Improvement of the Biodegradation Property and Biomineralization

Article pubs.acs.org/journal/abseba

Improvement of the Biodegradation Property and Biomineralization Ability of Magnesium−Hydroxyapatite Composites with Dicalcium Phosphate Dihydrate and Hydroxyapatite Coatings Yingchao Su,† Dayong Li,† Yichang Su,† Chengjia Lu,† Liyuan Niu,‡ Jianshe Lian,† and Guangyu Li*,† †

Key Laboratory of Automobile Materials, Ministry of Education, College of Materials Science and Engineering, Jilin University, 5988 Renmin Street, Changchun 130025, China ‡ Department of Material Engineer, Zhejiang Industry & Trade Vocational College,717 Fudong Street, Wenzhou 325003, China ABSTRACT: The application of calcium phosphate reinforced magnesium matrix composites has not achieved the expected effect to control the degradation rate of magnesium so far. Therefore, in order to enhance the corrosion resistance and further develop the surface bioactivity of the composites to meet specific requirements of bone tissue engineering applications, biocompatible dicalcium phosphate dihydrate (DCPD) and hydroxyapatite (HA) coatings have been deposited on homemade HA/Mg composites using a simple conversion coating method and a subsequent alkali posttreatment, respectively. The conversion coating mechanism was studied by comparing coating processes on the composites, pure Mg, and an AZ60 Mg alloy. Electrochemical results showed that polarization resistance of the optimum DCPD and HA coatings was about 15 and 65 times higher than that of the composites, respectively. Immersion tests in simulated body fluid revealed that both coatings could supply improved corrosion resistance and biomineralization ability for the HA/Mg composites. KEYWORDS: magnesium composites, calcium phosphate coating, biodegradable property, biomineralization

1. INTRODUCTION Magnesium (Mg) and its alloys have been considered as potential alternatives to conventional orthopedic implant materials because of their attractive mechanical and biodegradation properties.1−3Their physical and mechanical properties are similar to those of bones, which would decrease the incidence of stress shielding.3,4 In addition, their degradation properties can help to avoid the second removal surgery, and recently a long-term clinical study found that the controlled degradation of an MgZnCa alloy resulted in the formation of a “biomimicking calcification matrix” at the degrading interface to initiate the bone formation process.5 However, the degradable property is a double-edged sword. The extremely high degradation rate of the Mg-alloy implant in bodily fluid not only deteriorates its mechanical integrity before the sufficient tissue healing, which is especially important in the repair or replacement of damaged load-bearing bones, but also releases too much hydrogen gas, which results in subcutaneous bubbles6 and potentially delays the healing of the surgical region. It was anticipated that the application of calcium phosphate (CaP) reinforced Mg metal matrix composites (MMCs) could be an approach to control the in vivo corrosion rate of the Mg implant while not compromising the mechanical properties, or possibly even enhancing them.7 It has been found that the MMCs generally exhibit improved compression strength but reduced tensile strength and elongation than the master alloys, © XXXX American Chemical Society

whereas the corrosion resistance of MMCs could be either better or worse than the master alloy.8,9 CaP-based ceramic particles or scaffold can stabilize the corrosion rate of the composites and the surface can exhibit more uniform corrosion attack mainly due to the formation of a uniform passive layer on the MMCs surface.7,10−12 However, galvanic corrosion can easily occur along the particulate−matrix interface, especially in the early stage of corrosion, swelling the composites and allowing the electrolyte to penetrate deeper,13 and thus increasing the corrosion rate of the composites.14,15 It has been shown that producing protective bioactive ceramic coatings on Mg and its alloys can enhance their biocompatibility and slow down their corrosion rate in physiological environments.16−18 CaPs, such as dicalcium phosphate dihydrate (DCPD), tricalcium phosphate (TCP) and hydroxyapatite (HA), all have intrinsic bioactivity and biocompatibility because Ca and P are the main elements in bone minerals.19−21 Among these CaPs, HA is a preferred coating material, because it bears a chemical and structural resemblance to natural bone and can accelerate bone concrescence.22 In our previous work, DCPD and HA coatings were prepared on Mg−Al based alloys using a simple Received: January 10, 2016 Accepted: March 30, 2016

A

DOI: 10.1021/acsbiomaterials.6b00013 ACS Biomater. Sci. Eng. XXXX, XXX, XXX−XXX

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ACS Biomaterials Science & Engineering conversion coating technology.23,24 With the advantage of convenient operation to produce a uniform and well adhered coating, especially for the complex-shaped components of the orthopedic implant, this approach can be usefully extended to the Mg composites in order to enhance their corrosion resistance and surface biocompatibility. Equally, it is known that the second phase plays a critical role during the conversion coating process because of the different electrochemical potentials between the α-Mg phase and β phase in Mg alloys.23−26 The electrochemical heterogeneity also exists in the Mg composites, even with more anodic sites forming galvanic coupling during the corrosion process.14,15 However, this could possibly turn into a positive factor in the electrochemical reactions when producing the conversion coating on the Mg composites. Therefore, in the present study, a DCPD conversion coating was deposited on the surface of a homemade HA/Mg composite, and the coating deposition process was analyzed. The DCPD coatings with different deposition times were then transformed into different HA coatings using an alkali post treatment. The effects of the conversion coating time on the coating morphology and electrochemical polarization behavior were analyzed and then optimum DCPD and HA coatings were obtained. The electrochemical and immersion corrosion behaviors of the optimum coatings in simulated body fluid (SBF) were also systematically investigated.

Figure 1. (a) SEM image and (b−d) EDS elemental maps of the composites: (b) Mg, (c) Ca, and (d) P. MgSO4·H2O, 0.1 g/L KH2PO4·H2O and 0.06 g/L Na2HPO4·H2O, and its pH 7.4.27 A classical three-electrode cell was used with platinum as the counter electrode, a saturated calomel electrode (SCE, + 0.242 V vs SHE) as the reference electrode, and the sample as the working electrode (with 0.25 cm2 exposed area). Prior to the Electrochemical Impedance Spectroscopy (EIS) measurements, the sample was immersed in SBF for 30 min to establish OCP or the steady state potential. EIS measurements were carried out from 10 mHz to 100 kHz at OCP values, using 10 mV (root-mean-square, RMS) as the amplitude of the perturbation signal. The impedance data was shown as Nyquist plots. The potentiodynamic polarization (PDP) test was carried out at a scanning rate of 1 mV/s. The PDP and EIS results were analyzed using CorrView software and ZSimpWin software, respectively. 2.4. Immersion test. An immersion test was carried out in SBF for 7 days, and each sample was immersed in SBF separately (SBF volume (mL)/sample surface area (cm2) = 50). The temperature of the solution was controlled at 37 °C with a water bath. The hydrogen volume and the pH of SBF were recorded during the immersion. The surface morphologies and compositions of the samples after the immersion test were observed with SEM, EDS and XRD.

2. MATERIALS AND METHODS 2.1. Sample Preparation. HA/Mg composites were prepared by a conventional powder metallurgy method. The powder mixture containing 10 wt % of HA powder (99.5%, average particle size of 30 μm) and the balance of pure magnesium powder (99.9%, average particle size of 100 μm) was dried in a vacuum-dry oven at 120 °C for 2 h and then mixed by ball milling with agate balls in a polyethylene can for 1 h in argon. The powder mixture was then cold pressed into a cylindrical compact with a diameter of 40 mm at 500 MPa pressure level. The compact was hot pressed at 380 °C with 150 MPa for 20 min, and then hot extruded into a rod (10 mm in diameter) employing an extrusion ratio of 16:1. The extruded rods were machined to smaller cylindrical samples of approximately 8 mm in diameter and 5 mm in thickness. The SEM image and EDS elemental maps of the composites are shown in Figure 1, revealing that the HA particles are distributed homogeneously on the magnesium matrix. The detailed conversion coating process was described in our previous work.23,24 In brief, the samples were treated in a phosphating bath containing 11 g/L Ca(NO3)2·4H2O, 1.2 g/L CaO, 8 mL/L H3PO4(85% V/V), and 0.5 g/L Na2MoO4·2H2O at 37 ± 2 °C, for different deposition times (from 5 s to 20 min) to obtain DCPD conversion coatings. The pH of the phosphating solution was adjusted to 2.8−3.0 using sodium hydroxide (NaOH). The DCPD coatings with different deposition times were then alkali post-treated in 1 mol/ L NaOH solution at 80 °C for 2 h to obtain HA coatings. 2.2. Surface Characterizations. The surface morphologies and elemental compositions of the coatings were identified by scanning electron microscopy (SEM, ZEISS EV018, Germany) and energy dispersive X-ray spectrometry (EDS, Oxford instruments X-Max, England). The phases of the samples were characterized by X-ray diffraction (XRD, Rigaku Dymax, Japan) using Cu Kα radiation (λ = 0.154178 nm) and a monochromator at 40 kV, 200 mA with a scanning rate and step of 4°/min and 0.02°, respectively. 2.3. Electrochemical Test. The open circuit potential (OCP) of the Mg composites during the conversion coating process was monitored via an Electrochemical Analyzer (Versa STAAT3, Princeton Applied Research, USA) from the beginning of immersion to 1200 s. The electrochemical corrosion behavior was performed in SBF at 37 ± 0.5 °C. The SBF was composed of 8.0 g/L NaCl, 0.4 g/L KCl, 0.14 g/ L CaCl2, 0.35 g/L NaHCO3, 1.0 g/L C6H6O6 (glucose), 0.2 g/L

3. RESULTS AND DISCUSSION 3.1. Evolution of Conversion Coating with Coating Time. During the conversion coating process, the reactions on the sample surface were considered to take place correspondingly on different local polarization sites Mg → Mg 2 + + 2e−

(1)

−

2H 2O + 2e → H 2 + 2OH Mg + 2H 2O → Mg

2+

−

(2) −

+ 2OH + H 2

(3)

The hydrogen generation reaction (eq 2) promoted the local pH value at the solution-substrate interface. CaPs can be precipitated readily in the selected coating conditions when their solubility limits were exceeded as the pH value increased. During this process, the potential of the sample in the solution would change with time as the coating gradually grew on the surface. Therefore, to monitor the deposition process of the conversion coating, we recorded the open circuit potential (OCP) of the composites in the phosphating bath as a function of immersion time, as shown in Figure 2, and the OCP curves B


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corresponding to points A−F in Figure 2. Their corresponding EDS spectra are shown in Figure 4. It has been pointed out that

Figure 2. OCP curves during the conversion coating processes on the pure Mg, AZ60 alloy, and HA/Mg composites.

for pure Mg and an AZ60 alloy are also listed for comparison. All the OCP curves followed a similar trend, i.e., the OCP decreased quickly in the first several seconds before gradually increasing and becoming stable after a longer immersion time. This trend is in agreement with those obtained in previous conversion coating studies.28−30 The initial decrease in the OCP was probably caused by the surface activation of the substrate, and the dissolution of the loose surface oxide film in the acidic phosphating solution.28 The much sharper decrease observed for the composites indicates the quicker activation and dissolution than the pure Mg and AZ60 alloy. This is possibly because the microstructure of the composites can induce the galvanic corrosion in the present phosphating solution much more easily, because of the more evident electrochemical heterogeneity between the HA phase and the Mg matrix. Figure 3 shows the SEM micrographs of the DCPD conversion coatings prepared at different coating times,

Figure 4. EDS spectra of DCPD coatings formed at different deposition times (a) 5 s, (b) 1 min, (c) 5 min, (d) 10 min, (e) 15 min, (f) 20 min.

the minimum OCP value is associated with the end of the surface activation, while the increase of the potential value is associated with the start of the coating deposition.28,29 However, in the present study some leaflike flakes formed and even accumulated into large flake clusters on the reinforcement (HA phase) after immersion in the phosphating bath for 5 s, as shown in Figure 3 a (corresponding to point A in Figure 2). The EDS spectrum shows that Ca and P signals were also detected on the coating surface (Figure 4 a). Although the earlier coating nucleation has also been observed during the coating process on the AZ60 alloy, the larger flake clusters deposited on the composites indicate that the more evident electrochemical heterogeneity turns into a priority during the conversion coating process. Therefore, as shown in Figure 2, the increase of the OCP value following its initial decrease was much more rapid and dramatic. It was found that β-TCP in the Mg composites could promote the CaPs formation on the surface,31 and implanted calcium ions in titanium could also act as a catalyst to accelerate the CaPs precipitation.32 Therefore, the HA particles in the composites could also catalytically adsorb Ca2+ and PO43− to accelerate the DCPD deposition. After 1 min of immersion in the phosphating bath (Figure 3b), most of the substrate was covered with loosely packed flakes. The EDS spectrum in Figure 4 b shows that the Ca and P signals increased greatly. The flake sizes were not quite uniform, which was mainly due to the earlier nucleation and growth of the coating on the HA particles than on the Mg matrix. When the coating process lasted for 5 min (corresponding to point C in Figure 2), a plateau was reached in the OCP curve, suggesting that the coating reactions at the solution/substrate interface reached a steady state, i.e., a dynamic balance was established in the coating deposition process. The coating flakes increased considerably in size, but the flake edges slightly dissolved into a butterfly wing-like morphology, while some newly formed flowerlike flakes that appeared on the coating surface did not show evident dissolution (Figure 3c). This indicates that the formation and

Figure 3. SEM micrographs of DCPD coatings formed at different deposition times (a) 5 s, (b) 1 min, (c) 5 min, (d) 10 min, (e) 15 min, (f) 20 min. C


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should be especially noted that the coating at 10 min has more than 12 times reduced icorr and also 25 times increased polarization resistance than the uncoated composites. However, when the coating time increased to 15 min, Ecorr decreased mildly to −1409 mV/SCE and icorr also increased slightly. At 20 min, Ecorr further decreased to −1445 mV/SCE and icorr further increased to about two times that at 10 min. This indicates that the coating at 25 min has worse electrochemical corrosion resistance than that at 10 min, so the best corrosion resistance does not correspond to the longest treatment time. According to the microstructure evolution and electrochemical corrosion behaviors with the coating time, the conversion coating after 10 min of immersion in the phosphating bath owns the uniform and dense coating structure, and thus the optimum electrochemical corrosion resistance. However, the coatings at 5 and 15 min also showed uniform coating morphologies and good electrochemical behaviors, so these three DCPD coatings after 5, 10, and 15 min of deposition were employed in the following alkali posttreatment to obtain different HA coatings, labeled as HA-5 min, HA-10 min, and HA-15 min, respectively. 3.3. Effect of Alkali Post-Treatment on DCPD Coatings with Different Deposition Times. Typical surface morphologies and their corresponding EDS results of HA coatings with different DCPD deposition times are shown in Figure 6. It can be observed that the coating flakes of all the coatings were slightly dissolved to be clustered together, and then most micropores between the coating flakes were filled. Thus, the coatings became denser to cover the substrate more completely. The EDS results show that the Ca/P ratio for all the HA coatings are around 1.3, and the Ca and P peaks were lower than those of the corresponding precursor DCPD coatings. However, it seems that the coating flakes of the HA-5 min coating are bigger in size and thus give a better coverage on the substrate (Figure 6a), possibly because the DCPD-5 min coating has more newly formed flowerlike flakes on the surface, which are not easy to dissolve during the alkali post treatment. This is also the reason why the Ca and P peaks for the HA-5 min coating were close to those of the DCPD-5 min and HA10 min coatings. To obtain the optimum HA coating, the electrochemical corrosion behavior of the three HA-coated samples was assessed by PDP test, and the resulting curves are shown in Figure 7, and the PDP curves for the uncoated and corresponding precursor DCPD-coated samples are also listed for comparison. The values of the electrochemical parameters obtained from the curves are summarized in Table 2. For all the HA coatings, in comparison with their corresponding precursor DCPD coatings, the corrosion potential shifted greatly in the positive direction, the corrosion current density decreased (9, 4, and 2, times lower for the HA-5 min, HA-10 min, and HA-15 min, respectively), and the polarization resistance increased significantly. It is also shown in Figure 7 that all the HA coatings have a similar cathodic current density and corrosion potential, but the HA-5 min coating has the lowest corrosion current density and anodic current density, and also the much greater Rp (about 4 times more than its precursor DCPD coating). This indicates that the HA-5 min coating owns the optimum electrochemical corrosion resistance, and the performance of the HA coating is critically related to the coating surface structure, but it does not depend on its precursor coating or the coating thickness.

dissolution of the coating proceed simultaneously during the conversion coating process. The coating surface after 10 min of deposition showed a flattened morphology on the substrate surface (Figure 3d), which made the coating structure denser than that at 5 min. However, the newly formed flowerlike flakes reduced gradually with the coating time and disappeared for the coating at 15 min (Figure 3e), indicating that the coating dissolution was the dominant proceeding in this period. When the coating time increased to 20 min, as shown in Figure 3f, the coating showed an irregular fragment-like morphology, composed mainly of small chippings and some newly formed flakes. As for the EDS spectra for the coating from 5 to 20 min, as shown in Figure 4c−f, the peak strengths of Ca and P signals increased gradually. Therefore, immersion in the phosphating bath for a longer time could obtain thicker coating, but it is not beneficial for the coating structure and thus potentially influences its corrosion resistance. 3.2. Effect of Coating Time on Electrochemical Corrosion Behavior of DCPD Coating. Figure 5 shows

Figure 5. PDP curves for the uncoated and DCPD-coated HA/Mg composites in SBF.

the PDP curves for the uncoated and DCPD coated composites in SBF (at 37 °C). The values of the electrochemical parameters obtained from the PDP curves are summarized in Table 1. The polarization resistance Rp is taken from the fitting result of the data in the range of −30 mV to +30 mV versus OCP using CorrView software. Table 1. Values of Electrochemical Parameters Derived from the PDP Curves (Figure 5) uncoated DCPD-1 min-coated DCPD-5 min-coated DCPD-10 min-coated DCPD-15 min-coated DCPD-20 min-coated

Ecorr (mV/SCE)

icorr (μA/cm2)

Rp (Ω cm2)

−1525 −1504 −1468 −1384 −1409 −1445

343 132 37 27 31 50

91 280 1191 2291 1769 470

It can be seen that all of the coatings have more positive corrosion potentials, lower corrosion current densities and larger polarization resistances, compared to the uncoated composites. When the coating time was equal to or less than 10 min, Ecorr showed gradual shifts in the positive direction, icorr decreased and Rp increased gradually with coating time. It D


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Figure 6. Surface morphologies of HA coatings with different DCPD deposition times: (a) 5 min (HA-5 min), (b) 10 min (HA-10 min), (c) 15 min (HA-15 min), and their corresponding EDS results.

Table 2. Values of Electrochemical Parameters Derived from the PDP Curves (Figure 7) uncoated DCPD-5 min coated DCPD-10 min coated DCPD-15 min coated HA-5 min coated HA-10 min coated HA-15 min coated

Ecorr (mV/SCE)

icorr (μA/cm2)

Rp (Ω cm2)

−1525 −1468 −1384 −1409 −1373 −1341 −1400

343 37 27 31 4 7 15

91 1191 2291 1769 4242 3126 1952

HA coatings, respectively, to study the effects of these two optimum coatings on the corrosion behavior of the composites. 3.4. Microstructure and Phase Composition Analysis. Figure 8 illustrates the XRD patterns of the uncoated, DCPD, and HA-coated samples. The diffraction pattern of the uncoated HA/Mg composites exhibited a large number of strong Mg reflections and some less detectable HA reflections. The DCPD coating consists mainly of dicalcium phosphate dihydrate (DCPD, CaHPO4·2H2O). Because of the porous

Figure 7. PDP curves for the HA coatings with different DCPD deposition times in SBF. The curves for the uncoated and corresponding precursor DCPD coated samples are listed for comparison.

Therefore, the DCPD-10 min and HA-5 min coatings were employed in the following tests and labeled as the DCPD and E


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reaction (eq 1) and the hydrogen generation reaction (eq 2) mentioned above, Ca2+ and H2PO4−, which were produced during the phosphating solution preparation, reacted to form DCPD CaO + 2H3PO4 → Ca 2 + + 2H 2PO4 − + H 2O

(4)

H 2PO4 − + OH− → HPO4 2 − + H 2O

(5)

Ca 2 + + HPO4 2 − + 2H 2O → CaHPO4 ·2H 2O

(6)

As a precursor to HA, DCPD in the conversion coating is unstable in environments with pH greater than 6−7,33 so the alkali post-treatment transformed DCPD into the HA 10CaHPO4 ·2H 2O + 12OH− → Ca10(PO4 )6 (OH)2 + 4PO4 3 − + 30H 2O

Figure 8. XRD patterns of (I) uncoated, (II) DCPD, and (III) HAcoated HA/Mg composites.

(7)

Figure 9 shows the surface and cross-sectional morphologies of DCPD and HA coatings at high magnification, and their corresponding EDS results. As shown in Figure 9a, the newly formed flowerlike DCPD coating cluster is composed of flakes diverging from the center toward the periphery, and the porous coating is approximately 5 μm thick. After the alkali posttreatment, some coating flakes broke off or were dissolved, and then most pores between flakes were filled. Thus, the coating thickness decreased slightly (∼3 μm), and became denser and better adhered onto the substrate, as shown in Figure 9b. The EDS results show that the Ca/P ratio for the coating increased from 0.66 to 1.33, which is in accordance with previous studies,34,35 and less than that of stoichiometric HA (1.67),

structure of the coating, the reflections from the substrate were readily detected, but an HA phase in the composites was not detected, and the strength of the Mg peaks obviously decreased compared to pattern I for the uncoated composites. After the alkali post-treatment, HA became the major coating phase. The further decreased Mg peaks indicate a denser coating structure compared to the DCPD coating. On the basis of the phase transitions mentioned above, the following reactions could occur during the formation of DCPD and HA coatings.23,24 During the conversion coating process in the phosphating bath, except for the Mg substrate dissolution

Figure 9. Surface and cross-sectional morphologies of (a) DCPD and (b) HA coatings on the HA/Mg composite, and their corresponding EDS results. F


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(Cf) in parallel with a film resistance (Rf), and a charge transfer resistance (Rct) in parallel with the double layer capacity (Cdl). The EIS fitting results are listed in Table 3. The value of the admittance Y01 associated with the HF loop for the uncoated composites was 28 μΩ−1 cm−2 sn, which is greater than that of around 10 μΩ−1 cm−2 sn for the Mg alloys in previous studies,37−39 suggesting that the HA/Mg composites had a more defective corrosion product film than Mg alloys. This is possibly because ceramic particles deteriorated the corrosion film consistency and thus increased the film defects and roughness. Compared to the uncoated composites, the coated samples had greater sizes of capacitive loops. It is well-known that the diameter of the semicircle is related to the corrosion resistance of the sample. The HA coated sample showed a larger semicircle than the other samples, indicating the most improved corrosion resistance of the HA coating. In addition, n can be used as a measure of the surface in homogeneity. The smaller values of n, the higher the surface roughness. Both of the coatings had lower Y01 values and increased n1 values, as seen in Table 3, suggesting that the coatings were neater with less defects than the corrosion product film on the uncoated composites. The coating resistance and charge transfer resistance of the HA coating were 1828 and 4108 Ω cm2, respectively. It is noteworthy that the total resistance (Rtot) offered by the HA coating is greater than that of the DCPD coating, and about 6 times higher than that of the uncoated composites. According to the above results, the DCPD coating obtained in this study is porous in nature, but its passivation during the corrosion process could provide effective corrosion protection for the composites substrate, while the dense HA coating can act as a barrier layer and shows the most improved corrosion resistance. 3.6. Corrosion Behavior and Coating Evolution during Immersion Test. When the samples were immersed in SBF, the hydrogen evolution volume and pH of SBF were monitored, and the results are shown in Figure 11. During the first several hours of immersion, few hydrogen bubbles appeared on the surfaces of the coated samples, while large numbers of bubbles were observed rising from the surface of the uncoated sample. It can be seen that the H2 volume was reduced for all the coated samples, especially for the HA coated sample, being around seven times less than that of the uncoated sample over the whole period of 168 h (Figure 11a). The improved coating resistance and charge transfer resistance as shown in the above EIS analysis can explain the good corrosion protection ability for the DCPD and HA coated samples. The hydrogen evolution rates of the DCPD, and HA coated specimen are 2.33 mL cm−2 day and 0.91 mL cm−2 day, respectively, which are significantly lower than the uncoated HA/Mg composites (6.25 mL/cm2/day). The much decreased hydrogen evolution of the HA-coated sample could potentially be lower than the hydrogen absorption capacity of human body.6,40,41

indicating that the HA coating is calcium deficient. Many studies have indicated that the dissolution of HA in the human body after implantation is too slow to achieve optimal results, whereas Ca-deficient HA seems to be more soluble and may induce precipitation of a new bonelike apatite after implantation.35,36 3.5. EIS Test. The PDP curves for the uncoated, DCPD, and HA coated samples have been shown (Figures 5 and 7) and discussed in the above subsections. EIS was employed to further understand the electrochemical corrosion behavior of the coatings. Figure 10 shows the Nyquist impedance plots

Figure 10. EIS patterns for the uncoated, DCPD, and HA-coated HA/ Mg composites in SBF. Insert is an equivalent circuit for the EIS fitting (the elements in the parentheses are used to fit the plot for the uncoated sample).

obtained for the composites and the two coated samples. All the resulting plots show two capacitive loops, although the presence of two capacitive loops for the composites is not easily distinguishable because of the overlapping. An equivalent circuit used to fit the experimental impedance data is given in the inset diagram of Figure 10. As for the two coated samples, the high-frequency (HF) capacitive loop is related to the outer porous layer of the coating (characterized by Rp and Qp), while the low frequency (LF) capacitive loop is related to the inner barrier layer of the coating (described by Rb and Qb). Rs is the solution resistance. Q stands for CPE (constant phase element) and is defined by two values, admittance Y0 and the power index number n, given by Y = Y0(jω)n. If n = 1, CPE is identical to a capacitor. CPE accounts for the deviation from the ideal dielectric behavior and is related to the surface inhomogeneity, and thus commonly used in place of capacitance for a better quality fit. As for the uncoated composites, during the period for establishing a stable OCP in SBF before the EIS test, a thin film based on Mg(OH)2 would form on the surface. The capacitive loops for the composites were suggested to be related to the charge transfer and the film effect.37 Therefore, the similar electrical equivalent circuit to that for the coated samples was used: a film capacity

Table 3. Fitting Results of EIS Plots for the Uncoated, DCPD, and HA-Coated HA/Mg Composites in SBF

uncoated DCPD-coated HA-coated

Y01 (μΩ−1 cm−2 sn)

n1

Rp (Rf) (Ω cm2)

Cdl (μF/cm2)

28 4.0 3.0

0.70 0.88 0.83

757.4 1037 1828

13.7

G

Y02 (μΩ−1 cm−2 sn)

n2

Rb (Rct) (Ω cm2)

Rtot (Ω cm2)

60.3 39.7

0.55 0.71

229.1 2539 4108

986.5 3576 5936


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whereas the more stable pH values in the vicinity of the implant would reduce local alkalization and stabilize pH-dependent physiological processes, and thus could benefit osteoblast cell proliferations and help the healing of the surgical region. Osseointegration of the implant with the surrounding tissues is important for successful implantation and fixation. In this respect, the apatite-forming ability of the treated surface in SBF is regarded as an indication of osseointegration capability.43 SEM images and EDS spectra of the uncoated and coated samples after 7 days of immersion in pH of SBF are shown in Figure 12. The corresponding XRD patterns are shown in Figure 13. It can be seen that the uncoated sample was severely corroded and suffered from a large number of corrosion pits (Figure 12a). A corrosion layer with deep cracks can be observed on the uncoated composites after 7 days of immersion. Judging from the high contents of O and Mg as well as the low levels of Ca and P (below 1 at %) in the EDS result, the corrosion layer was composed of remarkable corrosion products based on Mg(OH)2 and CaP precipitates with a Ca/P molar ratio of approximately 0.85. The XRD result (Figure 13) confirmed that the corrosion products were mainly Mg(OH)2 and MgHPO4·3H2O and the small amounts of CaP precipitates were Ca-deficient HA. The strength of the Mg peaks decreased greatly because of the coverage by the substantial quantities of corrosion products. However, the surface of the DCPD-coated sample after 7 days of degradation was covered by a new mineralized layer composed mainly of clusters of white spherelike precipitates with a size of 1−2 μm in uniform distribution (Figure 12b), although a few cracks appeared on the surface because of the dehydration and shrinkage of the layer when dried in the air. The EDS result reveals that the Ca/P molar ratio of this mineralized layer was around 1.2, higher than the original DCPD coating. Although the minimum Ca/P molar ratio for apatite is 1.33,44 the lower Ca/P molar ratio is mainly due to the existence of the remaining DCPD after the flakelike coating structure was dissolved in SBF, as confirmed by the XRD pattern II in Figure 13. The stronger Mg peaks and weaker Mg(OH)2 peaks than those of the uncoated sample indicate fewer corrosion products on the surface and thus higher corrosion resistance for the DCPD coated sample. For the HA coated sample, the surface was also covered by a new mineralized layer, which was composed mainly of clusters of irregular particlelike precipitates (Figure 12c). Compared to the mineralized layer on the DCPD coated sample in Figure 12b, a denser surface structure with basically no cracks was obtained and the Ca/P molar ratio increased up to 1.57, which is much closer to stoichiometric HA (1.67). It can also be observed that more obvious HA peaks and weaker Mg(OH)2 peaks emerged in the XRD pattern III in Figure 13. As for orthopedic implants, a stable and favorable bone− implant interface is significantly important during the bone remolding process. Because of the easily formed galvanic couple at the ceramic/metal interface, the HA/Mg composites in the present study degraded too fast in a simulated biological environment and the precipitated CaP minerals were broken apart by corrosion products including Mg(OH)2 and MgHPO4· 3H2O as shown in Figure 12 b and Figure 13, leaving gaps at the bone−implant interface when applied as the implant material.3,45 Therefore, it is important to have a biocompatible coating to protect the composites and make a stable and

Figure 11. (a) Volume of released H2 and (b) pH of SBF as functions of immersion time in the immersion test.

The corrosion protection of the coatings could also be evidenced by the much smaller increase in pH of SBF (Figure 11 b). The pH of the solution with the uncoated sample increased rapidly from 7.4 to 9.0 on the second day of immersion, and then quickly stabilized at approximately 9.0, possibly because the HA particles could accelerate the coverage of the corrosion products when compared to the Mg alloy in our previous study.42 In contrast, the pH of the solution with the coated samples was quite similar and increased much more slowly. Although the pH increased to around 8.1 in the first 12 h, as quick as that with the uncoated sample, it then increased slowly to 8.4 in the first day, and remained stable at around 8.5 after that for both coatings. It is noteworthy that the solution pH of the DCPD-coated sample decreased slightly in the third and fourth days of immersion, possibly because the unstable DCPD transformed to HA in SBF following the reaction (eq 7), which decreased the OH− concentration and thus decreased the solution pH slightly. This could also explain why the solution pH of the coated samples was quite similar, although the hydrogen evolution rate of the DCPD-coated sample was quicker than that of the HA-coated sample. The much lower hydrogen evolution could significantly decrease the subcutaneous gas pockets, which can potentially loosen the connection between the implant and tissues,40 H


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Figure 12. Morphologies and EDS results of (a) uncoated, (b) DCPD, and (c) HA-coated HA/Mg composites after 7 days of immersion in SBF.

increase the amounts of calcium and phosphate ions available surrounding the implant to promote increased osseointegration.52 From the morphology evolution of DCPD and HA coatings in SBF (Figure 12), it was found that flowerlike coating flakes of both coatings were prone to dissolve first, and then apatite (mainly Ca-deficient HA) nucleation was promoted on the surface. This indicates that both CaP coatings could induce the fast precipitation of osteoconductive minerals on the HA/Mg composites.

compatible bone−implant interface in the early stage of operation. Previous in vitro and in vivo studies have shown DCPD and HA to be promising biocompatible materials for clinical use.46−49 As the naturally occurring mineral form of calcium apatite, HA is a stable coating material and can provide a bioactive and also corrosion protective surface for the Mg substrate, which will help to build a stable and favorable bone− implant interface. The situation is different for DCPD coating. DCPD exhibits greater solubility than most other CaP phases,50,51 which will cause concerns regarding long-term stability of the coated implants. However, it has been found that the CaP coatings will first dissolve and then mineralize in SBF, so the greater solubility of the DCPD coating could help to

5. CONCLUSIONS In the present study, DCPD coating was prepared on homemade HA/Mg composites using a simple conversion coating method, and then transformed to HA coating via an I


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Figure 13. XRD patterns of (I) uncoated, (II) DCPD, and (III) HAcoated HA/Mg composites after 7 days of immersion in SBF.

alkali post-treatment. Because of the more evident electrochemical heterogeneity between the HA phase and the Mg matrix, the composites substrate showed quicker activation and coating nucleation than the pure Mg and AZ60 alloy during the conversion coating process. By changing the conversion coating time, the morphologies and corrosion resistances of the DCPD and HA coatings could be regulated, and the optimum corrosion resistances of these two coatings were obtained at coating time of 10 and 5 min, respectively. Electrochemical and immersion tests showed that both the optimum coatings, especially the HA coating, dramatically reduced the degradation rate of the composites in SBF. In addition, both coatings could improve the biomineralization ability of the HA/Mg composites according to the higher Ca/P ratio of the surface precipitates during immersion in SBF.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: 86-431-85095875. Fax: 86431-85095876. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Nature Science Foundation of China (Grant 31070841). REFERENCES

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