Applied Surface Science 257 (2011) 4464–4467
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Corrosion behavior of TiO2 ﬁlms on Mg–Zn alloy in simulated body ﬂuid Shuai Chen a , Shaokang Guan a,∗ , Bin Chen a , Wen Li b , Jun Wang c , Liguo Wang a , Shijie Zhu a , Junhua Hu a a
Materials Research Centre, School of Materials Science and Engineering, Zhengzhou University, No. 97, Wenhua Road, Zhengzhou 450002, PR China Graduate School of Science and Technology, Shizuoka University, Hamamatsu 432-8011, Japan c Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan b
a r t i c l e
i n f o
Article history: Received 27 November 2010 Received in revised form 16 December 2010 Accepted 16 December 2010 Available online 23 December 2010 Keywords: Magnesium TiO2 ﬁlm Corrosion behavior Simulated body ﬂuid HAp
a b s t r a c t Magnesium alloys have been widely investigated in the ﬁeld of biomaterials due to their moderate mechanical properties close to human bone and gradual degradation in human physiological environment without second surgeries. But results from clinical studies showed that magnesium implants suffered from too rapid degradation in human physiological environment. To reduce the degradation rate of magnesium alloys, surface modiﬁcation is essential and effective besides element alloying. In this study, TiO2 ﬁlms were deposited on Mg–Zn alloy by direct current reactive magnetron sputtering. The morphology and structure of the ﬁlms were characterized by atomic force microscopy (AFM), scanning electron microscope (SEM) and X-ray diffraction (XRD). The corrosion resistance in simulated body ﬂuid (SBF) at 37 ◦ C was evaluated by potentiodynamic polarization and hydrogen evolution tests. The corrosion behavior of the samples was investigated by SEM with energy dispersive spectroscopy (EDS) after immersion for different periods. The results showed that the compact ﬁlms were composed of particles with the size of about 100 nm and could effectively improve the corrosion resistance in SBF. After immersion for 10 days, the corrosion rates of the substrates and samples with TiO2 ﬁlms were 4.13 mm/y and 1.95 mm/y, respectively. During the immersion, the TiO2 ﬁlms could induce the growth of hydroxyapatite (HAp) to improve the bioactivity of the samples. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved.
1. Introduction Magnesium alloys have been widely investigated as the promising candidates for bone implant application in the ﬁeld of biomaterials, because they possess good biodegradability in human physiological environment and excellent mechanical properties such as moderate elastic modulus which is close to that of natural bone [1–3]. But magnesium alloys are prone to rapid corrosion, especially in the environment rich in chloride ions including human body ﬂuid and blood plasma. This drawback has restricted their widespread application [4,5]. Therefore, it is necessary to reduce the corrosion rate of the magnesium alloys for the sake of wide clinical application. Methods to improve the corrosion resistance of magnesium alloys mainly involve element alloying and surface modiﬁcation [5–8]. Currently, most of the researches focused on surface modiﬁcation with biocompatible coatings [9–11]. At the same time, TiO2 ﬁlms prepared by physical methods have been also investigated as biocompatible coatings on permanent stents [12–14], and considered to be the most promising inorganic coatings due to their superior bioactivity performance and high adhesion strength with
∗ Corresponding author. Tel.: +86 0371 6388 7508; fax: +86 0371 6388 7508. E-mail address: [email protected]
the substrates . In this paper, TiO2 ﬁlms were deposited on a Mg–Zn alloy by direct current reactive magnetron sputtering. The ﬁlms were characterized by AFM, SEM, EDS and XRD. The corrosion properties of the substrates and samples with TiO2 ﬁlms were evaluated by potentiodynamic polarization and hydrogen evolution tests in SBF at 37 ◦ C. The morphology of the ﬁlms after immersion for different periods was also observed. 2. Experimental A Mg–2 wt.% Zn alloy was used as the substrate material in this study. The samples with dimensions of 4 mm × 10 mm × 25 mm were grounded by SiC emery papers to 5000 grit, and subsequently polished by diamond abrasive of W1.5 to obtain homogeneous roughness. Ultrasonic bath in acetone and ethanol was performed to degrease. TiO2 ﬁlms were deposited on the magnesium substrates (4 mm × 10 mm × 25 mm) with a direct current reactive magnetron sputtering system (JGP-450) by sputtering a high-purity titanium target (99.99%), and sustaining a constant ﬂow of argon (24 sccm, 99.99%) as sputtering gas and oxygen (6 sccm, 99.99%) as reactive gas. The working pressure was 0.75 Pa. The distance between the substrates and the target was 100 mm. The sputtering system was operated with a constant power of 325 W (325 V × 1 A). After deposition for 2 h, the thickness of the ﬁlms was about
0169-4332/$ – see front matter. Crown Copyright © 2011 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2010.12.093
S. Chen et al. / Applied Surface Science 257 (2011) 4464–4467
Table 1 Ions concentration in SBF. Ion
200 nm which was previously measured by ␣-step measurement system. The crystal structure of the ﬁlms was characterized by X-ray diffraction (XRD, X’pert PANalytical), the data was collected from 20◦ to 60◦ (2) with Cu K␣ at a step of 0.01◦ . The morphology and composition of the samples before and after immersion were identiﬁed by atomic force microscopy (AFM, XE-200) and scanning electron microscopy (SEM, JEOL-6700F) with an energy dispersive spectroscopy (EDS, Cambridge) system. The potentiodynamic polarization and hydrogen evolution tests were conducted in SBF at 37 ◦ C to evaluate the corrosion properties [1,6]. SBF was an aqueous solution with the ion concentrations are listed in Table 1. In order to prepare 1000 ml of SBF, reagents NaCl (8.035 g), NaHCO3 (0.355 g), KCl (0.225 g), K2 HPO4 ·3H2 O (0.231 g), MgCl2 ·6H2 O (0.311 g), 1.0 M-HCl (39 ml), CaCl2 (0.292 g) and Na2 SO4 (0.072 g) and Tris (6.118 g) were dissolved into the solution in the given order. The pH value of the solution was adjusted to 4.0 with 1.0 M-HCl (0–5 ml) at 37 ◦ C . The potentiodynamic polarization tests were carried out on an EG&G model potentiostat/galvanostat controlled by a computer with M352 corrosion analysis software. A three electrode cell was used in these tests with a sample as the working electrode, platinum electrode
as the counter electrode and saturated calomel electrode (SCE) as the reference electrode. After stabilization for 120 s, the potentiodynamic polarization tests were conducted with a constant voltage scanning rate of 5 mV/s. In the hydrogen evolution, each sample was immersed in a test tube (200 ml) containing SBF in a water bath at 37 ◦ C, the evolved gas bubbles were collected in a gas burette to monitor the volume. All two kinds of samples for the tests were embedded in epoxy resin with an area of 1.0 cm2 (1.0 cm × 1.0 cm) exposed to 40 ml ﬂuid according to ASTM G46 and G59. In assumption that the H2 volumes are measured under standard condition (25 ◦ C, 1 atm), the corrosion rate can be calculated according to the formula as follows:
where CR is the corrosion rate; t is the exposure time; M is the mole mass of the substrate; is the density of the substrate; n is the mole volume; V is the hydrogen volume per area that is equal to the hydrogen evolution substituted by the area exposed to the ﬂuid .
Fig. 1. Characterization of the TiO2 ﬁlms on the samples: (a) the AFM image; (b and c) the SEM images in low and high magniﬁcation; (d) the XRD pattern.
S. Chen et al. / Applied Surface Science 257 (2011) 4464–4467
Fig. 2. Immersion results of the substrates and samples with TiO2 ﬁlms in SBF at 37 ◦ C: (a) potentiodynamic polarization tests; (b) hydrogen evolution tests.
Fig. 3. (a–c) The morphology of the TiO2 ﬁlms after immersion for: 1, 3 and 5 days; (d) the EDS result of the particles on the ﬁlm.
3. Results and discussion Fig. 1(a) shows the two-dimensional AFM image of the sample with TiO2 ﬁlm. The studied area was 5 m × 5 m. It can be
seen that the TiO2 ﬁlm has a compact structure with several acute peaks which are packed densely on the surface. The AFM result also shows that the height difference in the area is within the size of 200 nm. Fig. 1(b) is the low-magniﬁcation SEM image, the TiO2
S. Chen et al. / Applied Surface Science 257 (2011) 4464–4467
ﬁlm are homogeneous without cracks. While, Fig. 1(c) is the highmagniﬁcation SEM image, it can be seen that the particles in the ﬁlm are within the size of about 100 nm. Fig. 1(d) shows the XRD pattern of the sample with TiO2 ﬁlm. Besides the typical peaks corresponding to magnesium, only diffraction peaks of TiO2 are observed in the pattern. The peaks at 25.3◦ , 48.0◦ , 38.0◦ can be perfectly in agreement with to the anatase (1 0 1), (2 0 0), (2 1 1) of TiO2 (JCPDS 84-1286). Previous studies showed that anatase TiO2 was more favorable to improve the bioactivity of the samples than the amorphous titanium oxides . Fig. 2(a) shows the polarization curves of the substrate and the sample with TiO2 ﬁlm. Compared to the substrate, the corrosion potential of the sample with TiO2 ﬁlm shifted positively from −1.8 V to −1.5 V, and the current density decreased by about two orders of magnitudes. It indicates that the TiO2 ﬁlm is more cathodic than the substrate and could improve the corrosion resistance in SBF. Fig. 2(b) shows the result of hydrogen evolution tests. The H2 volumes of the substrates exhibit approximately linear relationship with the immersion time. For the samples with TiO2 ﬁlms, the H2 volumes rise rapidly during the ﬁrst three days and seem to almost constant during the left time. Overall, the H2 volumes of the samples with TiO2 ﬁlms are much lower that of the substrate and tolerant to human body . It illustrated that the TiO2 ﬁlms prevented the substrates from rapid corrosion. By calculation, the average corrosion rates of the substrates and the samples with TiO2 ﬁlms in the ten days were about 4.13 mm/y and 1.95 mm/y, respectively. Fig. 3(a), (b), and (c) shows the morphology of the samples with TiO2 ﬁlms after immersion in SBF at 37 ◦ C for 1, 3 and 5 days, respectively. In Fig. 3(a), the ﬁlm is intact and smooth, which indicates that the ﬁlm has a good protective effect in the ﬂuid. In Fig. 3(b), the substrate in the area is still covered by the ﬁlm, but the ﬁlm ruptures and pitting corrosion in the substrate can be observed evidently. From the SEM images, the failure mechanism can be concluded as follows: ﬁrstly, the solution especially the chloride ions moved through the micro pores in the ﬁlms. Then, the solution corroded the substrates with the generation of H2 gas and corrosion products. With the increase in immersion time, the evolved H2 gas and expansion of corrosion products over the previous substrates resulted in the fracture of the ﬁlms under pressure. Additionally, galvanic cells might set up among the second phases, the substrates and the ﬁlms during the corrosion [6,18]. Consequently, Pitting corrosion contributed to an obvious increase in hydrogen evolution in Fig. 2(b). However, corrosion products were deposited on the corrosion pits and delayed the corrosion rate in some degree, which was exhibited after immersion for 3 days in Fig. 2(b). Fig. 3(c) shows some particles on the TiO2 ﬁlm, and the corresponding EDS result in Fig. 3(d) indicates the presence of carbon, sodium, magnesium, titanium, calcium, phosphorus and oxygen. Little dose of carbon and sodium came from the aqueous solution. While, magnesium and titanium were detected from the substrate and the ﬁlm, respectively. The calcium, phosphorus and oxygen with relatively higher contents were essential composition of calcium phosphate-based minerals. In this study, it was considered to be HAp, a main ingredient in natural bone and widely used in bone and dental implants. When exposed to the solution, TiO2 ﬁlms on titanium alloy induced the formation of large apatite particles by
the ion exchange between the ﬁlm and Na+ , Ca2+ , etc. in SBF in former studies [16,19]. Additionally, the ﬁne structure of the TiO2 ﬁlms was also favorable to the growth of HAp during the immersion in SBF [20,21]. Although the exact mechanism about the formation of HAp in detail is still needed to be further investigated, good bioactivity of the TiO2 ﬁlms has been demonstrated in the immersion tests. The results indicate that the ﬁlms can enhance the corrosion resistance, induce the formation of HAp in SBF and expected to be favorable to osteoblast adhesion in future application. 4. Conclusion TiO2 ﬁlms were deposited on magnesium substrates by direct current reactive magnetron sputtering. The ﬁlms were anatase and composed of particles within the size of about 100 nm. The results from potentiodynamic polarization and hydrogen evolution tests showed that the ﬁlms could effectively enhance the corrosion resistance in SBF. According to the corrosion surface, it was concluded that after immersion for a period, the failure of the ﬁlms could be ascribed to the micro pores in the ﬁlms. Additional, the immersion tests indicated that the ﬁne structure of the ﬁlms could induce the growth of HAp in SBF and showed good bioactivity. Acknowledgements This work was supported by National Natural Science Fund of China (30870634) and National Basic Research Program of China (2008CB617509). References  G.L. Song, Corros. Sci. 49 (2007) 1696–1701.  F. Witte, V. Kaese, H. Haferkamp, E. Switzer, A.M. Lindenberg, C.J. Wirth, H. Windhagen, Biomaterials 26 (2005) 3557–3563.  L. Xu, G. Yu, E. Zhang, F. Pan, K. Yang, J. Biomed. Mater. Res. A 3 (2007) 703–711.  F. Witte, N. Hort, C. Vogt, S. Cohen, K.U. Kainer, R. Willumeit, F. Feyerabend, Curr. Opin. Solid State Mater. Sci. 12 (2008) 63–72.  H. Hermawan, D. Dubé, D. Mantovani, Acta Biomater. 5 (2010) 1693–1697.  J. Wang, L. Wang, S. Guan, S. Zhu, C. Ren, S. Hou, J. Mater. Sci.: Mater. Med. 21 (2010) 2001–2008.  L. Xu, PanF F., G. Yu, L. Yang, E. Zhang, K. Yang, Biomaterials 30 (2009) 1512–1523.  Z. Li, X. Gu, S. Lou, Y. Zheng, Biomaterials 29 (2008) 1329–1344.  C.L. Wen, S.K. Guan, L. Peng, C.X. Ren, X. Wang, Z.H. Hu, Appl. Surf. Sci. 255 (2009) 6433–6438.  Y. Song, S.X. Zhang, J.N. Li, C.L. Zhao, X.N. Zhang, Acta Biomater. 6 (2010) 1736–1742.  H.M. Wong, K.W.K. Yeung, K.O. Lam, V. Tam, P.K. Chu, K.D.K. Luk, K.M.C. Cheung, Biomaterials 31 (2010) 2084–2096.  J.Y. Chen, G.J. Wan, Y.X. Leng, P. Yang, H. Sun, J. Wang, N. Huang, Surf. Coat. Technol. 186 (2004) 270–276.  N. Huang, P. Yang, Y.X. Leng, J.Y. Chen, H. Sun, J. Wang, G.J. Wang, P.D. Ding, T.F. Xi, Y. Leng, Biomaterials 24 (2003) 2177–2187.  P. Yang, N. Huang, Y.X. Leng, J.Y. Chen, H. Sun, J. Wang, F. Chen, P.K. Chu, Surf. Coat. Technol. 156 (2002) 284–288.  B. O’Brien, W. Carroll, Acta Biomater. 5 (2009) 945–958.  T. Kokubo, H. Takadama, Biomaterials 27 (2006) 2907–2915.  F. Witte, J. Fischer, J. Nellesen, H. Crostack, V. Kaese, A. Pisch, F. Beckmanne, H. Windhagen, Biomaterials 27 (2006) 1013–1018.  G. Wu, X. Wang, K. Ding, Y. Zhou, X. Zeng, Mater. Charact. 60 (2009) 803–807.  F. Xiao, K. Tsuru, S. Hayakawa, A. Osaka, Thin Solid Film 441 (2003) 271–276.  S.H. Oh, R.R. Finõnes, C. Daraio, L.H. Chen, S. Jin, Biomaterials 26 (2005) 4938–4943.  F. Wang, M. Li, Y. Lu, S. Ge, J. Mater. Sci. 40 (2005) 2073–2076.