Materials Science and Engineering C 47 (2015) 85–96

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In vitro degradation and electrochemical corrosion evaluations of microarc oxidized pure Mg, Mg–Ca and Mg–Ca–Zn alloys for biomedical applications Yaokun Pan, Siyu He, Diangang Wang ⁎, Danlan Huang, Tingting Zheng, Siqi Wang, Pan Dong, Chuanzhong Chen ⁎ Key Laboratory for Liquid–Solid Structural Evolution & Processing of Materials, Ministry of Education, School of Materials Science and Engineering, Shandong University, Ji'nan 250061, Shandong, PR China

a r t i c l e

i n f o

Article history: Received 11 September 2014 Received in revised form 26 October 2014 Accepted 12 November 2014 Available online 13 November 2014 Keywords: Magnesium alloy Coating Oxidation Electrochemical corrosion Interfaces

a b s t r a c t Calcium phosphate (CaP) ceramic coatings were fabricated on pure magnesium (Mg) and self-designed Mg– 0.6Ca, Mg–0.55Ca–1.74Zn alloys by microarc oxidation (MAO). The coating formation, growth and biomineralization mechanisms were discussed. The coating degradability and bioactivity were evaluated by immersion tests in trishydroxymethyl–aminomethane hydrochloric acid (Tris–HCl) buffer and simulated body fluid (SBF) solutions, respectively. The coatings and corrosion products were characterized by scanning electron microscope (SEM), X-ray diffractometer (XRD), X-ray photoelectron spectrometer (XPS) and fourier transform infrared spectrometer (FT-IR). The electrochemical workstation was used to investigate the electrochemical corrosion behaviors of substrates and coatings. Results showed that Mg–0.55Ca–1.74Zn alloy exhibits the highest mechanical strength and electrochemical corrosion resistance among the three alloys. The MAO-coated Mg–0.55Ca–1.74Zn alloy has the potential to be served as a biodegradable implant. © 2014 Elsevier B.V. All rights reserved.

1. Introduction In recent years, the need for load-bearing materials in temporary implant applications has grown. Over the past few years, the traditional metallic implants (e.g. stainless steels, cobalt–chromium and titanium alloys, and other alloys) used for load-bearing applications are typically biocompatible but not biodegradable [1,2]. Their non-degradability, the mismatched mechanical properties with nature bones, stress shielding effects, second surgical procedure and possible toxic metal ion release may lead to bone loss, higher morbidity and health care costs [2–4]. To avoid these problems, the materials that can provide short-term structural support and then be gradually absorbed and metabolized during the healing process are being sought. The application of Mg alloys as potential biodegradable and bioabsorbable medical implants has gained significant attention in the field of novel metallic implants. The elastic modulus and density of magnesium alloy are closer to those of the cortical bone tissue, so the stress shielding effect can be reduced or even eliminated [3–5]. The mechanical properties of Mgimplants are well matched with the bone tissue in vivo [6]. Besides, magnesium is an essential element for the human body, and its corrosion products generated by in vivo biodegradation can be easily absorbed and metabolized by the human body [7,8]. However, the use ⁎ Corresponding authors at: Jinan 250061, Jing Shi Road # 17923, Shandong, PR China. E-mail addresses: [email protected] (C. Chen), [email protected] (D. Wang).

http://dx.doi.org/10.1016/j.msec.2014.11.048 0928-4931/© 2014 Elsevier B.V. All rights reserved.

of magnesium for biomedical applications is still limited due to their rapid corrosion, which will reduce the mechanical integrity of magnesium before the host tissue is sufficiently healed [3–6]. Moreover, high corrosion rate will lead to unphysiological pH and rapid production of hydrogen gas (Mg + 2H+ → Mg2+ + H2), and then delay tissue healing [9–11]. Nowadays magnesium alloys under investigation as potential biodegradable implant materials were mostly commercial alloys developed for the needs in transportation industry [12], such as AZ series (Mg– Al–Zn, AZ31, AZ31B, AZ91, AZ91D, etc.) [13–15], AM series (Mg–Al– Mn, AM60, etc.) [16,17], AE series (Mg–Al–RE, AE21, LAE442, etc.) [3, 18], WE series (Mg–RE–Zr, WE43, etc.) [3,19,20], ZE series (Mg–Zn– RE, ZE41, etc.) [21], and ZK series (Mg–Zn–Zr, ZK60, ZK61, etc.) [22, 23]. However, almost none of the above mentioned alloys were originally developed to be served as biodegradable implant material. Many aluminum (Al) and (or) heavy metal elements-containing Mg alloys may have latent toxic effects [24,25]. Previous studies showed that aluminum element is a risk factor in the generation of Alzheimer's disease, muscle fiber damage and osteoclast viability decrease [12]. Besides, mild foreign body reactions were observed in vivo during the implantation of AZ91D alloy [26]. Witte et al. [12] recommended that Mg–Al system alloys should be just used as experimental alloys to investigate the improvement of processing and surface modification technology in biomedical applications. Al-free Mg alloys are recommended for the potential use in the human body. Taking these factors into consideration,

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novel Mg alloys with good biocompatibility, high strength and corrosion resistance should be developed. The alloying elements, lithium (Li), calcium (Ca), zinc (Zn), manganese (Mn) and perhaps a very small amount of rare earth (RE) elements with low toxicity, can be used. Currently, many novel Mg alloys have been developed for bone implant applications, such as Mg–Ca, Mg–Zn, Mg–Mn, Mg–Zr, Mg–Zr, Mg–Y, Mg–Zn– Si, Mg–Zn–Mn, and Mg–Zn–Mn–Ca [27–32]. Results showed that Ca and Zn are two suitable alloying elements for the improvement of strength and corrosion resistance of Mg alloys [12,27–32]. However, the anticorrosion and surface morphology of bare Mg alloys are still not ideal. MAO is one of the most prospective surface treatments for Mg alloys. Biocompatible and bioactive ceramic coatings with unique features including inner firmly-adhered, outer porous, and anticorrosive properties can be obtained by MAO technique [2,5,22,23]. In this paper, Ca and Zn were chosen as alloying elements. Two kinds of Mg alloys, Mg–0.6Ca and Mg–0.55Ca–1.74Zn, were developed. CaP ceramic coatings were fabricated on pure Mg, Mg–0.6Ca and Mg– 0.55Ca–1.74Zn alloys in the self-designed electrolyte solution by micro-arc oxidation. The primary objectives of this study include the following three respects: (1) to investigate the effect of Ca and Zn on mechanical properties and corrosion behavior of Mg alloy; (2) to evaluate the in vitro degradation and electrochemical corrosion behavior of Mg alloy and its surface MAO coating; and (3) to discuss the formation and growth mechanism of CaP-MAO coating and coating biomineralization mechanism in SBF solution. 2. Experimental 2.1. Preparation of Mg alloys and MAO coatings In order to investigate the effects of the addition of Ca and Zn on phase composition, microstructure and mechanical property of Mg alloy, pure Mg, Mg–Ca and Mg–Ca–Zn were cast using commercial purity Mg (99.99%) ingot (supplied by Feixian Yinguang Magnesium Co., Ltd. Shandong, China), Ca (99.9%) and Zn (99.9%) blocks (supplied by Beijing Cuibolin Non-ferrous Technology Developing Co., Ltd. Beijing, China) in an electronic resistance furnace under the protection of covering flux. Firstly, Mg ingot was melted at 680 °C, and alloying elements with designed contents (Ca and Zn contents were 0.6 and 1.75 wt.%, respectively) were added to the melt at 700–720 °C, then the temperature was increased to 750 °C, holding and stirring for 15 min to ensure that all alloying elements were dissolved into the melt alloy. Secondly, before casting, the melts were cooled down to 680–700 °C and cast into a graphite mold preheated to 200 °C. Thirdly, the alloy ingots were homogenized at 400 °C for 16 h. The chemical compositions of the as-cast Mg–Ca and Mg–Ca–Zn alloy ingots were analyzed by the energy dispersive spectrum (EDS) in different region with an area of 1 mm × 1 mm, and the final value was an average of five replicate measurements in different regions. For the binary Mg–Ca alloy, the actual Ca content was 0.6 ± 0.08 wt.%. For the ternary Mg–Zn–Ca alloy, the actual Ca and Zn contents were 0.55 ± 0.10 and 1.74 ± 0.23 wt.%, respectively. The alloys were cut into rectangular specimens with the size of 8 × 8 × 6 mm3. The self-designed electrolyte was prepared from a solution containing 0.005 mol/L (C6H5O7)2Ca3·4H2O, 0.005 mol/L Na3PO4, 0.0891 mol/L KOH, 0.1227 mol/L NH4HF2, 0.5% C3H8O3 (volume fraction), 0.5% N (CH2CH2OH)3 (volume fraction) and 0.75% H2O2 (volume fraction) in distilled water. All drugs used were of analytical grade. Prior to coating preparation, each sample was polished up to 1000 grit, cleaned in acetone, absolute ethanol and distilled water respectively, and then dried in warm air. The MAO device consists of a pulse power supply unit, a stirring and a cooling circulating system and a stainless steel container which served as cathode. The specimens were served as anode. The applied positive voltage, negative voltage, frequency, positive duty ratio, negative duty ratio and ratio of positive and negative pulses were 450 V, 0 V, 600 Hz, 30%, 20% and 1:1, respectively. The oxidation was carried out for 10 min. The electrolyte temperature was kept at about

30 °C. The MAO-coated samples were flushed with distilled water after the treatment and dried with a blower. 2.2. Tensile and hardness tests The tensile samples were machined according to ASTM-E8-04 standard [33]. The tensile test was carried out on a WDW-100D electronic universal testing machine at a displacement rate of 1 mm/min at room temperature. Three parallel samples were used for each alloy. The alloy hardness was measured by a DHV-1000 micro-Vikers hardness tester at a load of 100 g for 15 s. The final microhardness value was the average of ten replicate measurements. 2.3. Microstructural characterization X-ray diffractometer (Shimadzu XRD-6100, Japan) with Cu-Kα radiation was used to analyze the phase components of alloys, coatings and their corrosion products after immersion at a scanning speed of 2°/min. Scanning electron microscopes (JEOL JSM-6380LA and Hitachi S3400 N, Japan) equipped with energy-disperse spectrometer attachment, were used to study surface and cross-section morphologies and element distributions of the samples before and after immersion tests. Fourier transform infrared spectroscopy (Bruker Tensor-37, Germany) was used to analyze the phase and structure of the MAO coatings after Tris–HCl and SBF immersion tests in the scanning range of 4000–400 cm−1 at a resolution of 4 cm−1 using the KBr pellet method. The surface chemical composition of MAO coating after SBF immersion test was probed using X-ray photoelectron spectroscopy (PE PHI-5300, USA) with Al Kα (1486.6 eV) radiation. The anode powder of 250 W (12.5 kV, 20 mA) was used. The generated photoelectrons were analyzed with a hemispherical analyzer and the core level XPS spectra for Ca 2p, P 2p and C 1s were measured. Energy calibration was achieved by setting the hydrocarbon C 1s line at 284.6 eV. The data were analyzed with software Xpspeak 4.1. 2.4. Electrochemical corrosion tests Potentiodynamic polarization and electrochemical impedance spectroscopy (EIS) experiments were performed in SBF solution using an electrochemical workstation (Princeton Parstat 2273, USA). All measurements were carried out in a 300 mL beaker using typical threeelectrode cell at 36.5 ± 0.5 °C. Saturated calomel electrode (SCE) was used as reference electrode, a platinum wire was used as counter electrode, and sample with 0.64 cm2 surface area exposed to solution was used as working electrode. Potentiodynamic polarization experiments were carried out at a scan rate of 1 mV/s. The EIS experiments were performed in the frequency range from 105 to 10−2 Hz with alternating current (AC) amplitude of 10 mV. EIS spectra were recorded as a function of different immersion time in SBF solution. The SBF solutions were renewed every 12 h and were kept colorless and deposit-free. All EIS data were analyzed using Zview software. 2.5. Degradation tests For degradability studies, all samples were immersed in Tris–HCl buffer solution (0.05 mol/L) with a pH value of 7.3 at 36.5 ± 0.5 °C. 1 L Tris–HCl buffer solution was prepared by dissolving 6.057 g Tris ((CH2OH)3CNH2) and 43.4 mL HCl (1 mol/L) into distilled water. Triplicate samples were used during the immersion, and the ratio of surface plane area to solution volume was 0.08 cm2/mL. The buffer solution was renewed every other day. At the end of each period of immersion time (6, 12 and 18 days), samples were collected from solution, washed in distilled water and dried to constant weight, and measured on an electronic balance. The weight loss of samples and pH value of Tris–HCl solution was monitored during the immersion test.

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2.6. In vitro bioactivity tests The in vitro bioactivity assessments of all samples were carried out by SBF immersion test, proposed by Kokubo et al. [34]. SBF solution has similar composition and concentration to those of the inorganic part of human plasma [34]. 1 L SBF solution was prepared by dissolving the reagents in the following content and order: 7.996 g NaCl, 0.350 g NaHCO3, 0.224 g KCl, 0.228 g K2HPO4·3H2O, 0.305 g MgCl2·6H2O, 40 mL HCl (1 mol/L), 0.278 g CaCl2 and 0.071 g Na2SO4 in distilled water at 36.5 ± 0.5 °C, then the pH was adjusted to 7.3 with Tris and HCl at 36.5 °C. The detailed ion concentration of SBF in this paper can be referred to the reference [23]. Triplicate samples were used during the immersion, and the ratio of surface plane area to solution volume was 0.08 cm2/mL. The SBF solution was renewed every other day and kept colorless and deposit-free during immersion test. Samples were collected from solution after 6, 12, and 18 days immersion, respectively, cleaned in a chromic solution (200 g/L CrO3 + 10 g/L AgNO3) to remove the corrosion products, then washed with distilled water and air-dried. The weight loss of samples and the pH value of SBF solution were monitored during the immersion test. 3. Results 3.1. Optical microstructure and mechanical properties of alloys The optical microstructure and elemental mapping of alloys are shown in Fig. 1. It can be observed that the grain size decreases with the addition of 0.6 wt.% Ca in pure Mg (Fig. 1(a, b)), and decreases further with the addition of 0.55 wt.% Ca and 1.74 wt.% Zn in pure Mg (Fig. 1(c)). The refining effect of Ca and Zn in pure Mg is proved. The elemental mapping of Mg–0.55Ca–1.74Zn alloy in Fig. 1(d) shows the element distribution of Mg, Ca and Zn, and we observe that there are relatively high Ca and Zn contents in grain boundaries. It indicates that Ca- and Zn-containing phases precipitate in grain boundaries. Microhardness and tensile properties of the alloys are listed in Table 1. With the addition of Ca and Zn, microhardness and tensile strength of Mg–0.55Ca–1.74Zn alloy are improved significantly. 3.2. Morphology and elemental composition Surface morphologies and elemental compositions of MAO coatings before and after SBF immersion are shown in Fig. 2. The coating surfaces are porous and the total sizes of micropores are about 5–15 μm. The pore size and distribution of the coating on pure Mg are relatively homogeneous in Fig. 2(a1). The surfaces of coatings on Mg–0.6Ca and

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Table 1 Microhardness and mechanical properties of the as-cast pure Mg, Mg–0.6Ca and Mg– 0.55Ca–1.74Zn alloys. Alloys

Microhardness (HV0.1) Xa ± SDa

Ultimate strength (MPa) Xb ± SDb

Elongation (%) Xb ± SDb

Pure Mg Mg–0.6Ca Mg–0.55Ca–1.74Zn

28.8 ± 0.6 44.6 ± 1.1 49.1 ± 0.9

58 ± 1.8 116 ± 2.6 148 ± 3.8

7.2 ± 0.8 8.7 ± 1.1 10.6 ± 1.5

1 a X in Xa ± SDa refers to the mean value of ten replicate measurements, and SDa refers to the standard deviation. 2 b X in Xb ± SDb refers to the mean value of three replicate measurements, and SDb refers to the standard deviation.

Mg–0.55Ca–1.74Zn are rough in Fig. 2(b1) and (c1). As shown in Fig. 2 (a2), (b2) and (c2), the porous morphologies of the MAO coatings are not evident and numerous spherical and blocky particles with different size are found after immersion in SBF for 18 days. According to EDS results, these particles have an elemental composition of C, Mg, F, O, Ca and P. The element Zn is detected in Fig. 2 (c2). The cross-section image and element distribution of the coatings are shown in Fig. 3(a), (b) and (c). There is no apparent discontinuity in the alloy/coating interface in Fig. 3(a) and (c). Visible microcrack in bonding zone is observed in Fig. 3(b). It indicates that the interface bond is relatively weak for Mg–0.6Ca alloy. The thickness of coatings on pure Mg, Mg–0.6Ca and Mg–0.55Ca–1.74Zn are about 35, 37, 43 μm, respectively. The elements of Mg, F, O, Ca and P are detected in the coating. 3.3. Phase analysis before and after immersion tests 3.3.1. XRD analysis XRD patterns of bare alloys and coatings before and after immersion in Tris–HCl and SBF solution for 18 days are shown in Fig. 4. From Fig. 4 (a), α-Mg and Mg2Ca phases are observed for Mg–0.6Ca alloy. Two more phases, Ca2Mg6Zn3 and Ca2Mg5Zn13, are observed for Mg– 0.55Ca–1.74Zn alloy. From Fig. 4(b), it can be observed that the MAO coatings are mainly composed of metal oxides (MO, M = Mg, Ca, Zn) and fluorides (MF2, M = Mg, Ca, Zn), besides, β-tricalcium phosphatein (β-TCP, β-Ca3(PO4)2) and calcium pyrophosphate (CPP, Ca2P2O7) are detected in coatings. Owing to the porous structure and thin thickness of the coatings, the peak intensity of Mg is still strong. As to Tris–HCl immersion test, for the alloys in Tris–HCl in Fig. 4(c), the corrosion products for each kind of alloy are mainly hydroxides (M (OH)2, M = Mg, Ca, Zn) and oxides, besides, basic magnesium chloride (MgxCly(OH)2x − y·zH2O) and magnesium carbonate hydroxide (Mg5

Fig. 1. Optical microstructure and elemental mapping of the as-cast alloys: (a) Pure Mg; (b) Mg–0.6Ca; (c) Mg–0.55Ca–1.74Zn; (d) elemental mapping of Mg–0.55Ca–1.74Zn.

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Fig. 2. Surface morphologies and elemental composition of the MAO coatings before (a1, b1, c1) and after (a2, b2, c2) immersion in SBF for 18 days: (a1–a2) Pure Mg; (b1–b2) Mg–0.6Ca; (c1–c2) Mg–0.55Ca–1.74Zn.

(CO3)4(OH)2) are observed. For the coatings, as shown in Fig. 4(d), βTCP and CPP disappear. The pre-existing phases before immersion can be observed, but their peak intensities become weak. No MgxCly(OH)2x − y · zH2O is found in corrosion products. As to SBF immersion test, for the alloys in Fig. 4(e), their metal hydroxides are main corrosion products. Moreover, hydroxyapatite (HA, Ca10(PO4)6(OH)2), β-TCP and CPP phases are found on alloys surfaces. For the MAO coatings in Fig. 4(f), their pre-existing phases are noticeable. However, no obvious hydroxides corrosion products are detected. What is more, CaP phases, HA, β-TCP, CPP and carbonated hydroxyapatite (CHA, Ca10(PO4)6 − 2x(CO3)2x(OH)2) are found. Besides, the presence of amorphous region below 20° (2θ) indicates the formation of amorphous phases.

3.3.2. FT-IR analysis The FT-IR spectra of Mg–0.55Ca–1.74Zn alloy and its surface coating after immersion in Tris–HCl and SBF solutions for 18 days are shown in Fig. 5. It provides more information about amorphous phases and other components of precipitates after in vitro immersion. As shown in Fig. 5, the strong, sharp band at 3698 cm−1 is the characteristic frequency of the O–H stretching vibration in Mg(OH)2, while the weak band at 3644 cm−1 in Fig. 5(a-2) is due to the presence of Ca(OH)2 [35]. The strong, broad band at 3390–3428 cm−1 is assigned to the stretching absorption band of free water and the medium band at 1625–1646 cm−1 is attributed to the H–O–H bending mode of crystal water [36]. The split bands at 1480, 1433 cm−1 (Fig. 5(a-1)), 1480, 1440 cm−1 (Fig. 5(a-2)), 1438, 1511 cm−1 (Fig. 5(b-1)) and 1456, 1530 cm−1 (Fig. 5(b-2)) are

Fig. 3. Cross-section morphologies and elemental composition of the MAO coatings before immersion: (a) Pure Mg; (b) Mg–0.6Ca; (c) Mg–0.55Ca–1.74Zn.

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Fig. 4. XRD patterns of bare alloys and CaP coatings before (a, b) and after (c, d, e, f) immersion in Tris–HCl (c, d) and SBF (e, f) solutions for 18 days: (a, c, e) bare alloys; (b, d, f) CaP coatings.

ascribed to the v3 antisymmetric stretching mode of carbonate (CO2− 3 ) specie. Bands v2 (out-of-plane deformation) and v4 (in-plane bending) of CO2− are recorded at approx. 857–861 and 745–756 cm−1, respec3 tively. The vibration (2346 cm−1) is assigned to v3 antisymmetric stretching mode of CO2 gas molecule in the air [36]. The shoulder at approx. 934–937 cm−1 is ascribed to the presence of hydroxyl groups (bending vibration, δOH) [36,37]. Other bands in the range of

450–570 cm− 1 in Fig. 5(a) and (b) are due to the presence of M–O or M–F (M = Mg, Ca, Zn) [38]. Moreover, the main difference between Fig. 5(a) and (b) is the presence of bands at 1045 and 1011–1120 cm− 1 in Fig. 5(b). The band at 1045 cm−1 is assigned to the v3 asymmetric stretching mode of PO34 − while the band at 1011–1120 cm− 1 indicates the presence of PO34 −, HPO2− and (or) 4 P2O4− species [36,39]. 7

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Fig. 6. XPS analysis of MAO-coated coating on Mg–0.55Ca–1.74Zn alloy after immersion in SBF solution for 18 days: (a) high-resolution spectrum of Ca 2p; (b) high-resolution spectrum of P 2p.

Fig. 5. FT-IR spectra of bare Mg–0.55Ca–1.74Zn alloy and MAO-coated coating after 18 days immersion in Tris–HCl (a) and SBF (b) solutions: (1) bare alloys; (2) CaP coating.

3.3.3. XPS analysis The chemical states of Ca and P were investigated by XPS. Fig. 6 displays the XPS spectra of MAO-coated CaP coating on Mg–0.55Ca–1.74Zn alloy after immersion in SBF solution for 18 days. High-resolution X-ray photoelectron spectra are acquired for Ca 2p and P 2p. The broad peaks of Ca 2p and P 2p indicate the presence of different Ca- and P-containing species. This may be the reason for the absence of Ca 2p spectra with doublet feature. High-resolution spectrum of Ca 2p splits into five peaks at 351.1 (2p1/2), 349.6 (2p3/2), 348.4 (2p), 347.3 (2p3/2) and 345.0 eV (Fig. 6(a)), respectively. The Ca 2p3/2 at 347.3 eV is the satellite peak and the others are energy peaks. These peak positions correspond to a divalent oxidation state (Ca2+) in inorganic Ca-containing compounds. As shown in Fig. 6(a), the binding energy of Ca 2p3/2 peaks at 347.3 eV can be attributed to the presence of Ca3(PO4)2, CaO, Ca(OH)2, CaCO3, CaHPO4 and Ca10(PO4)6(OH)2. The other binding energies of Ca peaks at 351.1, 349.6 and 348.4 eV can correspond to CaCO3 (2p1/2), CaF2 (2p3/2), CaF2 (2p), respectively [40–43]. The P 2p spectrum shows five peaks at 135.5 (2p), 134.3 (2p), 133.0 (2p), 131.4 (2p) and 130.0 eV (Fig. 6(b)), respectively. The satellite peak at 133.0 eV is attributed to the presence of Ca10(PO4)6(OH)2 (P 2p) and Ca3(PO4)2 (P 2p) [41,44]. The peaks at 135.5 and 134.3 eV are

corresponding to Na3PO4 (P 2p) and CaHPO4 (P 2p3/2), respectively [45–47], which is coincident with the analysis of Ca 2p in Fig. 6(a). On the basis of published data [44–48], it is suggested that complex and amorphous CaP compounds are formed during the MAO or (and) immersion process, which contributes to the presence of Ca 2p peak at 345.0 eV (Fig. 6(a)) and P 2p peaks at 130.0 and 131.4 eV (Fig. 6(b)). The results of XPS are basically coincident with XRD and FT-IR analysis. It can be concluded that the particles in Fig. 2(a2), (b2) and (c2) are actually calcium phosphates. 3.4. Electrochemical corrosion behaviors Typical potentiodynamic polarization curves of alloys and coatings are shown in Fig. 7. The corrosion potential (Ecorr) and corrosion current density (Icorr) are also given in the figure. Compared with the Ecorr and Icorr values of pure Mg and Mg–0.6Ca alloys, the Ecorr of Mg–0.55Ca– 1.74Zn alloy is higher and the Icorr is lower, which indicates that the corrosion resistance of pure Mg alloy is enhanced by the addition of Ca and Zn. Compared with the Ecorr and Icorr values of the alloys, the coating Ecorr of the coating shifts to the positive direction with a varying degree, and the coating Icorr decreases by 40–115-fold. Results show that the corrosion resistance of alloys is improved significantly by MAO coatings. The MAO-coated Mg–0.55Ca–1.74Zn exhibits the lowest corrosion current density.

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Fig. 7. Potentiodynamic polarization curves of the bare (a) and MAO-coated (b) samples in SBF solution: (1) pure Mg; (2) Mg–0.6Ca; (3) Mg–0.55Ca–1.74Zn.

EIS spectra of the bare and MAO-coated Mg–0.55Ca–1.74Zn in SBF solution as a function of immersion time are shown in Fig. 8. For the bare alloy exposed in SBF for 0, 6, 12, 24 h and the coating before SBF immersion, one high frequency capacitive loop and one low frequency capacitive loop, are observed in Fig. 8(a) and (b). Their experimental data can be accurately fitted to the equivalent circuit (EC-1): Rs(Qc(Rp(RtQdl)))Lhigh (as shown in Fig. 9(a)). Besides, for the

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alloy exposed in SBF for 6 h, a small inductive loop in low (Llow) frequency region is observed, as the exposure prolongs to 12 and 24 h, the inductive loop Llow becomes invisible. For the coating exposed in SBF for 6, 24 and 48 h, as shown in Fig. 8(c) and the partial enlargement in high frequency region, one high frequency capacitive loop, one medium frequency capacitive loop and one low frequency capacitive loop, are observed. The small inductive loop in high and low frequency regions disappears. Their experimental data can be fitted to the following EC (EC-2): Rs(Qc(Rpo(Qf(Rf(RtQdl))))) (as shown in Fig. 9(b)). In the above two equivalent circuits, Rs represents the solution resistance between the reference electrode and working electrode. Qc, one of the constant phase element (CPE) components, is related to the capacitance of the corrosion products layer. Rp is the polarization resistance of the alloy. Rt is the charge transfer resistance relating to the electrochemical reaction. Qdl, another CPE component, corresponds to the capacitance of the double layer between alloy and the electrolyte solution. Rpo is the relevant resistance relating to micropore or ionic conducting defect resistance. Qf, the third CPE component, represents the capacitance of the coating. Rf represents the resistance of the coating. The presence of CPE can be explained by dispersion effect caused by microscopic inhomogeneity and roughness of electrode surface [49]. The source of element Lhigh in EC-1 is not exactly clear, according to the published data, it is possibly resulting from wiring and measuring equipment components [50,51]. The presence of the inductive loop Llow is probably resulting from intense microgalvanic corrosion in the boundaries between the matrix and second phases of the alloy at the initial immersion period, and with the accumulation of the corrosion products, this galvaniccouple effect is depressed, and Llow becomes invisible with immersion proceeding. The high frequency capacitive loop is generally associated with the electrolyte penetration, charge transfer and the formation and alteration of corrosion products layers [49]. The low frequency region in EIS contains important information on the electrode controlled process together with the contribution from localized defects to the overall impedance [52]. The capacitive loop in low frequency in this study is probably attributed to changes in corrosion product layer and corrosion mode. The fitting results of EIS spectra for EC-1 and EC-2 are listed in Tables 2 and 3, respectively. For bare alloy in Table 2, after 6 h immersion in SBF, Rp and Rt increase to some extent, which indicates the formation of partially covered corrosion product layer on sample surface. As the exposure prolongs to 12 h, Rp and Rt decrease slightly, but still higher than those of the bare alloys before immersion. It implies the dissolution of corrosion products and continuous corrosion of the alloy substrate. When the exposure prolongs to 24 h, both Rp and Rt increase significantly, owing to the formation of protective corrosion products layer. For the MAO coating in Tables 2 and 3, it shows that the MAOcoated sample before immersion exhibits high Rp and Rt. The MAO coating can effectively improve the electrochemical corrosion resistance of the bare alloy. However, in the following immersion in SBF for 6 h, a corroded layer forms on coating surface, and the corrosion mode changes. A four-phases–three-interfaces corrosion system (SBF solution/corrosion products layer/outer porous layer of MAO coating/inner dense layer of MAO coating) is formed. The protective capability of corrosion products and coating porous layers is limited. The decrease of Rf and Rt is probably due to the penetration of electrolyte into the inner dense layer of MAO coating through micropores and cracks. As the immersion prolongs to 24 h, Rpo, Rf and Rt increase to some extent, which is due to the formation of large amount of corrosion products. Micropores and cracks are covered by these corrosion products. At 48 h, Rpo, Rf and Rt decrease again, which implies the dissolution of corrosion products and continuous corrosion of MAO coating. It can be observed that Rp Rpo, Rf and Rt of substrate or coating in SBF vary dramatically with elongated immersion time. The above results prove that corrosion products and protective layer can be formed on Mg–0.55Ca–1.74Zn alloy and its surface coating during SBF immersion process. The formation and dissolution of corrosion products is a dynamic process.

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Fig. 8. EIS spectra of bare (a) and MAO-coated (b, c) Mg–0.55Ca–1.74Zn samples as a function of immersion time in SBF solution.

3.5. In vitro degradation in Tris–HCl and SBF Fig. 10 shows pH value, weight loss percentage changes in Tris–HCl and SBF solutions. From Fig. 10(a) and (c), it can be observed that the pH value increases with prolonging exposure to Tris–HCl and SBF solutions in one immersion cycle. For each sample with the same immersion time, the pH value of the Tris–HCl solution is higher than that of the SBF solution. After the same exposure time, the pH value of the SBF solution in which the bare sample was immersed is relatively higher than that of the SBF solution in which the MAO-coated sample was immersed. In Tris–HCl solution, except for the solution in which the Mg–0.55Ca– 1.74Zn alloy was immersed, the others exhibit a similar tendency to that in SBF solution. The pH values of the SBF solutions in which the bare samples were immersed fluctuate obviously, while the pH values of the SBF solutions in which the MAO-coated samples were immersed are relatively smooth. The MAO-coated Mg–0.6Ca exhibits lower but fluctuant pH values in Tris–HCl solution, while the MAO-coated Mg– 0.55Ca–1.74Zn exhibits lower and stable pH values in SBF solution, almost equal to that of MAO-coated pure Mg. The MAO-coated pure Mg exhibits the lowest pH values in two solutions. From Fig. 10(b) and (d), it can be observed that the weight loss increases with prolonging immersion in Tris–HCl and SBF solutions. As shown in Fig. 10(b), the MAO-coated Mg–0.6Ca and Mg–0.55Ca–1.74Zn show higher weight loss percentage than that of the bare alloy, which is probably due to the massive spalling of MAO coating when the Tris–HCl solution penetrates the coating to the alloy substrate. It indicates that the MAO coating exhibits inadequate corrosion protection in this kind of corrosive

Fig. 9. The equivalent circuit: (a) EC-1; (b) EC-2.

buffer solution. However, as shown in Fig. 10(d), the MAO-coated pure Mg, Mg–0.6Ca and Mg–0.55Ca–1.74Zn show lower weight loss percentage than that of their bare alloys, especially for the MAOcoated Mg–0.55Ca–1.74Zn, which presents the lowest weight loss percentage (3.19%) after 18 days immersion in SBF. It can be observed that the MAO coating provides the most effective protection for Mg– 0.55Ca–1.74Zn alloy. 4. Discussion Studies have indicated that calcium is an effective grain growth inhibitor in Mg alloy [53]. Laves phase Mg2Ca in the binary Mg–0.6Ca alloy is formed and the addition of Ca refines the grain. Zn can hinder the movement of the recrystallized grain boundaries and the grain size can be refined [31,32]. Besides, Ca and Zn contribute to solid solution and precipitation strengthening. Moreover, Ca is the most abundant mineral element in the human body (about 1–1.1 kg) and mainly stored in bone and teeth. Its normal serum level is 0.919–0.993 mg/L [12]. Zn is also an essential element in the human body, which is essential for the immune system. It is a co-factor for specific enzymes in bone and cartilage [54]. However, it is neurotoxic at higher concentration. Its normal blood serum level is 12.4–17.4 μmol/L [55]. Adequate amounts of Ca and Zn elements can be added to pure Mg to improve the biocompatibility and mechanical property of Mg–Ca–Zn alloys. MAO coating can be fabricated on Mg alloy to improve its corrosion resistance and surface morphology. Fig. 11 shows the voltage–, current– time response of MAO process and surface morphologies of different films formed with different treatment time. As shown in Fig. 11(a), (b) and (c), with the increase of voltage, the current increases firstly, then decreases slightly, and subsequently increases to the highest value, and drop to zero at last. Corresponding to the current changes, the formed films show different morphologies. In the initial stage (0–27 s), the Mg alloy was firstly dissolved and cations (Mg2+, Ca2 +, Zn2 +) were released. A very thin and glossy oxide film was formed on the substrate surface (Fig. 11(c)—20s) and began to grow (Fig. 11(c)—27 s). The kinetics of the electrode processes at low voltage conforms to Faraday's and Ohm's laws [56]. Thus, the increase in voltage led to a proportional rise in the current in this stage. In the following stage (27–31 s), oxygen (4OH− − 4e → O2 + 2H2O) gas was adsorbed on the preformed film, accumulating in the positions of defect parts. The current rise was limited by this partial shielding action. In the other areas where the electrode remained in contact with the electrolyte solution, the current density continued to rise, leading to local boiling of the electrolyte adjacent to the electrode. When the voltage reached to the value of 31 s, the sample surface was enshrouded by a continuous gaseous envelope of low electrical conductivity. So the total current dropped slightly. At the same time, this gaseous envelope provided

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Table 2 Fitting results of EIS spectra for the bare alloy and MAO coating before immersion (fitting to EC-1). t (h)

Rs (Ω·cm2)

Rp (Ω·cm2)

Qc (Ω−1·cm−2·Sn)

0 6 12 24 0 (MAO)

21.2 18.0 21.3 17.4 21.0

50.7 140.4 100.5 204.0 1.48 × 104

3.169 6.199 5.039 4.859 1.100

× × × × ×

10−5 10−5 10−5 10−5 10−6

nc

Rt (Ω·cm2)

Qdl (Ω−1·cm−2·Sn)

0.8378 0.7954 0.8144 0.8254 0.7466

15.0 40.5 38.3 88.2 3739.0

6.868 5.854 8.318 3.571 7.952

× × × × ×

10−3 10−3 10−3 10−3 10−6

ndl 0.8164 0.9422 0.8255 0.8069 0.7011

Table 3 Fitting results of EIS spectra for the MAO coating after SBF immersion (fitting to EC-2). t, h 6 24 48

Rs, Ω·cm2 26.2 26.4 25.6

Rpo, Ω·cm2 60.0 185.2 79.1

Qc, Ω−1·cm−2·Sn −6

2.763 × 10 6.657 × 10−6 8.642 × 10−6

nc 0.6426 0.6569 0.6325

Rf, Ω·cm2 1105 2233 2097

favorable conditions for the formation of plasma. Additionally, the high electric field strength (106 and 108 V/m) in the near-electrode region leads to the ionization processes of plasma [56]. The weakest areas with the lowest resistance were broken down preferentially, accompanied by spark discharge phenomenon. The gases were pulled out because of the high temperature and pressure, so a lot of micropores penetrating to substrate were left on the oxide coating surface (Fig. 11 (c)—31 s). So in this stage (31–35 s), the current increased sharply.

Qf, Ω−1·cm−2·Sn −5

1.652 × 10 1.257 × 10−5 1.300 × 10−5

nf

Rt, Ω·cm2

Qdl, Ω−1·cm−2·Sn

ndl

0.8752 0.9352 0.8568

304.3 569.7 524.8

4.300 × 10−3 2.289 × 10−3 2.369 × 10−3

0.9986 0.9352 0.8966

The micropores became subsequent discharge channels. With further increase of voltage, molten ceramic phases were formed by migrated cations (Mg2 +, Ca2 +, Zn2 +, etc.) and anions (O2 −, F−, PO34 −, etc.) under the influence of the electric field and ion concentration gradient. These molten oxides condensed rapidly by exposure to the cool electrolyte. Then the oxides formed in different discharge areas combined together gradually. The aforementioned process continued and the coating began to grow until the pre-set voltage could not break down

Fig. 10. pH value (a, c) and weight loss percentage (b, d) changes in Tris–HCl (a, b) and SBF (c, d) solutions as a function of immersion time. Error bars represent standard error of mean (N = 5).

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Fig. 11. Voltage–, current–time response of MAO process and surface morphologies of different films formed with different treatment time (a) Pure Mg; (b) Mg–0.6Ca; (c) Mg–0.55Ca– 1.74Zn.

the coating any more. The current dropped to zero at last. Because much more energy was required to break down the preformed coating, the coating surface became more porous and rough with the prolongation of treatment time (Fig. 11(c) 35–60 s). The MAO coating formation and growth process is a complex process relating to chemical, electrochemical, thermochemical and plasma reactions. The outer porous structure may be valuable as a depot for growth factors or bone morphogenic proteins (BMP), and can provide interfacial fixation between the implant and tissue [5]. The MAO coatings on Mg–0.6Ca and Mg– 0.55Ca–1.74Zn alloys are uniform, and without visible defects. There is a big pit on the MAO coating surface of pure Mg in Fig. 11(a), which is caused by continuous partial discharges. Owing to the extremely high thermal effect, the elements in electrolyte and alloy can be incorporated into coating through MAO reactions. Many ceramic phases, such as MgO, MgF2, CaO, CaF2, ZnO, ZnF2, β-TCP and CPP and some amorphous CaP compounds can be formed in the coating. It is reported that β-TCP and CPP are biocompatible and bioresorbable materials [57]. Tris and HCl is an acid–base pair in the Tris–HCl buffer solution. A large amount of chloride (Cl−) ions is also involved in this solution. When the bare alloys were immersed in the solution, the acid condition in the metal/solution interface caused partial dissolution of the alloys. Mg2+, Ca2+ and (or) Zn2+ ions were produced, accompanied by hydrogen evolution. The pH value increased at the same time. Then metal ions would react with OH− to form Mg(OH)2, Ca(OH)2 and Zn(OH)2. Cl− ions accelerated the reaction by reacting with Mg(OH)2 and forming a relatively soluble MgCl2, MgCl(OH), etc. Moreover, CO2 in the air would dissolve into the solution and CO23 − ions were formed. Mg5 (CO3)4(OH)2 phase was formed by the following reaction:

− 5Mg2+(aq) + 4CO2− 3 (aq) + 2OH (aq) → Mg5(CO3)4(OH)2.

The XRD and FT-IR results clearly show the presence of CO2− 3 ions in corrosion products. Previous studies showed that the surface with thick magnesium hydroxy carbonate film was passive [58]. The effect of CO2 on the atmospheric corrosion behavior of Mg and its alloys should be considered. The electrochemical corrosion resistance of pure Mg is improved by the addition of 0.55 wt.% Ca and 1.74 wt.% Zn in pure Mg, which is due to the formation of a compact passive film. In addition, it has been widely accepted that Zn can reduce the effects of iron (Fe), nickel (Ni) on corrosion and then improve the anticorrosion ability of pure Mg [30–32]. The coating is a barrier to hold back the corrosive chloride ions. The anticorrosion abilities of Mg alloy can be improved by the MAO coating owing to the blocking effect in SBF solution. The porosity and microcracks are two important factors affecting the corrosion behavior of MAO coating. During the SBF immersion process, when the bare and MAO-coated samples were immersed in SBF solution, the degradation occurred. The alloy and coating degradability was influenced by the dissolution of the alloy and coating, and the formation of corrosion products and CaP apatites. The formation and growth of CaP apatites on sample surface could compensate the weight loss of alloys and coating dissolution. In the present work, the biomineralization mechanism of MAO-coated Mg alloy was discussed. The schematic diagram of the formation of CaP apatites on the surface of CaP-MAO coating in SBF is shown in Fig. 12. The apatite formation ability is related to ions exchange in SBF solution. SBF is a supersaturated solution with Ca2+ and HPO2− ions. 4

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Fig. 12. Schematic diagram of the apatites formation on CaP-MAO coating surface in SBF.

As shown in Fig. 12, In the initial stage of SBF immersion, many phases in the coating, such as β-TCP, CPP, MgO, MgF2, CaO and CaF2 reacted with H+ ions to release Mg2+, Ca2+, F− and PO3− 4 ions into the solution. A lot of corrode pits formed in the coating/solution interface with the help of corrosive Cl− ions. The consumption of H+ ions caused an increase of − 3− pH value. Thus, Ca2+, HPO2− 4 , PO4 , OH and the other ions concentrated in the pitting/solution interfaces. The local saturation of Ca2+, PO3− 4 and OH− ions increased in this alkaline solution, at the same time, rough areas with large specific surface areas, such as micropores, corrosion pits and products, provided favorable nucleation sites for HA and other apatites. Some amorphous CaP phases could be formed. Once the nuclei formed, they grew gradually at the expense of Ca2+, HPO2− 4 and OH− ions absorbed in apatites/solution interface. Moreover, CO2− 3 ions in SBF solution could occupy the monovalent anionic (OH−) sites (designated as type A) or trivalent anionic (PO3− 4 ) sites (designated as type B) in HA to form carbonate hydroxyapatite (CHA) [59]. F− ions could replace the OH− sites in HA to form fluor-hydroxy-apatite (FHA). When these apatites grew to some degree, they began to dissolve into the solution again. The dissolution and precipitation of apatites is a reversible reaction. The movement of dissolution–precipitation balance depends on the coating corrosion resistance and apatites forming ability.

Zn–Ca alloys. The addition of Ca and Zn elements in pure Mg with proper content is favorable for the improvement of electrochemical corrosion resistance. (2) Bioactive CaP coatings with outer porous and inner dense layers are fabricated on pure Mg, Mg–0.6Ca and Mg–0.55Ca–1.74Zn alloys by MAO in Ca-, P-containing electrolytes. The electrochemical corrosion resistance of the alloys is improved by MAO coating. The porosity and microcracks of the coating are two important factors affecting the coating electrochemical corrosion behavior. (3) The in vitro Tris–HCl immersion assessments show that the pure Mg exhibits the best corrosion resistance, but its surface MAO coatings do not provide effective protection. However, the corrosion resistance of MAO-coated samples is significantly better than that of bare alloys in SBF immersion tests. The MAO coating provides the most effective protection for Mg–0.55Ca–1.74Zn alloy. (4) MAO coating formation and growth process is a complex process relating to chemical, electrochemical, thermochemical and plasma reaction. The dissolution and precipitation of apatites on MAO coating is a reversible reaction. The movement of dissolution–precipitation balance depends on the coating corrosion resistance and apatites forming ability.

5. Conclusion Element alloying and surface modification are taken into consideration to develop new Al-free Mg alloys and to improve the surface morphology and corrosion resistance of the alloys further. The following conclusions can be drawn: (1) Ca and Zn are effective grain refiners in pure Mg. Adequate amounts of Ca and Zn elements can be added to pure Mg to improve the biocompatibility and mechanical properties of Mg–

Acknowledgments This work is financially supported by two projects. One is the Science and Technology Supporting System Item of Resource-Conserving Society of Shandong Province (Grant No. 2007JY05) and the other is the Development Project of Science and Technology (Grant No. 2010GSF10627) of Shandong Province.

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