JOURNAL OF RARE EARTHS, Vol. 33, No. 1, Jan. 2015, P. 93 Improving in-vitro biocorrosion resistance of Mg-Zn-Mn-Ca alloy in Hank’s solution through addition of cerium ZHANG Fan (张 凡)1, MA Aibin (马爱斌)1, 2,*, SONG Dan (宋 丹)1, 2,*, JIANG Jinghua (江静华)1, 3, LU Fumin (卢富敏)1, ZHANG Liuyan (张留艳)1, YANG Donghui (杨东辉)1, 2, CHEN Jianqing (陈建清)1 (1. College of Mechanics and Materials, Hohai University, Nanjing 210098, China; 2. Changzhou Research Institute of Hohai Technology Co., Ltd., Changzhou 213164, China; 3. Marine and Offshore Engineering Research Institute, Hohai University, Nantong 226000, China) Received 28 April 2014; revised 27 August 2014 Abstract: Two kinds of Mg-Zn-Mn-Ca alloys with and without cerium were designed and fabricated. In-vitro degradation tests and electrochemical evaluations were carried out to compare their biocorrosion behavior in Hank’s solution at 37 ºC. After adding cerium, the continuous network distributed Ca2Mg6Zn3 phases in Mg-2Zn-0.5Mn-1Ca alloy (Alloy I) were separated due to the emerging non-continuously distributed Mg2Ca phase and Mg12CeZn phase. This change led to corrosion acceleration of Mg matrix at the initial stage but also sped up the formation of compact corrosion products for Mg-2Zn-0.5Mn-1Ca-1.5Ce alloy (Alloy II), and therefore enhanced its biocorrosion resistance. Cerium containing Alloy II has the potential to be used as future biomaterials. Keywords: biomedical Mg alloy; rare earth alloying; in-vitro degradation; corrosion resistance; rare earths Due to its low density and high specific strength, magnesium and its alloys are widely used in many industrial sections including aerospace component, computer parts, mobile phones, etc. Compared to other commonly used biomaterials, magnesium and its alloys have excellent biocompatibility and biodegradability as well as an elastic modulus which is closer to that of human bones and also a possibility of being gradually dissolved and absorbed after implanting in a physiological environment. As a result, more and more attention has been paid to their potential applications as biomedical materials. Recent applications include vascular stents and bone implants[1–3]. However, the Mg alloys commonly used in recent researches were originally developed for structural materials and therefore did not consider its bio-safety necessary for use as biomaterial. Moreover, the degradation rate of these alloys in corrosive medium, especially in Cl–-containing physiological environment, is so fast that the implants can not hold their mechanical integrity before restoration of the human tissue[4–6]. Furthermore, the volume of hydrogen generated during the corrosion process of the common magnesium alloys generally exceeds human tolerance. These disadvantages limit the further applications of commonly used Mg alloys as biomaterials. Research has shown that the alloying elements can be an important factor to influence the corrosion resistance of magnesium alloys. The common alloying elements include Al, Zn, Mn, Ca and rare-earth (RE) elements. Among these alloying elements, Al has been regarded as a neurotoxin which can induce dementia. While toxicity researches about Zn, Mn, Ca have shown no harm to the human health[8–10]. Although the toxicity of RE on the human body is not entirely clear, one thing is reasonable that toxic effects could be avoided by controlling the released concentration of the RE ions below the critical toxic concentration. Mn is mainly used to enhance ductility. More important is the formation of intermetallic phases which can pick up iron (Fe) and can therefore be used to control the corrosion of magnesium alloys due to the detrimental effect of Fe on the corrosion behavior. In smaller amounts, Zn contributes to strength due to solid solution strengthening. It can also improve the castability. But larger amounts (>2 wt.%) of Zn leads to a decrease of the corrosion resistance. Ca contributes to solid solution strengthening and precipitation strengthening. It also acts, to some extent, as a grain refining agent and additionally contributes to grain boundary strengthening. In binary Mg-Ca alloys, the Mg2Ca phase is formed which can improve creep resistance due to solid solution strengthening, precipitation strengthening and grain boundary pinning. Larger amounts of Ca (>1 wt.%) can lead to problems like hot tearing or sticking during Foundation item: Project supported by National Natural Science Foundation of China (51141002), Natural Science Foundation of Jiangsu Province (BK2011249) and the Fundamental Research Funds for the Central Universities of China (2011B08214) * Corresponding authors: MA Aibin, SONG Dan (E-mail: [email protected], [email protected]; Tel.: +86-25-83787239) DOI: 10.1016/S1002-0721(14)60388-4 94 JOURNAL OF RARE EARTHS, Vol. 33, No. 1, Jan. 2015 casting. RE elements are introduced into magnesium alloys normally by master alloys. In general, the RE can be divided into two groups: elements with large solid solubilities in Mg (such as Y, Gd, Tb, Dy, Ho, Er, Tm, Yb, and Lu) and elements with limited solubility in Mg (Nd, La, Ce, Pr, Sm, Eu). All RE elements can form complex intermetallic phases with Mg, which can act as obstacles for impeding dislocation movement at elevated temperatures and cause precipitation strengthening. The RE elements with limited solubility form intermetallic phases early during solidification. Thus, RE elements can arrest grain boundaries at elevated temperatures and contribute to higher strength mainly by precipitation strengthening. This mechanism increases the service temperature of Mg alloys in transportation industry and improves creep resistance as well as corrosion resistance. In this work, one novel kind of Mg-Zn-Mn-Ca alloy (nominal composition: Mg-2Zn-0.5Mn-1Ca, designated as Alloy I) and another RE-containing Mg alloy (nominal composition: Mg-2Zn-0.5Mn-1Ca-1.5RE, designated as Alloy II) were designed and fabricated after considering the influence of alloying elements on both human health and corrosion properties of magnesium alloys. Here, Ce was chosen as an additional RE element by considering its availability and its reported benefits for health[14–16]. In-vitro degradation tests and electrochemical evaluation (including open circuit potential, potentiodynamic polarization curves and electrochemical impedance spectra) together with corrosion morphologies observation were carried out to investigate the biocorrosion behaviour of two new alloys in Hank’s solution at 37 °C. The influence of the Ce alloying on magnesium alloy was researched by analysis of the microstructure change and degradation products of the novel alloy. Table 1 Chemical composition of experimental alloys (wt.%) 1 Experimental 1.1 Materials Two kinds of alloys were prepared from high purity Mg (99.99%), pure Zn (99.9%), Mg-5 wt.%Mn master alloy, Mg-30 wt.%Ca master alloy and Mg-25 wt.%Ce master alloy. Melting and alloying operations were carried out in a stainless steel crucible, under the protection of a mixed gas atmosphere of SF6 (0.3 vol.%) and CO2 (Bal.). Pure Zn and master alloys were added to the pure Mg melts at 730 ºC. The melts were then kept for 30 min at 730 ºC to ensure that all the required alloying elements had dissolved in the molten alloy. It was then cooled to 700 ºC and poured into a steel mold, which had been pre-heated at 200 ºC, to form an ingot. The chemical compositions of the alloys were analyzed by inductively coupled plasma optical emission spectrometry (ICP-OES) (Iris Advantage 1000, USA), as listed in Table 1. Alloys Zn Mn Ca Ce Mg Alloy I 2.13 0.43 Alloy II 2.00 0.50 0.90 – Bal. 1.02 1.35 Bal. 1.2 Microstructure characterization The microstructures of the alloys were examined by BX51M (Olympus, Japan) optical microscopy (OM) and S340-N (Hitachi, Japan) scanning electron microscopy (SEM) with energy dispersive spectrum (EDS). The preparation of OM/SEM samples consisted of grinding on SiC paper up to 2000 grit and subsequent mechanical polishing. Before observation, the samples were etched using 4 vol.% nitric acid in ethanol for 5–8 s, thoroughly washed with water and alcohol, and then dried by hot air. For the phase analysis, X-ray diffraction (XRD) measurements were performed using a Bruker D8 Advance (Bruker AXS, Germany). The samples were investigated in parallel beam geometry, using Cu Kα1 radiation. The diffraction patterns were measured from 2θ (10º–90º) for each sample with the scanning rate of 2 (º)/min. 1.3 Immersion test Table 2 presents the chemical composition of the Hank’s solution used in this work. Hank’s solution was made using BS004 Hank’s Balance Salts (without sodium bicarbonate, Merck Millipore Beijing Skywing), sodium bicarbonate (reagent grade) and distilled water. The initial pH value of Hank’s solution was 7.4, which is similar to that of human’s blood plasma. The immersion tests were conducted for 10 d at 37±0.5 ºC using water bath to evaluate the rate of in-vitro degradation, under the condition of refreshing the Hank’s solution each day. Samples for immersion testing with a dimension of 10 mm×10 mm×5 mm were cut from the cast ingot. Pure magnesium sample was used here as contrast. The specimens were polished with emery papers up to 2000 grit, cleaned ultrasonically in alcohol, and then dried in open air. All specimens were immersed in 500 mL beaker containing 300 mL Hank’s solution. Hydrogen bubbles from each specimen were collected into a burette. 1.4 Corrosion morphology observation After the immersion test, the specimens were cleaned with chromate acid (200 g/L CrO3+10 g/L AgNO3) for 5 min to remove the corrosion products, and then cleaned with distilled water and dried in air. The surface morTable 2 Chemical composition of Hank’s solution used in this work (mmol/L) Solution NaCl CaCl2 MgSO4 KCl KH2PO4 Na2HPO4 D-Glucose NaHCO3 Hank’s 137 1.261 0.814 5.33 0.44 0.338 5.56 4.17 ZHANG Fan et al., Improving in-vitro biocorrosion resistance of Mg-Zn-Mn-Ca alloy in Hank’s solution through … 95 phologies before and after removal of corrosion products were observed by a HIROX KH-7700 digital microscope regularly. The chemical compositions of the corrosion products were examined using a PHI 5000 VersaProbe (UlVAC-PHI, Japan) X-ray photoelectron spectroscope (XPS) and EDS. For the XPS analysis, the specimen was removed from the solution, washed with distilled water, dried and then some corrosion products were collected for XPS analysis by scraping the surface. tential through a frequency domain from 100 kHz down to 10 mHz. The EIS was recorded after achieving a steady state OCP. The potentiodynamic polarization tests were carried out with a scanning rate of 1 mV/s over the potential range from –1.8 V vs. SCE to different anodic limits. 1.5 Electrochemical measurement Fig. 1 shows the microstructures of Alloy I and Alloy II. The grain size of Alloy I is about 125 μm, but that of, the Ce-containing, Alloy II was reduced to 65 μm. For both alloys, second phases, mainly strip-like, distributed along the grain boundaries while some granular phases were also found within the inner grains. It is obvious that the non-continuously distributed part of second phases increased after rare earth alloying. Fig. 2 shows the chemical composition of the second phases identified by EDS. For Alloy I, the second phases at the grain boundary and inner grain were composed of Mg, Zn and Ca elements as shown in Fig. 2(a), and the atomic ratio of Zn to Ca was about 1.25. According to the ternary phase diagram of Mg-Ca-Zn and the XRD patterns shown in Fig. 3, the second phase was identified as ternary Ca2Mg6Zn3 phase. However, after the addition of Ce, additional Mg2Ca and Mg12CeZn phase appeared in Alloy II. The addition of Ce changed the composition and the distribution of the second phase, thus further affecting the corrosion properties of the alloy. Rectangular specimens were molded into epoxy resin with only one side, with area of 1 cm2, exposed for the test. Pure magnesium samples were used here as contrast. The working surface was ground with SiC emery papers up to 2000 grit. Before the electrochemical measurement, all the specimens were immersed in Hank’s solution from 1 h up to 10 d. Electrochemical tests were carried out at 37 °C in a beaker containing 250 mL Hank’s solution on a PARSTAT 2273 advanced electrochemical system using a standard three-electrode configuration. The configuration was made up of a saturated calomel electrode (SCE) as a reference, a platinum electrode as a counter and the sample as a working electrode. All the samples were first used to measure electrochemical impedance spectroscopy (EIS), and then their potentiodynamic polarization curves were obtained. Before the EIS step, the open circuit potential (OCP) curve of each sample was recorded up to 1800 s. The impedance diagrams were recorded at the OCP by applying a 5 mV sinusoidal po- 2 Results 2.1 Microstructure characterization Fig. 1 Microstructures of Alloy I and Alloy II (a) OM image of Alloy I; (b) OM image of Alloy II; (c) SEM image of Alloy I; (d) SEM image of Alloy II with the arrows indicating the non-continuous second phase 96 JOURNAL OF RARE EARTHS, Vol. 33, No. 1, Jan. 2015 Fig. 2 EDS analysis of the secondary phases in the cast samples (a) Alloy I; (b) Alloy II Fig. 3 XRD patterns of Alloy I and Alloy II 2.2 Immersion test Fig. 4 presents the hydrogen evolution data of the samples immersed in Hank’s solution. The curves can be summed up into three stages. In the initial stage (less than 50 h), the corrosion products formed gradually on the surface and all the samples showed low degradation rate. During the immersion period from 50 to 150 h, the degradation rate of all the samples increased. This may be related to the local pitting corrosion which occurred on the surface of corrosion products. After 150 h of immersion, the degradation rate of HP Mg sample increased further, whereas the other two alloys showed almost constant degradation rates. This phenomenon is probably due to the corrosion barrier effect caused by the second phases in the alloys. Table 3 lists the average degradation rate of each sample. The mean degradation rate of HP Mg is the highest, while the Ce-containing Alloy II Fig. 4 Hydrogen evolution of as cast samples immersed in Hank’s solution Table 3 Degradation rate (mL/cm2/d) of the samples Samples Steady period Average Alloy I 3.25 2.79 Alloy II 2.36 2.02 HP Mg 5.09 3.76 shows the best corrosion resistance. Fig. 5 shows the surface morphologies of each sample immersed in Hank’s solution. It is obvious that there are white corrosion products heterogeneously distributed on each sample surface. Fig. 6 shows the surface morphologies of each immersed sample after the removal of corrosion products. The HP Mg sample showed an almost entirely corroded morphology with several obvious corrosion pits (dark pits in Fig. 6(c)). Alloy II maintained the most “no-corroded” area (bright area in Fig. 6(b)), thus showed the best corrosion resistance after a long period ZHANG Fan et al., Improving in-vitro biocorrosion resistance of Mg-Zn-Mn-Ca alloy in Hank’s solution through … 97 Fig. 5 Macro morphologies of the samples after immersion in Hank’s solution for 10 d (a) Alloy I; (b) Alloy II; (c) HP Mg Fig. 6 Macro morphologies of the samples after removal of corrosion products (a) Alloy I; (b) Alloy II; (c) HP Mg of immersion which is in accordance with the hydrogen evolution data. 2.3 Electrochemical measurements Fig. 7 illustrates the OCP curves of the alloy samples compared with HP Mg. All the OCP curves of the samples experienced two stages, which are, the rapid growth at initial stage and the subsequent stage of slow decline until balance. Both alloy specimen possessed higher initial OCP values than that of HP Mg, and the initial OCP value of Alloy II sample was the highest due to the addition of Ce. Furthermore, the required time for the cast alloy samples to arrive at their peak OCP value are less than that of HP Mg, and this time tends to decrease with Fig. 7 OCP curves of the samples immersed in Hank’s solution for the initial 1800 s the addition of Ce. This initial potential increase to a maximum is consistent with an incubation period for the initiation of corrosion. Considering that the steady OCP value represents the corrosive tendency of a sample at the initial stage, the lowest steady OCP value for Alloy II sample indicates that it will easily undergo corrosion during the first hours of immersion. This phenomenon is consistent with the hydrogen evolution data. Fig. 8 presents the potentiodynamic polarization curves of the samples. Both polarization curves of the two alloys are similar in shape, but are different from that of pure magnesium. The corrosion potential (Ecorr) and corrosion current density (icorr) were calculated from the intersection of the anodic and cathodic Tafel line extrapolations as listed in Table 4. Although the Ecorr of the two alloys Fig. 8 Potentiodynamic polarization curves for the samples after immersion in Hank’s solution for 1 h 98 JOURNAL OF RARE EARTHS, Vol. 33, No. 1, Jan. 2015 Table 4 Ecorr and icorr of as-cast samples and HP Mg related with the pitting corrosion of magnesium alloy[20–22]. It is obvious that the capacitive arc diameters of the as-cast alloy samples are much larger than that of HP Mg not only at the early stage but also after a long period of immersion. Moreover, the Alloy II sample always had a maximum diameter of the capacitive arc. In order to provide better explanation of the corrosion behavior of each sample in Hank's solution, Randles equivalent circuit was used to fit the electrochemical impedance spectroscopy data as shown in Fig. 9(d). Wherein Rs, Rct and C represent the solution resistance, the charge transfer resistance and the double layer capacitor or an oxide film capacitance, respectively. Taking into account the inductive reactance behavior exhibited by the specimen in the low frequency area, the inductive components L was added in the Randles equivalent circuit. Zsimpwin impedance fitting software was used to obtain the parameters of each component in equivalent circuit. The charge transfer resistance is the most important parameter, which can be used to characterize the corrosion resistance of each sample. Table 5 is a list of the Rct values for each sample, Alloy II sample had the highest Rct value which indicates the best corrosion resistance in Hank’s solution. Samples Ecorr/V icorr/(A/cm2) HP Mg –1.481 5.74×10–5 Alloy I –1.465 2.65×10–5 Alloy II –1.458 1.14×10–5 show only a slight difference (–1.458 V for Alloy II and –1.465 V for Alloy I), the icorr of Alloy II (1.15×10–5 A/cm2) is much lower than that of Alloy I (2.65×10–5 A/cm2). In the Tafel extrapolation method for measuring the corrosion rate, the corrosion current density, icorr (mA/cm2) is estimated by Tafel extrapolation of the cathodic branch of the polarization curve, and icorr is related to the average corrosion rate using Pi=22.85icorr. Therefore, the icorr of the alloys can characterize the corrosion resistance and the relative relationship of icorr is in accordance with the result of immersion test. Fig. 9 represents the Nyquist plots of as cast samples and HP Mg after immersion in Hank’s solution. All the samples have similar shape of the Nyquist spectra, which consist of a capacitive arc in the high frequency stage and an inductive loop in the low frequency stage. Nadine et al. pointed out that the capacitive arc was caused by dissolution of metal in corrosive media and the diameter of the capacitive arc was associated with charge transfer resistance which could indirectly characterize the corrosion resistance. The larger the capacitive arc diameter, the better the corrosion resistance of the material. The inductive loop is more complex and it is proposed to be Table 5 Fitting values of Rct (Ω·cm2) for the samples Time HP Mg Alloy I Alloy II 1h 804 3336 3761 2d 720 1092 1728 10 d 194 302 411 Fig. 9 Nyquist plots of the samples after immersion in Hank’s solution for 1 h (a), 2 d (b), 10 d (c) and equivalent circuit of the samples (d) ZHANG Fan et al., Improving in-vitro biocorrosion resistance of Mg-Zn-Mn-Ca alloy in Hank’s solution through … 99 Fig. 10 EDS analysis of corrosion products of the samples immersed in Hank’s solution for 1 d (a) Alloy I; (b) Alloy II 2.4 Corrode surface characterization Fig. 10 illustrates the EDS analysis of corrosion products of the as-cast alloy samples immersed in Hank’s solution for 1 d. For the Alloy I, the corrosion products showed a morphology with coarse grains and some cracks on the surface. The EDS spectra indicates that in the rough area of corrosion morphology, C, O, Mg were the main elements constituting the corrosion products, in addition to a small amount of Cl and Ca. As for the Alloy II, the product had formed relatively dense agglomerates as illustrated in Fig. 10(b). From the result of the spectra, its corrosion product mainly contains C, O, P, Mg, and Ca. The appearance of P in Alloy II indicates that there are some phosphates in the corrosion products, which made them more dense. Fig. 11 shows the XPS spectra for corrosion products of AlloyⅠafter immersion in Hank’s solution for 10 d. The main elements of the corrosion products in this stage for Alloy I samples are O, Mg, Ca, and C. Fig. 12 represents the XPS peak fittings for Ca 2p, C 1s, O 1s and P 2p orbital. Wherein, Ca element is mainly present in the form of a divalent corrosion product, C element is in the form of carbonate ion, whereas Mg mainly includes two forms, as carbonate and magnesium hydroxide. P element is mainly present as phosphate and the content is relatively small. The analysis of Alloy II sample showed a similar result. It can therefore be deduced that the corrosion products of the two alloys had no obvious difference after long period of immersion. Fig. 11 XPS survey spectra for corrosion products of Alloy I after immersion in Hank’s solution for 10 d 3 Discussion As mentioned above, the addition of Ce resulted in the change of the composition and distribution of the second phase. After adding cerium, the second phase in the as-cast alloy changed from a single Ca2Mg6Zn3 distributed along the grain boundary to the discontinuous intergranular mixture of a new RE-containing phase and Ca2Mg6Zn3. The Ce alloying broke the distribution form of the second phase, thus formed a new intermittent-like distributed Ce-containing phase. These two factors resulted in the different performance in subsequent immersion tests and electrochemical tests for the as-cast alloys. Since the electrode potential of magnesium is very negative, there usually exists a potential gap between 100 JOURNAL OF RARE EARTHS, Vol. 33, No. 1, Jan. 2015 Fig. 12 XPS peak fittings of Alloy I (a) Ca 2p; (b) C 1s; (c) O 1s; (d) P 2p magnesium matrix and the second phase, which constitutes a micro electrical coupling and results in a preferential corrosion of magnesium matrix around the second phase when the magnesium alloy was immersed in corrosive medium. For the above reasons, the volume fraction and distribution of second phases in magnesium alloys had a significant impact on its corrosion resistance. When the second phase in the alloy exists in the form of isolated coarse particles, it plays the role of accelerating the corrosion of substrate; when the second phase is distributed along the grain boundary as a continuous network structure, it acts as a barrier to prevent the further development of corrosion[23–25]. In Alloy I, cathodic Ca2Mg6Zn3 phase was a continuous network distribution around the grain boundary. When Ce was added to the alloy, as illustrated by the XRD results, the Mg2Ca and Mg12CeZn appeared in the alloy wherein Mg12CeZn was a relatively cathodic phase and the Mg2Ca an anodic phase. On the one hand, these intermittent-like distributed phases broke the netlike structure and weakened the corrosion barrier. Furthermore, the Mg12CeZn phase accelerated the corrosion of the magnesium matrix as a cathode while the Mg2Ca phase corroded preferentially as anode. The role of the new phases produced two consequences: Firstly, an increase of hydroxyl ions in the solution which promoted the generation of phosphate ions; secondly, an increase in calcium and magnesium ions in the solution. These two consequences together lead to the generation of a more compact calcium and magnesium phosphate. The protective effects of phosphate together with the accelerated corrosive effect of the cathodic phase formed a competitive relationship. In the case of Alloy II, the phosphate formed was compact and the protective effect dominated, thus resulted in its better corrosion resistance. EDS analysis of the corrosion products formed at the initial stage of immersion in Fig.10 showed that the addition of Ce had a certain impact on the formation of corrosion product in the early stage. Due to the corrosion of the magnesium matrix, the first reaction which occurred in solution was the generation of magnesium hydroxide and hydrogen. At this stage, the number of the hydroxide ions in the solution determined the time when the phosphate ions appear. Due to the role of the discrete second phases, Alloy II generated, in the initial stage, more hydroxide ions with calcium and magnesium ions, thereby promoting the formation of phosphate. Thus, the addition of Ce accelerated the formation of phosphate, thereby accelerating the formation of a dense film of corrosion product. Due to the protective effect of phosphate, Cecontaining Alloy II exhibited a better corrosion resistance. 4 Conclusions (1) The addition of cerium in Mg-2Zn-0.5Mn-1Ca al- ZHANG Fan et al., Improving in-vitro biocorrosion resistance of Mg-Zn-Mn-Ca alloy in Hank’s solution through … loy (Alloy I) enhanced its in-vitro biocorrosion resistance and reduced its degradation rate. The alloy containing 1.35 wt.% Ce (Alloy II) had the potential to be used as future biomaterials. (2) Adding cerium in Mg-Zn-Mn-Ca alloy resulted in a change of the composition and distribution of the second phases. The second phase in Alloy I was mainly Ca2Mg6Zn3 phase with a continuous network distribution, while a discrete distribution of Mg2Ca phase and Mg12CeZn phase appeared in the Ce-containing Alloy II. 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