alloys were cast and hot rolled into 0.1-mm thick sheets. The main secondary phase in Zn-0.8%Li-(Mg, Ag) alloys was the LiZn4 nanoprecipitate. Following the addition of minimal amounts of Mg, the tensile strength of the Zn-0.8% Li-0.2%Mg alloy improved, albeit with a greatly reduced elongation and corrosion resistance. The addition of minimal amounts of Ag refined the microstructure, producing fine equiaxed grains (2.3 μm) in the Zn-0.8%Li0.2%Ag alloy, and promoted a uniform distribution of LiZn4 nanoprecipitates with increased density and refined size. Therefore, the Zn-0.8%Li-0.2%Ag alloy exhibited optimal tensile strength and the highest corrosion resistance, with its elongation reaching 97.9 ± 8.7%. The corrosion products of Zn-0.8%Li-(Mg, Ag) alloys immersed in Ringer's solution for 35 days mainly consisted of zinc oxide and zinc carbonate. In addition, the cytotoxicity test using L929 cells and the evaluation of bone marrow mesenchymal stem cell proliferation indicated that the Zn-0.8%Li-0.2%Ag alloy had good biocompatibility.

1. Introduction Guided bone regeneration (GBR) is a clinical therapy for periodontal bone regeneration, based on providing the space for new bone formation by employing barrier membranes [1]. GBR membranes must fulfill certain design criteria, including biocompatibility, occlusivity, space shielding, clinical manageability, and the osseo-integration [2]. However, GBR membranes currently available for clinical applications are not simultaneously biodegradable and biocompatible, while providing space shielding. Collagen membranes, which are typical representatives of biodegradable membranes, have high biocompatibility and biodegradability [3,4]. However, a single collagen membrane does not possess the sufficient mechanical strength to provide stable space shielding of the defect area [5]. Titanium membranes are among the most widely used non-biodegradable membranes in clinical implantation dentistry due to their high mechanical strength. Nevertheless, when bone regeneration is completed, a second surgery is required to remove the

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non-biodegradable implant, which is harmful to the body [2]. Therefore, the development of a novel biodegradable metallic membrane is able to provide the bone defect area with the necessary space for tissue ingrowth and possessing good biocompatibility, is urgently required. In recent years, investigations into metallic biodegradable materials have mainly considered Fe- and Mg-based alloys as potential candidate materials [6]. However, previous studies have shown that Fe- and Mgbased alloys do not fully satisfy the requirements for implantation. Fe corrodes at a slow rate, yet the large amount of corrosion products repel neighboring cells and inhibit the formation of a biological matrix. Additionally, the released Fe2+ is not easily excreted from the body, resulting in iron poisoning [7]. Pure Mg and its alloys possess good biocompatibility and biodegradability, yet they rapidly corrode in body fluids, resulting in a reduction of the mechanical properties. Furthermore, an associated release of hydrogen gas during degradation is also harmful for tissue regeneration [8]. Zn-based alloys have recently emerged as novel metallic

Corresponding authors. E-mail addresses: [email protected] (Y. Kang), [email protected] (K. Yu).

https://doi.org/10.1016/j.msec.2019.01.120 Received 11 August 2018; Received in revised form 15 January 2019; Accepted 25 January 2019 0928-4931/ © 2019 Published by Elsevier B.V.

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550 °C, achieving the ingot dimensions Ø60 mm × 200 mm. The actual composition of all experimental alloys was analyzed by inductively coupled plasma atomic emission spectrometry (ICP-AES). Annealing at 400 °C was conducted for 24 h in order to homogenize the cast structure, followed by water quenching. The annealed alloys were then hot rolled with intermediate annealing to obtain sheets of 0.1 mm in thickness. Before every rolling procedure, sheets were placed in a box furnace at 250 °C for several minutes until reaching the furnace temperature.

biodegradable materials [9]. They appear to be free of the abovementioned limitations of Fe- and Mg-based alloys. The standard corrosion potential of pure Zn is −0.763 V, which is higher than that of pure Mg (−2.37 V) and lower than that of pure Fe (−0.44 V). Thus, Znbased alloys have a respectively slower and faster corrosion rate than those of Mg- and Fe-based alloys [6,10]. As a biodegradable metallic stent, Zn harmlessly degrades at a rate of 10–20 μm/year, making it an ideal biodegradable material [11,12]. Additionally, Zn is the second most nutritionally necessary element in the human body [13], being present in all tissues and in > 200 types of enzymes and being involved in physiological metabolic processes of the human body. The daily intake value of Zn is 10 mg/day for adults and 2 mg/day for infants [14]. Further, Zn plays critical roles in immune system, the nervous system, metabolism, and body growth [15,16]. Therefore, Zn is relatively nontoxic and biocompatible for humans. With regards to orthopedic implants, pure Zn exhibits a very low tensile strength unable to meet the mechanical requirements [17,18]. Adding alloying elements is therefore an effective way to improve the mechanical properties of pure Zn [19–21]. Li is one of the commonly used elements to enhance the mechanical properties of pure Zn. Zhao et al. [22,23] studied Zn-Li alloys as endovascular stents and found that, with the addition of 6 at.% Li into Zn, the ultimate tensile strength (UTS) increased from 120 MPa (pure Zn) to 560 MPa. Furthermore, they implanted 10 mm Zn-Li wire segments into the abdominal aorta of rats, showing an excellent biocompatibility and promising corrosion rate. Therefore, herein, Zn-Li-based alloys was selected to be studied as a potential biodegradable membrane for periodontal tissue regeneration. Mg has a good biocompatibility and can stimulate new bone formation [24,25]. Additionally, Mg has a strong strengthening effect on Zn-based alloys due to the formation of hard intermetallic phases (Mg2Zn11 and MgZn2). However, a high Mg content in Zn alloys results in a reduction of mechanical properties due to the excessive eutectic intermetallic phases [26,27]. As another important alloying element, Ag can enhance the mechanical properties of Zn while preserving its biocompatibility and its addition can effectively refine grain size and enhance the mechanical properties of Zn alloys [28]. Furthermore, given its good antibacterial properties, Ag-containing materials have been used for dental implants in the clinic [29,30]. Of note, the reference dose for Ag is 0.35 mg/day for a standard 70 kg person, and excessive Ag can lead to argyria [31]. Thus, minimal amounts of Mg and Ag elements can be added to Zn-Li alloys to enhance their mechanical properties and improve their biocompatibility. To the best of our knowledge, the formation of Zn-Li-X ternary alloys has not been previously reported in the literature. Therefore, herein, we assess the effects of the addition of minimal amounts of Mg/ Ag on the mechanical properties and corrosion behavior of Zn-Li alloys, as well as the corrosion behavior of Zn alloys in various simulated body fluids. For this purpose, as-rolled Zn-0.8%Li, Zn-0.8%Li-0.2%Mg, and Zn-0.8%Li-0.2%Ag alloys were studied in detail, including their microstructure, mechanical properties, corrosion behavior, cytotoxicity, and biocompatibility.

2.2. Microstructure characterization The microstructure was examined using optical microscopy, scanning electron microscopy (SEM, Quanta 200) with an energy dispersive spectrometer (EDS), and transmission electron microscopy (TEM, Tecnai G2 F20). All samples for microscopic observation were polished and etched with an etchant consisting of 1.5 g sodium sulfate, 20 g chromium oxide, and 100 mL distilled water. Additionally, X-ray diffraction (XRD), using a DMAX-2500× apparatus, was employed for the identification of constituent phases. Diffraction patterns were generated with the values of 20–90° at a scanning speed of 4°/min. Specimens for TEM observation mechanically thinned to 80 μm, and then reduced by electrolytic jet polishing in a solution of 10% HClO4 + 90% C2H5OH at 20 V between −30 °C and −20 °C. 2.3. Mechanical tests Mechanical tests of samples were performed through tensile tests using a universal testing machine (Instron 3369). Samples were cut from the sheet with a gauge length of 50 mm. Testing results were averaged over three samples. The fracture morphologies were investigated by SEM. 2.4. Immersion tests Immersion tests were performed in Ringer's solution (pH 7.4) at 37 ± 0.2 °C and 5% CO2 atmosphere for 35 days to evaluate the pH variation and weight change. The samples were cut into 1 × 1 cm2 pieces. The exposed surface was then ground using 200–1000# SiC papers, washed by distilled water, and then air dried and weighed (W0). All samples were sterilized by ultraviolet radiation and tests were performed on a super-clean worktable. The ratio of the solution volume to sample surface area was 20 mL/cm2 [32]. 2.4.1. Surface corrosion morphology and corrosion product Samples immersed in Ringer's solution after 7 days and 35 days were removed from solutions and observed by SEM. Prior to observation, samples were ultrasonically cleaned in ethanol for 5 min and dried at room temperature. For all samples, Au spraying was used on the surface to elevate their conductivity. All observations were performed using secondary electron images. Functional groups in the corrosion products were recorded by Fourier transform infrared analysis (FTIR, Nicolet 6700). The spectra were recorded from 4000 to 800 cm−1.

2. Materials and methods 2.4.2. Electrochemical measurements Electrochemical measurements were performed at 37 ± 0.2 °C using a standard three-electrode configuration consisting of a saturated calomel as the reference electrode, platinum mesh as the counter, and the sample as the working electrode. A CHI660C potentiostat/galvanostat system was used for all electrochemical tests. The exposed surface was ground by 200–1000# SiC papers and the surface was then polished with ethanol. Potentiodynamic polarization curves were tested at a scan rate of 0.5 mV s−1. Prior to measurement, the open circuit potential was tested for the sample to reach a stable state. The results, including corrosion potential (Ecorr) and corrosion current density (icorr), were calculated by the Tafel extrapolation method, and the

2.1. Material preparation Zn-0.8%Li (wt), Zn-0.8%Li-0.2%Mg (wt), and Zn-0.8%Li-0.2%Ag (wt) alloys were prepared under an Ar atmosphere using pure zinc (> 99.99%, Huludao Zinc Industry Co.,China), pure Li (> 99.9%, Hunan Rare Earth Metal Material Research Institute, China.), pure Mg (> 99.9%, YinGuang Magnesium Industry [group] Co., Ltd., China), and pure Ag (> 99.9%, Hunan Rare Earth Metal Material Research Institute, China.). Samples were melted at 580 °C in graphite crucibles in a resistance furnace under an Ar atmosphere, and then cast into a preheated cylindrical steel mold at a temperature of approximately 1022

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corrosion rate was derived from current density icorr (mA·cm−2) using the equation [33]:

Corrosion rate = K ∗ icorr ∗ EW/ρ

(1)

−3

where K is 3.27 × 10 in mm·g/μA·cm·yr, icorr is the corrosion current density in μA/cm2, EW is the equivalent weight of Zn (with a value of 32.68), and ρ is the density of Zn-based alloys (7.14 g/cm3). 2.5. Cytotoxicity tests and proliferation evaluation L929 and bone marrow mesenchymal stem cells (BMSCs) were used to assess cytotoxicity and effect on proliferation respectively evaluated by cell morphological observation and CCK8 assay. Extracts were taken from samples incubated in Dulbecco's Modified Eagle Medium supplemented with 10% fetal bovine serum for 72 h in a humidified atmosphere containing 5% CO2 at 37 °C, according to ISO 10993-5: 1999 [34]. The culture medium was respectively diluted to 50% and 10% concentration, and then refrigerated at 4 °C before the cytotoxicity test. Cells were seeded into three 96-well culture plates at a cell density 5 × 104 of cells/mL in each well and incubated in a humidified atmosphere containing 5% CO2 at 37 °C for 24 h. The cell culture medium was used as a negative control. For the cytotoxicity test, the medium was successively replaced with 100 μL by 100%, 50%, or 10% extracts, or a negative control (culture medium), or a positive control (containing 0.64% phenol). As for proliferation evaluation, the medium in each well was substituted by 100 μL of 100% extracts or negative control (culture medium) and then incubated for 1, 3, and 5 days, respectively. At each time point, one plate was selected and cell morphology was observed by optical microscopy. Then, 10 μL CCK8 solution (5 mg/mL) was added to each well and further incubated for 4 h. Finally, the medium was replaced by 100 μL of dimethyl sulfoxide and a spectrophotometer was used to calculate the absorbance of the supernatant at 475 nm. The cell relative growth rate (RGR) and cell viability were calculated as follows:

RGR = ODtest /ODnegative × 100%

(2)

Cell viability = ODtest /ODnegative × 100%

(3)

Fig. 1. XRD patterns of the Zn-Li-(Mg, Ag) alloys.

dissolved in the Zn matrix. The optical metallographic microstructure of the three alloys showed a fibrous structure along with the rolling direction (Fig. 2a, c, e), where the grains and the secondary phase were crushed, lengthened into a fibroid, and distributed along the rolling direction. From the microstructures in the cross-section perpendicular to the rolling direction (Fig. 2b, d, f), the alloys presented fine equiaxed grains, with a calculated average size below 5 μm. These results indicate that dynamic recrystallization occurred during the rolling process. The average grain size of the Zn-Li-Ag alloy (2.3 μm) was smaller than that of the Zn-Li alloy (4.1 μm), indicating that adding a Ag alloying element obviously refined the grains of the Zn-Li alloy. Conversely, in the Zn-Li-Mg alloy, an intermetallic phase (Mg2Zn11) with an elliptical shape uniformly appeared in the Zn matrix. However, the secondary phase (LiZn4) could not be distinctly observed, and therefore SEM and TEM were used. 3.2. SEM and TEM of the microstructure The back scattered electron images and EDS analyses of the Zn-Libased alloys in the rolling direction showed a similar structure of black stripe in the rolling direction (Fig. 3). Despite XRD (Fig. 1) showing the presence of a LiZn4 secondary phase in the alloys, the morphology and distribution of the LiZn4 is not clear in the SEM images. Since the Li atom, which has a small atomic radius, cannot be detected by the EDS. For the Zn-Li-Mg alloy, several black elliptical particles were observed along the rolling direction (Fig. 3b). According to both XRD and EDS results (Fig. 3d, e), the secondary phase was confirmed as Mg2Zn11. Conversely, a secondary phase was not observed for the Zn-Li-Ag alloy (Fig. 3c). Further, the Ag content in the black stripe area and the grey matrix was very similar, implying that Ag is uniformly distributed across the Zn matrix. The bright-field TEM images of the Zn-Li alloy showed fine grains and small subgrains (Fig. 4a), implying that recrystallization occurred during the rolling progress. The recrystallization temperature of Zn is so low that the thermal effect during rolling may lead to a recrystallization of the Zn-Li alloy, and thus a small grain size. Further, dislocation tangles and dislocation/precipitation interactions appeared in interior of the grain (Fig. 4b), showing a typical characteristic of dislocation pinned by precipitates. Finally, a large number of spherical particles were homogeneously distributed inside the grain (Fig. 4c), attributed to the secondary LiZn4 phase. The high magnification image of the Zn-Li alloy in the state A (Fig. 4c) is shown in Fig. 5a, indicating that the size of the precipitates is of approximately 10–30 nm, with a high density. Such a high density of fine dispersive precipitates could effectively impede the dislocation movements, leading to a strength increase in for alloy. The HRTEM

3. Results and discussion 3.1. Microstructure The composition of the study alloys was analyzed by ICP-AES (Table 1), showing that the alloying elements Li, Mg, and Ag were close to their nominal compositions. The XRD patterns showed that Zn-Li and Zn-Li-Ag alloys were composed of a Zn matrix with a secondary LiZn4 phase, whereas the Zn-Li-Mg alloy consisted mainly of a Mg2Zn11 phase (Fig. 1), in agreement with previous results [22,23,35]. For the Zn-LiAg alloy, no diffraction peak arising from Ag or Ag-Zn secondary phase was detected due to the relatively low Ag content. Since Ag has a high solid solubility (about 8.0 wt% at 431 °C) in Zn according to the Zn-Ag equilibrium phase diagram [36]. This result is consistent with the observations of Sikora-Jasinska et al. [28], wherein the Ag-Zn secondary phase cannot be observed in Zn-Ag binary alloys with Ag content lower than 2.5%. In which case, Ag can be considered to be completely Table 1 The actual compositions of all experimental alloys analyzed by ICP-AES. Alloy

Zn-0.8%Li Zn-0.8%Li-0.2%Mg Zn-0.8%Li-0.2%Ag

Composition (wt%) Li

Mg

Ag

Zn

0.74 0.73 0.76

– 0.22 –

– – 0.23

Bal. Bal. Bal.

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Continue

Fig. 2. Microstructure of the as-rolled Zn-Li-(Mg, Ag) alloys along the rolling direction and in the cross-section perpendicular to the rolling direction: (a), (b) Zn-Li alloy; (c), (d) Zn-Li-Mg alloy; (e), (f) Zn-Li-Ag alloy.

intermetallic Mg2Zn11 phase. Following Ag alloying, the elongation of Zn-Li-Ag alloy was enhanced to 97.9 ± 8.7%, with a slight improvement in YS and UTS (196.2 ± 5.1 MPa, 254.7 ± 3.9 MPa), likely attributed to grain refinement. The tensile fracture morphologies of the three as-rolled alloys were also assessed (Fig. 7). The Zn-Li alloy showed a ductile fracture surface with dimples and a few shear lips (Fig. 7a, b), indicating that the plasticity of the Zn-Li alloy can be theoretically improved. The fracture surface of the Zn-Li-Mg alloy was essentially a mixed ductile-brittle mode, with dimples decreasing and a small amount of a brittle secondary phase emerging (Fig. 7c, d), which might contribute to the reduction of ductility. Finally, the Zn-Li-Ag alloy showed an intergranular fracture surface as well as several dimples (Fig. 7e, f), similar to a fracture surface experiencing superplasticity deformation. Refined grains distributed uniformly on a fracture surface result from grain boundary sliding (GBS) [37,38]. Therefore, GBS was possibly involved

microstructure of the precipitate with the corresponding fast Fourier transform (FFT) patterns (Fig. 5b) show a very clear particle–matrix interface, indicating that the precipitates have a regular shape. The FFT-filtered high-solution TEM image and the FFT pattern of the state A (Fig. 5b) show that the precipitates are LiZn4 with hcp structure, whereas the matrix is Zn in [0 1 –1 0]. 3.3. Mechanical properties and fracture surface The three alloys exhibited the good strength and high elongation (Fig. 6 and Table 2). The Zn-Li alloy showed a high yield strength (YS), ultimate tensile strength (UTS), and elongation of 183.5 ± 5.3 MPa, 238.1 ± 4.7 MPa, and 75 ± 6.0%, respectively. Following Mg alloying, the YS and UTS of the Zn-Li-Mg alloy increased to 253.7 ± 4.7 MPa and 341.3 ± 4.8 MPa, respectively, accompanied by a decrease in elongation to 30.6 ± 5.8% due to the presence of large 1024

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Fig. 3. The SEM micrograph and EDS analysis of the Zn-Li-(Mg, Ag) alloy: (a) Zn-Li; (b), (d), (e) Zn-Li-Mg; (c), (f), (g) Zn-Li-Ag.

Fig. 4. Bright-field TEM images of the Zn-Li alloy: (a) fine grains and subgrains, (b) morphology of dislocations, (c) distribution of the precipitates. 1025

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Fig. 5. (a) TEM images of the precipitates, (b) a high-solution TEM image of the interface between the precipitate and the matrix as well as the corresponding FFT (fast Fourier transform) patterns, (c) and (d) FFT-filtered high-solution TEM images of the state A and B in (b) respectively.

in plastic deformation of the Zn-Li-Ag alloy. The as-rolled Zn-Li-(Mg, Ag) alloys prepared herein exhibited excellent tensile properties (especially plasticity), superior to those of reported Zn-based alloys such as Zn-Mg or Zn-Ca alloys [19,39]. The excellent mechanical properties observed can be ascribed to the following factors.

Table 2 Mechanical properties of Zn-Li-(Mg, Ag) alloys.

(1) Precipitation strengthening. According to the phase diagram, Li has a relatively low solubility in Zn (about 0.12 wt% at 403 °C) and therefore the Zn-0.8%Li alloy investigated herein is hypereutectic. Under the casting conditions, a eutectic mixture was formed, resulting in the formation of a Zn matrix and LiZn4. During the process of annealing, LiZn4 dissolved in the Zn matrix to form the supersaturated solid solution, followed by the formation of LiZn4 nanoprecipitates and precipitation in the Zn matrix during the rolling process, as shown in Fig. 5a. Therefore, the LiZn4

Alloy

YS(MPa)

UTS(MPa)

δ(%)

Zn-0.8%Li Zn-0.8%Li-0.2%Mg Zn-0.8%Li-0.2%Ag

183.5 253.7 196.2

238.1 341.3 254.7

75.0 30.6 97.9

nanoprecipitates can be considered the main strengthening phases. Homogeneously distributed LiZn4 nanoprecipitates can play an effective role in strengthening the materials by impeding dislocation sliding. Moreover, during the process of rolling and tensile testing the precipitates likely act as the core of heterogeneous nucleation to refine the grain, resulting in a further strength increase. (2) Grain refinement. According to the Hall–Patch law (σs = σ0 + Kd−1/2), a fine grain would enhance the mechanical

Fig. 6. (a) Strain-stress curves and (b) Tensile properties of the as-rolled Zn-Li-(Mg, Ag) alloys. 1026

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Fig. 7. Tensile fracture morphologies of Zn-Li-(Mg, Ag) alloys: (a), (b) Zn-Li alloy; (c), (d) Zn-Li-Mg alloy; (e), (f) Zn-Li-Ag alloy.

superplasticity [40]. Herein, the grain of the alloys was < 5 μm, and therefore GBS may have occurred during the deformation, leading to high ductility. Particularly, the Zn-Li-Ag alloy showed superior ductility with a high elongation value (97.9 ± 8.7%), nearly achieving superplasticity at room temperature at a high strain rate (1.67 × 10−2 mm/s). The refined grains were uniformly distributed on the fracture surface (Fig. 7f), also showing solid evidence of GBS. Therefore, Zn-Li-(Mg, Ag) alloys exhibit high ductility due to the fine grains that could help hinder grain slide. (3) Dynamic recrystallization.

properties of an alloy. Additionally, refined grains can promote the plastic deformability. If the grain size is very small such that the grain boundary can slide, move, and rotate, it will be involved in the plastic deformation process, increasing the alloy plasticity or even leading to superplastic deformation. GBS is commonly the dominant deformation mechanism in superplasticity, and therefore high tensile ductility may be obtained in traditional materials when the grain size is decreased down to the micrometer range, typically < 10 μm [40]. Huang et al. reported that, while a refinement of grain size down to 1.3 μm in a Zn-22wt%Al alloy led to a room-temperature elongation of 125% at a strain rate of 10–3 mm/ s, a grain size decrease to submicron level (550 nm) led to a 275% elongation at the same test conditions [41]. A grain refinement is especially associated with maintaining room-temperature

The recrystallization temperature of zinc is so low, Zn alloys could have a fully recrystallization at room temperature [42]. The texture microstructures and fine equiaxed grains of alloys (Fig. 2) indicated the 1027

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samples undergone a strong plastic deformation during the rolling process, suggesting dynamic recrystallization occurred in the alloys. Fine grains increased, the orientation optimized, the alloy has a good plasticity for grain boundary effect that bearing stress concentration and crack extension. Shi et al. [43] reported that the as-rolled Zn0.34Mn and Zn-0.76Mn alloys showed excellent elongations (88.8% and 94.0%) at room temperature. The extraordinary ductility of the alloys was ascribed to the fine grains before the tensile test and dynamic recrystallization during the tensile test. Dynamic recrystallization during the tensile test may also have happened to help improve the plasticity for the alloys in this study, and this work is also being further studied. The addition of 0.2% Mg resulted in an increase in strength and a drop in elongation, attributed to the coarse and brittle Mg2Zn11 intermetallic phase. The Mg2Zn11 not only played the role of second-phase strengthening to enhance the strength of the alloy, but also provided the opportunity for crack initiation at the matrix/particle interface during tensile testing. As reported by Shen et al. [44], a Mg content of 1 wt% and above leads to an increase in the volume fraction of the brittle eutectic and, therefore, a deterioration in alloy ductility and fracture toughness. The addition of 0.2% Ag promotes both tensile strength and elongation due to further grain refinement, making the Hall–Patch effect more obvious. Ag has a high solubility in Zn alloys (about 8 wt% at 431 °C) and Ag atoms were completely dissolved in the Zn matrix in the Zn-0.8%Li-0.2%Ag alloy. Due to the solid solution strengthening, Zn-0.8%Li-0.2%Ag alloy had a higher strength compared with that of the Zn-0.8%Li alloy.

Fig. 9. Polarization curves of Zn-Li-(Mg, Ag) alloys in the Ringer's solution. Table 3 Electrochemical data of Zn-Li-(Mg, Ag) alloys in Ringer's solution. Alloy

Ecorr(V)

Icorr(μA/cm2)

Vcorr(mm/year)

Zn-Li Zn-Li-Mg Zn-Li-Ag

−1.29 −1.32 −1.21

8.24 11.31 7.67

0.12 0.17 0.11

the corrosion product prior to weighing, therefore more clearly reflecting the integrity of the sample. A lower weight change indicated a better integrity and a lower strength loss. The weight change decreased in the order of Zn-Li-Mg > Zn-Li > Zn-Li-Ag, showing that the Zn-LiAg alloy had the highest corrosion resistance. The polarization curves of three alloys in Ringer's solution (Fig. 9) and the electrochemical parameters and corrosion rates obtained from electrochemical measurement (Table 3) showed that, compared to the corrosion potential and current density of the Zn-Li alloy (−1.29 V, 8.24 μA/m2), the Zn-Li-Mg alloy had a relatively negative potential (−1.32 V) and higher current density (11.31 μA/m2), whereas the ZnLi-Ag alloy had the most positive corrosion potential (−1.21 V) and the lowest current density (7.67 μA/m2). The addition of Mg to the Zn-Li alloy led to the formation of galvanic effects in the Zn matrix and Mg2Zn11 phase, accelerating the corrosion rate. Conversely, the addition of Ag had a positive effect on the corrosion resistance of the Zn-Li binary alloy, attributed to the finer and homogeneous microstructure of the Zn-Li-Ag alloy. Of note, the three curves exhibited similar

3.4. Immersion test and electrochemical measurement The pH value variation of the three alloys in Ringer's solution as a function of immersion time (Fig. 8a) was measured. The lower increasing rate of pH value reflects a higher corrosion resistance of the alloys since the generation of OH– ions would lead to an increase in pH [45]. For the three alloys, the pH value showed a sharp increase over the initial 7 days followed by a slower increase, reaching a peak value approximately on day 21. This behavior indicated an initial dissolution of the Zn matrix followed by the formation of a protective passive film. The pH value at culmination was ranked as Zn-Li-Mg > Zn-Li > ZnLi-Ag. The maximum pH value among the alloys was 9.3, which was lower than the pH value that could lead to cell death and bone tissue inflammation [45,46]. Furthermore, the weight change of the three alloys immersed in Ringer's solution for 35 days was calculated by the formula: C = (m1–m0)/m0 [47], where C is the weight change, m1 is the weight following immersion, and m0 is the original weight (Fig. 8b). In this experiment, the samples were dried without removing

Fig. 8. (a) The pH value variation and (b) the weight change of Zn-Li-(Mg, Ag) alloys in Ringer's solution with a function of immersion time. 1028

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Fig. 10. The surface morphologies of three alloys immersed in Ringer's solution before and after removal of corrosion products, and the EDS analyses of corrosion product as well as the morphologies of cross section for 35 days. From (a) to (c), representing Zn-Li, Zn-Li-Mg and Zn-Li-Ag alloy; 1 and 2 mean the immersion time are 7 days, 3–5 mean the immersion time are 35 days.

were characterized (Fig. 10). Additionally, the EDS analyses of the corrosion products as well as the morphologies of the cross-section after 35 days were performed. Small and limited amounts of corrosion products were observed on the surface after 7 days. The initial corrosion products on the surface of the Zn-Li alloy were round and cloud-like aggregations (Fig. 10a1). Following removal of the corrosion product, the Zn-Li alloy showed localized corrosion with localized corrosion pits (Fig. 10a2). The Zn-Li-Mg alloy exhibited a surface composed of elongated crystals with the accumulation of voluminous corrosion products in local area (Fig. 10b1). Following corrosion product removal, the surface exhibited characteristic non-uniform corrosion, accompanied by a large area of corrosion pits (Fig. 10b2). The Zn-Li-Ag surface displayed cotton-like structures with a uniform distribution (Fig. 10c1), with the Zn-Li-Ag alloy substrate exhibiting finer corrosion pits and less corrosion area (Fig. 10c2) compared to the two other alloys. After 35 days of immersion, the corrosion products of the three alloys were accumulated. The cloud-like corrosion products on the surface of the Zn-Li alloy became larger and sparse filled with some large holes (Fig. 10a3). The Zn-Li alloy substrate indicated severe corrosion as exhibited by large and severe corrosion pits and pit interconnection, resulting in corrosion holes (Fig. 10a4). The surface of the Zn-Li-Mg alloy was covered with a large corrosion layer, in addition to cracks appearing on the corrosion layer (Fig. 10b3). The corrosion area was also large and severe as observed by the surface of the Zn-Li-Mg alloy following corrosion product removal (Fig. 10b4). Particularly, the presence of uniformly distributed pores on the substrate was observed, attributed to the high corrosion rate. The surface of Zn-Li-Ag was covered with a compact layer of globular corrosion products, 10–15 μm in size, homogenously distributed and closely bonded (Fig. 10c3). The corrosion pits were comparably small and uniform (Fig. 10c4). The corrosion depth of the alloys effectively reflected the corrosion resistance of the alloys; thus, the maximal corrosion depth on the crosssection of the Zn-Li-(Mg, Ag) alloys after immersion for 35 days were measured in Fig. 10a5–c5. Both the Zn-Li and Zn-Li-Mg alloys showed deep corrosion pits (Fig. 10a5, b5). The deepest corrosion pits were approximately of 24 μm, observed on the surface of the Zn-Li-Mg alloy (Fig. 10b5), indicating that it had the lowest corrosion resistance among the three alloys. The Zn-Li alloy showed a uniform corrosion depth, covered with a layer of corrosion products, likely protecting the inner Zn matrix from further corrosion (Fig. 10c5). The EDS analyses showed that the corrosion product of the three alloys mainly consisted of Zn, O and C (Fig. 10a6–c6), suggesting that the corrosion products were composed mainly of zinc carbonate and zinc oxide. The FTIR spectra of the corrosion products of the three alloys following immersion for 35 days (Fig. 11) showed a broad absorption from 3600 to 3100 cm−1, ascribed to the OeH stretching vibration of the hydroxyl group [49]. The peak at 1640 cm−1 was attributed to H2O bending vibration and rotation modes, indicating the presence of crystal water [49,50]. The peaks observed within the 830–890 cm−1 wavenumber range were attributed to CO32– (v2) band, as previously reported [51,52]. The weak peaks between 1470 and 1420 cm−1 were assigned to the stretching vibration of absorbed CO32– from the solution [52,53], observed with a slightly stronger intensity in the corrosion products of the Zn-Li-Mg alloy. XRD was performed to further confirm the composition of the corrosion products of the Zn-Li-(Mg, Ag) alloys after 35 days of immersion (Fig. 12). The combination of the results of EDS, FTIR, and XRD indicated that the corrosion products were mainly composed of zinc carbonate and zinc oxide. Therefore, according to the above results, a mechanism of corrosion of the Zn-Li-(Mg, Ag) in Ringer's solution was proposed. For Zn immersed in neutral and alkaline solutions, the corrosion reactions are

Fig. 11. FTIR spectra of the corrosion products of Zn-Li-(Mg, Ag) alloys after immersion for 35 days.

Fig. 12. The XRD patterns of the Zn-Li-(Mg, Ag) alloys after immersion for 35 days.

characteristics, with a successive long current plateau being observed in the anodic branch of the polarization curves, representing a high corrosion resistance of alloys through the formation of a dense passive film on the sample surface. The starting point of the current plateau indicated the beginning of the formation of a passive film, whereas the plateau region signifies thickening and growth of the passive film and the end point of the current plateau represents breakdown of the passive film. The stability of the passive film can also be characterized by the potential at which it breaks down. Generally, the more positive potential when the passive film breaks, the more stable it is to corrosion [48]. The Zn-Li-Ag alloy exhibited the most positive potential at the end point of the current plateau, implying the highest corrosion resistance among the three alloys. In addition, the corrosion rates calculated from the corrosion current density were ranked in decreasing order, as follows: Zn-Li-Mg > Zn-Li > Zn-Li-Ag, in agreement with the results of corrosion rate following immersion in Ringer's solution. 3.5. Corrosion behavior and mechanism The surface morphologies of three alloys immersed in Ringer's solution for 7 and 35 days before and after removal of corrosion products 1030

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Fig. 13. Cytotoxicity results of the Zn-Li-Ag alloy: (a) Relative growth rate (RGR) of L-929 cells cultured in different extracts. Optical images of L-929 cells that cultured in (b) negative control, (c) 10% extraction, (d) 50% extraction, (e) 100% extraction (f) positive control (0.64% phenol) for 5 days.

[39]:

Cathodic reaction: 2H2 O + 2e = H2 + 2OH− Anodic reaction: Zn = Zn2 + + 2e Total reaction: Zn + 2H2 O = H2 + Zn(OH)2

(4)

Zn(OH)2 = ZnO + H2 O

(5)

Based on the Pourbaix diagram, Zn tends to be passivated in neutral or slightly alkaline environments. Zn(OH)2 precipitates on metal surfaces due to its low solubility, and the accumulation of Zn(OH)2/ZnO tends to form a protective passive film, retarding the corrosion process. However, in aqueous environments containing Cl−, Zn(OH)2 may be transformed into ZnCl2. Subsequently, ZnCO3 will spontaneously precipitate on the surface due to its low solubility [54] and Zn5(OH)8Cl2 (simonkolleite) will form on the surface. This kind of intermediate corrosion product can prevent the substrate from direct exposure to the environment, improving the corrosion resistance for the following reactions [39]:

Fig. 14. Cell viability of the Zn-Li-Ag alloy after 1, 3 and 5 days. 1031

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Fig. 15. Optical morphologies of BMSCs after 5 days of incubation. (a, c, e) Negative Control; (b, d, f) Zn-Li-Ag alloy extracts.

Zn(OH)2 + 2Cl− = ZnCl2 (soluble)

(6)

Zn2 + + H2 CO3 = ZnCO3 + 2H+

(7)

4Zn(OH)2 + ZnCl2 = Zn5 (OH)8 Cl2

(8)

flaws on the corrosion products, which increase the possibility of Cl− ion penetration and of other ions. On the other hand, the corrosion product layer of Zn-Li-Ag alloy was compact, inhibiting the Zn matrix from further corrosion. Additionally, the presence of a coarse Mg2Zn11 intermetallic phase in the Zn-Li-Mg alloy led to an acceleration of corrosion of the Zn matrix due to their potential difference [55]. The formation of pores on the surface of the Zn-Li-Mg alloy may be attributed to the H2 flow created by the dissolution of Zn or Mg, implying the high corrosion rate of the Zn-Li-Mg alloy. The corrosion of Zn-Li-Ag alloy was weak and uniform, with shallow and even corrosion pits, attributed to its homogeneous microstructure since the compact passive corrosion product layer formed on the surface protected the substrate. Although the corrosion products were largely zinc oxide and zinc carbonate, a small quantity of Cl− and Mg2+ were also detected in the EDS. Thus, we believe the simonkolleite and magnesium hydroxyl carbonate are also likely corrosion products.

As for Zn-Li-Mg alloy, magnesium hydroxyl carbonate may be formed on the surface, according to Eq. (9) [39]:

2Mg2 + + 2OH− + CO32– = Mg 2 (OH)2 CO3

(9)

The Zn-Li-(Mg, Ag) alloys are mainly composed of Zn matrix and a LiZn4 secondary phase, thus this micro-galvanic corrosion might be the main corrosion pathway. Additionally, the LiZn4 secondary phase will act as the cathode in the galvanic effect between the Zn matrix and the LiZn4 secondary phase, due to the more noble potential of LiZn4 over that of Zn [22]. During the beginning of corrosion, the Zn matrix around the secondary phase was preferentially corroded due to microgalvanic corrosion and the corrosion products attached on the surface. However, the corrosion products did not initially cover (Fig. 10) the substrates completely and the surface exposed to the solution could be further attacked. Subsequently, the corrosion increased, leading to a greater amount of corrosion products being accumulated and an increase in number and depth of corrosion pits with increasing time. Finally, the corrosion product layer acted as a good barrier for substrate protection from further corrosion. Furthermore, the low corrosion resistance of Zn-Li and Zn-Li-Mg alloys might be related to the pits and

3.6. Cytotoxicity test According to the above studies, the Zn-Li-Ag alloy demonstrated the highest corrosion resistance as well as superior mechanical properties, and was therefore selected for cytotoxicity testing. The morphology of L929 cells after 5 days of incubation in different extract mediums with concentrations of 10, 50, and 100% showed that L929 cells were healthy and exhibited a flattened spindle shape, similar to that in the 1032

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negative control (Fig. 13b–f). The RGR of L929 cells after 1, 3, and 5 days of incubation in the extract medium showed no statistical difference between the different extract mediums and the negative control (p > 0.05) (Fig. 13a), while there was a significant difference between the extract groups and the positive control (p < 0.05). According to ISO 10993-5: 1999, the cytotoxicity of these extracts of Zn-Li-Ag alloy was of Grade 0–1. In other words, the Zn-Li-Ag alloy is innocuous and satisfactory for bone tissue engineering.

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3.7. Proliferation evaluation of BMSCs BMSCs have the potential to differentiate into osteoblastic cells under specific conditions. Scaffold materials containing BMSCs have a good therapeutic effect on bone defects, promoting the formation of new bone [56]. Thus, it is important to study the proliferation of BMSCs on Zn-Li-Ag alloy. The calculated cell viability of Zn-Li-Ag alloy after 1, 3, and 5 days showed a high cell viability value at each time point (Fig. 14), suggesting that the Zn-Li-Ag alloy did not play a negative effect on the proliferation of BMSCs. The optical images of the morphologies of BMSCs after 1, 3, and 5 days of incubation in the extraction medium and the negative control showed normal and healthy cells, exhibiting well spread and affluent pseudopods (Fig. 15), indicating a high viability in the extract medium compared with in the negative control. In conclusion, the Zn-Li-Ag alloy has a good biocompatibility with BMSCs, making it a potentially suitable scaffold material for BMSCs for application in bone tissue engineering.

4. Conclusion In this study, newly developed Zn-0.8%Li, Zn-0.8%Li-0.2%Mg, and Zn-0.8%Li-0.2%Ag alloys were investigated as potential biodegradable materials for GBR membranes. The following conclusions can be drawn. 1) The as-rolled Zn-Li, Zn-Li-Mg, and Zn-Li-Ag alloys were mainly composed of a Zn matrix and a LiZn4 secondary phase, with a small amount of Mg2Zn11 intermetallic phase identified in the Zn-Li-Mg alloy. The LiZn4 precipitate was uniformly distributed in the Zn matrix, with a particle size of 10–30 nm. 2) With the addition of minimal amounts of Ag, the Zn-Li-Ag alloy showed an optimal tensile property, with an elongation of 97.9 ± 8.7%. This was due to the combination of grain strengthening and precipitation strengthening. 3) With the addition of a minimal amounts of Mg, the Zn-Li-Mg alloy exhibited the highest yield strength and ultimate tensile strength, but a lower elongation than other alloys. The coarse intermetallic Mg2Zn11 phase contributed to the increase of tensile properties and the reduction of elongation. 4) The corrosion rates of the Zn-0.8%Li-(Mg, Ag) alloys immersed in Ringer's solution for 35 days were ranked as Zn-Li-Mg > ZnLi > Zn-Li-Ag. The corrosion products mainly consisted of zinc oxide and zinc carbonate. 5) The Zn-Li-Ag alloy had a good biocompatibility, along with suitable mechanical properties and corrosion resistance. Therefore, the ZnLi-Ag alloy shows great potential for biomedical applications.

Acknowledgements The authors acknowledge of the 2015 ShanDong Province project of outstanding subject talent group, the Natural Science Foundation of ShanDong Province of China (ZR2017MEM005), the project (2017GK2120) supported by the Key Research and Development Program of Hunan Province and the Natural Science Foundation of Hunan Province of China (2018JJ2506). 1033

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