Original Article
Effect of TiO2 addition on surface microstructure and bioactivity of fluorapatite coatings deposited using Nd:YAG laser
Proc IMechE Part H: J Engineering in Medicine 2014, Vol. 228(4) 379–387 Ó IMechE 2014 Reprints and permissions: sagepub.co.uk/journalsPermissions.nav DOI: 10.1177/0954411914528307 pih.sagepub.com
Chi-Sheng Chien1,2, Yu-Sheng Ko3, Tsung-Yuan Kuo3, Tze-Yuan Liao4, Tzer-Min Lee5 and Ting-Fu Hong6
Abstract To study the effect of titania (TiO2) addition on the surface microstructure and bioactivity of fluorapatite coatings, fluorapatite was mixed with TiO2 in 1:0.5 (FA + 0.5TiO2), 1:0.8 (FA + 0.8TiO2), and 1:1 (FA + TiO2) ratios (wt%) and clad on Ti-6Al-4V substrates using an Nd:YAG laser system. The experimental results show that the penetration depth of the weld decreases with increasing TiO2 content. Moreover, the subgrain structure of the coating layer changes from a fine cellular-like structure to a cellular-dendrite-like structure as the amount of TiO2 increases. Consequently, as the proportion of TiO2 decreases (increase in fluorapatite content), the Ca/P ratio of the coating layer also decreases. The immersion of specimens into simulated body fluid resulted in the formation of individual apatite. With a lower Ca/P ratio before immersion, the growth of the apatite was faster and then the coating layer provided a better bioactivity. X-ray diffraction analysis results show that prior to simulated body fluid immersion, the coating layer in all three specimens was composed mainly of fluorapatite, CaTiO3, and Al2O3 phases. Following simulated body fluid immersion, a peak corresponding to hydroxycarbonated apatite appeared after 2 days in the FA + 0.5TiO2 and FA + 0.8TiO2 specimens and after 7 days in the FA + TiO2 specimen. Overall, the results show that although the bioactivity of the coating layer tended to decrease with increasing TiO2 content, in accordance with the above-mentioned ratios, the bioactivity of all three specimens remained generally good.
Keywords Laser cladding, fluorapatite, titania, Ti-6Al-4V, simulated body fluid
Date received: 9 February 2013; accepted: 18 February 2014
Introduction Hydroxyapatite (HA, Ca5(PO4)3OH) and other calcium phosphates are the main inorganic phases of human hard tissue.1,2 HA is frequently used as a coating material for surgical implants since its chemical, structural, and biological similarities to human bone prompt the formation of chemical bonds with the body tissue.3,4 However, when pure HA coating by sol-gel method is exposed to human body fluid, its relatively high dissolution rate presents long-term performance problems, resulting in the disintegration and detachment of the coating layer (CL) from the underlying implant.5,6 Moreover, HA coatings tend to crack and become detached from the metallic substrates during the deposition process due to the large mismatch between the thermal expansion coefficients of the two materials.7 Some weaknesses such as poor toughness
and low bonding strength have limited the usage of HA in the orthopedic and dentistry fields.8,9 It has been shown that these problems can be resolved by replacing 1
Chimei Foundation Hospital, Tainan, Taiwan, R.O.C. Department of Electrical Engineering, Southern Taiwan University of Science and Technology, Tainan, Taiwan, R.O.C. 3 Department of Mechanical Engineering, Southern Taiwan University of Science and Technology, Tainan, Taiwan, R.O.C. 4 Department of Materials Engineering, National Cheng Kung University, Tainan, Taiwan, R.O.C. 5 Institute of Oral Medicine, Medical College, National Cheng Kung University, Tainan, Taiwan, R.O.C. 6 Graduate Institute of Materials Engineering, National Pingtung University of Science and Technology, Pingtung, Taiwan, R.O.C. 2
Corresponding author: TY Kuo, Department of Mechanical Engineering, Southern Taiwan University of Science and Technology, Tainan 710, Taiwan, R.O.C. Email: [email protected]
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the OH21 ions in the HA structure with F21 ions to form fluorapatite (FA, Ca10(PO4)6F2). FA not only provides better protein adsorption and cell attachment capability than those of HA, but also improves alkaline phosphatase activity.10 FA has good phase stability even at higher temperatures.11,12 Its high chemical stability13 and low fluoride solubility could improve bone formation with comparable bioactivity and biocompatibility with that of HA.14 Finally, the F2 ions in the FA structure improve the mineralization and crystallization of calcium phosphate during the boneforming process.15,16 FA has thus gradually replaced HA as the preferred coating material for metal implants. However, pure FA coatings have high brittleness and low wear resistance and are therefore unsuitable for most artificial hard tissue applications.17 The mechanical properties of coating materials can be improved by adding bioinert ceramics as reinforcement within the coating.18,19 Coatings on a Ti alloy substrate with pure TiO2 have been shown to have better bonding strength than that of other biomaterial coatings while retaining good biocompatibility.20 Coating a metal implant with a thin intermediate layer of TiO2 before the deposition of the HA layer not only increases the bonding strength between the coating and the substrate, but also enhances the osteoblast adhesion of the medical implant and prompts more vigorous cell growth following implantation.21,22 Furthermore, it has been shown that the corrosion resistance of HA coatings increases with increasing thickness of the TiO2 intermediate layer.23 Accordingly, many studies have attempted to improve the mechanical properties, bioactivity, and corrosion resistance of surgical implant coatings by combining the advantages of both HA and TiO2 by adding TiO2 powder to HA coatings.15,20,24–26 However, the addition of TiO2 powder to FA coatings has received little attention in the literature. Moreover, in the studies mentioned above, the coating was generally deposited using sol-gel or plasma spraying techniques, which produce a relatively weak mechanical bond between the coating and substrate. Laser beams have coherence and directionality and generate a strong metallurgical bond between the coating and substrate.27–30 Accordingly, in this study, FA coatings with various levels of TiO2 are prepared and deposited on Ti-6Al-4V substrates using an Nd:YAG laser system. The coatings are prepared using anatase phase TiO2 since anatase has a high photocatalytic reaction efficiency31 and is highly conducive to the formation of apatite.32 Moreover, the Nd:YAG laser cladding method is chosen in preference to sol-gel or plasma spraying techniques due to the higher strength of the resulting metallurgical bond. Laser-clad specimens were immersed in simulated body fluid (SBF) for 2, 7, or 14 days. The surface properties and chemical compositions of the specimens were examined using optical microscopy (OM), scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), and X-ray diffraction (XRD) in
Figure 1. XRD pattern of FA powder. FA: fluorapatite.
order to investigate the effect of TiO2 addition on the microstructure and bioactivity of FA coatings.
Experiments TiO2 powder used in this study was supplied by Ishihara Inc., Japan. FA powder (3Ca3(PO4)2 + CaF2 ! Ca10(PO4)6F233) was prepared by mixing b-tricalcium phosphate (TCP) and CaF2 in a stoichiometric ratio of 3:1 followed by annealing in air for 3 h at a temperature of 1000 °C. The atomic structure of the asprepared FA powder was examined using XRD (Figure 1). Coating materials were prepared by mixing FA and TiO2 in 1:0.5, 1:0.8, and 1:1 ratios (wt%). The resulting powders, designated as FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2, respectively, were mixed with a polyvinyl alcohol binder material ((C2H4O)n) in a 1:1 ratio (wt%) and then thoroughly stirred until a slurry-like consistency was obtained. Substrates for the laser cladding process were prepared by machining Ti-6Al-4V alloy with the chemical composition shown in Table 1 into thin plates with dimensions of 100 mm 3 60 mm 3 3.8 mm. To remove any traces of surface oxidation, the plates were ground using 120-grit SiC paper, cleaned in an acetone solution, and dried in an electric oven. One slot (80 mm 3 40 mm 3 0.8 mm) was milled on the Ti-6Al-4V plate and refilled with the FA/TiO2 + binder slurry, and then the slurry was leveled by a stainless steel scraper.34 All the pre-placed specimens were dried in an oven under atmospheric conditions at 100 °C for 40 min. Once this pre-placed layer was dry, the specimens were then laser clad using an Nd:YAG laser (ROFIN CW025, 2500 W; Rofin Sinar Technologies Inc., Germany) set to continuous-wave mode with an output power of 850 W. The laser beam was guided to the workstation by an optical fiber (core diameter: 600 mm) and focused by a lens with a 120 mm focal length. The cladding process was performed in an Ar-shielded atmosphere (Ar flow rate: 25 L/min)
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Table 1. Chemical composition (wt%) of Ti-6Al-4V. Al
V
O
Fe
C
N
H
Ti
6.1
4.24
0.152
0.16
0.017
0.008
0.0006
Balance
Table 2. Inorganic ion concentrations of SBF and human blood plasma.
SBF (mM) Blood (mM)
Na +
K+
Ca2 +
Mg2 +
Cl2
HCO 3
HPO2 4
142.0 142.0
5.0 5.0
2.5 2.5
1.5 1.5
148.8 103.0
4.2 27.0
1.0 1.0
SBF: simulated body fluid.
Results and discussion Morphology and microstructure of weld beads
Figure 2. Experimental setup for Nd-YAG laser cladding process. CNC: Computer Numerical Control.
with a 5° laser incident angle, a 15-mm plus defocus length, and a 600-mm/min cladding travel speed. The laser spot size was about 3 mm in diameter. The experimental setup is shown schematically in Figure 2. The microstructures of the clad specimens were observed using OM and SEM. The Ca/P ratios of the various coatings were calculated by measuring the atomic percentages of Ca and P within the structure using a SEM (JEOL JSM-6390LV, JEOL Ltd., Japan) equipped with an EDS detector. After the microstructures of the various coatings had been examined, the specimens were cleaned, sterilized, and immersed in a standard SBF solution prepared in accordance with Kokubo’s protocol.35 The ion concentrations in the SBF solution and human blood plasma, respectively, are presented in Table 2. The specimens were soaked in SBF for 2, 7, and 14 days at a temperature of 37 6 0.1 °C and then examined for the nucleation and growth of apatite using SEM with EDS. In addition, the phases within the coatings were identified using XRD (Cu Ka radiation; Rigaku D/Max III.V, Rigaku Ltd., Japan). The XRD patterns were collected in a 2u range of 20°– 70° with a scanning rate of 2 °min21.
Figure 3 shows the morphologies and microstructures of the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 weld beads. In each case, the weld bead comprises a CL at the top of the bead and a transition layer (TL) at the bottom of the bead. Moreover, in each specimen, an obvious heat-affected zone (HAZ) formed in the substrate beneath the TL. A comparison of the three OM images indicates that the weld penetration depth decreases with increasing TiO2 content. The melting point of TiO2 (1855 °C) is higher than that of FA (1650 °C). Thus, with a higher proportion of TiO2, the average melting point of the FA/TiO2 coating material increases. As a result, a greater amount of laser energy is required to melt the coating material. Consequently, for a given laser output power, the laser energy transmitted to the substrate decreases with increasing TiO2 content and thus the weld depth decreases. The SEM images in Figure 3 show that a distinct interface formed between the CL and TL regions of the three weld beads. In each case, the morphologies tend to become coarser toward the surface of the bead. In the FA + 0.5TiO2 specimen, the microstructure in the middle of the CL has a fine cellular-like appearance. In contrast, in the FA + 0.8TiO2 and FA + TiO2 specimens, the microstructures are predominantly coarse cellular-like and cellular-dendrite-like structures, respectively. In other words, as the TiO2 content in the coating material increases, the subgrain structure changes progressively from a fine cellular-like structure to a cellular-dendrite-like structure. The microstructural change induced by TiO2 addition is the result of increased constitutional supercooling.36
In vitro bioactivity of laser-clad coatings Figure 4 shows the surface morphologies of the CL regions of the three laser-clad coatings before and after SBF immersion. It can be seen that after 2 days of SBF immersion, the surface of the FA + 0.5TiO2 specimen was covered with a dense arrangement of fine
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Figure 3. OM and SEM metallographs of weld bead, CL/TL interface, and mid-point position of CL in specimens with various levels of TiO2 addition. FA: fluorapatite; CL: coating layer; TL: transition layer; HAZ: heat-affected zone; OM: optical microscopy; SEM: scanning electron microscopy.
precipitates. Precipitates also formed on the surface of the FA + 0.8TiO2 specimen, but with a lower density than that for the FA + 0.5TiO2 specimen. The surface of the FA + TiO2 specimen contained only a very small number of precipitates. Following SBF immersion for 7 days, precipitates grew significantly with increased density. The size and density of the precipitates on the FA + 0.8TiO2 specimen after 7 days of immersion were similar to those for the FA + 0.5TiO2 specimen after 2 days of immersion. Similarly, the size and density of the precipitates on the FA + TiO2 specimen after 7 days of immersion resembled those for the FA + 0.8TiO2 specimen after 2 days of immersion. After 14 days of SBF immersion, the size of the precipitates on all of the specimen surfaces further increased slightly. The EDS spectrum of the CLs in the three specimens before immersion, as shown in Figure 5, indicates Ca, P, O, F, Ti, and Al as the main components. In the succeeding section, the XRD analysis shown in section ‘‘XRD patterns of CL surface before and after SBF immersion’’ comprises the FA, CaTiO3, and Al2O3 phases. Figure 6 presents the EDS analysis results for the precipitates formed on the surfaces of the three specimens following 14 days of immersion in SBF. In each case, dominant peaks corresponding to Ca, P, and O
appear. These elements are the main components of apatite.37–39 The XRD analysis discussed in section ‘‘XRD patterns of CL surface before and after SBF immersion’’ confirms that the precipitates are main forms of hydroxycarbonated apatite (HCA, Ca10(PO4)3(CO3)3(OH)2), which is an essential characteristic of bioactivity. The C peaks in the EDS spectra are most likely the result of adsorbed gas molecules (contaminants) on the specimen surface.40 In addition, the small quantity of Mg within the apatite phase is the result of ionic migration from the SBF solution.39 In general, the results presented in Figures 4 and 6 show that the amount of apatite phase formed, and thus, bioactivity on the FA/TiO2 surface decreases with increasing TiO2 content. The reduction in bioactivity is due to the addition of TiO2 to pure FA increasing the Ca/P ratio of the CL (see section ‘‘Ca/P ratio before and after SBF immersion’’), prompting a greater formation of the CaTiO3 phase during the laser cladding process (see section ‘‘XRD patterns of CL surface before and after SBF immersion’’).
Ca/P ratio before and after SBF immersion Figure 7 shows the Ca/P ratios of the CLs in the three specimens before and after SBF immersion. Prior to
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Figure 4. Surface morphologies of three coating layers following immersion in SBF for 0–14 days. FA: fluorapatite.
Figure 5. EDS analysis results of CLs for three specimens before SBF immersion: (a) FA + 0.5TiO2, (b) FA + 0.8TiO2, and (c) FA + TiO2. FA: fluorapatite.
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Figure 6. EDS mapping analysis results of CLs for three specimens following immersion in SBF for 14 days: (a) FA + 0.5TiO2, (b) FA + 0.8TiO2, and (c) FA + TiO2. FA: fluorapatite.
Figure 7. Ca/P ratios of three coating layers before and after SBF immersion. FA: fluorapatite.
immersion, the Ca/P ratio had a relatively high value in all three specimens. Furthermore, the Ca/P ratio increased with increasing TiO2 content. However, after 2 days of immersion, the Ca/P ratio fell markedly in all three specimens. Specifically, the Ca/P ratio for the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens decreased from 6.32 to 1.95, from 8.38 to 2.13, and from 9.02 to 4.43, respectively. In other words, the higher the amount of FA in the coating (i.e. the lower the amount of TiO2), the more rapidly the Ca/P ratio dropped toward the optimal bioactivity value of 1.67. After 7 days of SBF immersion, the Ca/P ratio fell to \ 1.67 in all three specimens. However, in each case, the Ca/P ratio increased slightly when the immersion time was extended from 7 to 14 days. Specifically, the final Ca/P ratios for the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens were 1.71, 1.83, and 2.17, respectively. The SBF immersion test is primarily used to explore apatite growth conditions for a variety of coating
material (layer) surfaces. It is also the determining factor for bioactivity. In this study, a correlation between the Ca/P ratio (6.32–9.02) of the specimen CL before soaking and the growth of apatite after soaking was found. During SBF soaking, it was observed that for a lower Ca/P ratio of the CL, apatite growth was faster, demonstrating better bioactivity. This is due to the superiority of the FA bioactivity to that of TiO2. Higher FA content in the coating material leads to a better environment for bone growth on the CL after laser cladding. The Ca and P in the coating material come from FA. Therefore, the Ca/P ratio on the CL before SBF immersion depends on the FA (Ca/P = 1.67) content, in addition to the P evaporation due to laser irradiation (this effect significantly increases the Ca/P ratio). The higher the FA content, the closer the Ca/P ratio in the CL is to that of FA, which improves the bioactivity during SBF soaking. The results presented in Figure 7 show that for all specimens, the Ca/P ratio increased slightly when the immersion time was increased beyond that required to obtain the ideal value of 1.67. Previous investigations20,39,41 into the mechanism of apatite formation on the surface of NaOH-treated TiO2 or HA/TiO2 coatings following immersion in SBF reported similar findings. In the immersion process, an amorphous calcium phosphate (ACP) phase with a low Ca/P ratio (\ 1.67) initially forms on the coating surface. The ACP phase grows gradually as the soaking time is increased and develops a granular morphology. Since apatite is the stable phase in the body environment, the ACP phase spontaneously transforms into bone-like apatite as the soaking time is further increased due to the consumption of calcium and phosphate ions from the SBF. Thus, the Ca/P ratio increases following a prolonged immersion time. In addition, when the NaOH-treated
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substrate coatings mentioned above are soaked in SBF, Ti–OH groups form on their surface. These Ti–OH groups can induce ACP precipitation on the surface of the NaOH-treated TiO2 or HA/TiO2 coatings. However, after the substrate in this study underwent laser cladding, the TiO2 completely dissolved, leaving no residue. Furthermore, no ACP was detected in the XRD patterns, as shown in section ‘‘XRD patterns of CL surface before and after SBF immersion.’’ The Ca/ P ratio thus increased after it had decreased. This phenomenon requires further research.
XRD patterns of CL surface before and after SBF immersion Figures 8–10 show the XRD analysis results for the CL surfaces of the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens before and after SBF immersion. Prior to SBF immersion, the three coatings consisted mainly of FA, CaTiO3, and Al2O3. In other words, following the laser cladding process, some of the FA in the original coating material remained. The presence of CaTiO3 can be attributed to the interaction between FA (Ca-rich phase) and TiO2.42 The amount of CaTiO3 formed in the laser cladding process increases with increasing TiO2 content in the original coating material. Thus, the amount of residual FA decreases as the TiO2 content increases. Although some studies43,44 have reported that CaTiO3 is conducive to the growth of apatite, its effectiveness in prompting apatite precipitation and growth is not as great as that of FA. As a result, the bioactivity of the present specimens decreases with increasing TiO2 content due to the corresponding reduction in FA. Nonetheless, all three specimens still have good bioactivity (i.e. Ca/P ratios of 1.71, 1.83, and 2.17 wt%, respectively, following SBF immersion for 14 days, see Figure 7). The Al2O3 peak observed in the three specimens prior to SBF immersion is the result of a reaction between the O atoms from the coating materials and the Al atoms released from the Ti-6Al-4V substrate during the laser cladding process. A comparison of the XRD patterns of the three specimens indicates that the intensity of the Al2O3 peak increases with decreasing TiO2 content. The average melting point of the FA/ TiO2 coating material decreases as the proportion of TiO2 decreases. Thus, a greater amount of the laser input energy is transmitted to the Ti-6Al-4V substrate. Consequently, a greater number of Al atoms are released by the substrate during the cladding process, and thus, the amount of Al2O3 increases. The XRD patterns for the FA + 0.5TiO2 and FA + 0.8TiO2 specimens show an obvious peak corresponding to HCA (Joint Committee on Powder Diffraction Standards (JCPDS) card no. 19-272) following 2 days of immersion in SBF. A similar peak appears in the XRD pattern of the FA + TiO2 specimen after 7 days. As the immersion time was increased,
Figure 8. XRD analysis results for FA + 0.5TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
Figure 9. XRD analysis results for FA + 0.8TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
the intensity of the HCA peak increased notably in all three specimens. The strengthening of the HCA peak was accompanied by a weakening of the FA, CaTiO3, and Al2O3 peaks. This suggests that the CL becomes progressively covered by HCA as the SBF immersion time increases. Thus, after 14 days of immersion, Al2O3 peaks were no longer detected in any of the specimens. In general, the XRD results presented in Figures 8– 10 show that the specimens with a greater amount of FA in the coating material had improved bioactivity. F2 has a high affinity to oxygen ions.3 Thus, as the proportion of FA in the coating material increases, the number of F2 ions available to react with the OH2 ions to form negatively charged F–OH and Ti–OH groups on the surface increases. The 2OH groups promote the attachment and adhesion of cells.45 Thus, the HCA layer formed on the surface of coatings with a
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Proc IMechE Part H: J Engineering in Medicine 228(4) FA, CaTiO3, and Al2O3 peaks became weaker and were replaced by a dominant peak corresponding to HCA. The HCA peak appeared following a shorter immersion time in the specimens with lower TiO2 content. Thus, based on the FA/TiO2 ratios used in this study, it can be inferred that the bioactivity of the FA/TiO2 coatings increases with decreasing TiO2 content. Declaration of conflicting interests The authors declare that there is no conflict of interest. Funding
Figure 10. XRD analysis results for FA + TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
greater proportion of FA is more conducive to the attachment of proteins such as collagen, fibronectin, and vitronectin.46 This in turn enhances the osteoblast adhesion of the coating47 and improves the bonding strength between the implant and the surrounding tissue.48
Conclusion 1.
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TiO2 has a higher melting point than that of FA. Thus, as the amount of TiO2 in the coating material increases, the amount of laser energy transmitted to the substrate decreases. As a result, the penetration depth of the weld bead decreases. Furthermore, a higher TiO2 content increases the supercooling effect. Consequently, the subgrain structure of the CL changes from a fine cellularlike structure to a cellular-dendrite-like structure as the amount of TiO2 in the coating material increases. With SBF immersion, the rate at which apatite precipitated on the coating surface increased with decreasing TiO2 content. However, following prolonged SBF immersion (14 days), all three specimens were covered with a dense arrangement of HCA precipitates, indicating good bioactivity. The Ca/P ratios of the three specimens dropped rapidly following SBF immersion. The rate at which the Ca/P ratio falls increased with decreasing TiO2 content. However, following 7 days of SBF immersion, the Ca/P ratios of all three specimens were close to the ideal bioactivity value of 1.67. After 14 days, the Ca/P ratios further increased slightly due to the spontaneous transformation of the ACP phase into bone-like apatite. The XRD analysis results show that the laser-clad coatings consisted mainly of FA, CaTiO3, and Al2O3 phases. Following immersion in SBF, the
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photocatalytic activity of TiO2 thin films prepared by liquid phase deposition. J Phys Chem B 2003; 107: 13871–13879. Uchida M, Kim HM, Kokubo T, et al. Structural dependence of apatite formation on titania gels in a simulated body fluid. J Biomed Mater Res A 2003; 64: 164–170. Kim HW, Kong YM, Koh YH, et al. Pressureless sintering and mechanical and biological properties of fluorhydroxyapatite composites with zirconia. J Am Ceram Soc 2003; 86: 2019–2026. Chien CS, Han TJ, Hong TF, et al. Effects of different adhesive of hydroxyapatite power on weld morphology, Ca/P value and hardness of cladding layer by Nd-YAG deposition on Ti-6Al-4V substrate. Mater Trans 2009; 50: 2852–2857. Kokubo T and Takadama H. How useful is SBF in predicting in vivo bone bioactivity. Biomaterials 2006; 27: 2907–2915. Kuo S. Welding metallurgy. New York: John Wiley & Sons, 1987. Han JY, Yu ZT and Zhou L. Hydroxyapatite/titania composite bioactivity coating processed by sol-gel method. Appl Surf Sci 2008; 255: 455–458. Kurella A and Dahotre NB. Laser induced hierarchical calcium phosphate structures. Acta Biomater 2006; 2: 677–683. Kokubo T, Kim HM and Kawashita M. Novel bioactive materials with different mechanical properties. Biomaterials 2003; 24: 2161–2175. Trepanier C, Tabrizian M, Yahia LH, et al. Effect of modification of oxide layer on NiTi stent corrosion resistance. J Biomed Mater Res 1998; 43: 433–440. Takadama H, Kim HM, Kokubo T, et al. An X-ray photoelectron spectroscopy study of the process of apatite formation on bioactive titanium metal. J Biomed Mater Res 2001; 55: 185–193. Lusquin˜os F, De Carlos A, Pou J, et al. Calcium phosphate coatings obtained by Nd:YAG laser cladding: physicochemical and biologic properties. J Biomed Mater Res A 2003; 64: 630–637. Webster TJ, Ergun C, Doremus RH, et al. Increased osteoblast adhesion on titanium coated hydroxylapatite that forms CaTiO3. J Biomed Mater Res A 2003; 67: 975–980. Asami K, Saito K, Ohtsu N, et al. Titanium-implanted CaTiO3 films and their changes in Hanks’ solution. Surf Interface Anal 2003; 35: 483–488. Zhu X, Ohtsubo M, Bohmer RM, et al. Adhesion-dependent cell cycle progression linked to the expression of cyclin D1, activation of cyclin E-cdk2, and phosphorylation of the retinoblastoma protein. J Cell Biol 1996; 133: 391–403. Jona´sova´ L, Mu¨ller FA, Helebrant A, et al. Biomimetic apatite formation on chemically treated titanium. Biomaterials 2004; 25: 1187–1194. Webster TJ, Ergun C, Doremus RH, et al. Specific proteins mediate enhanced osteoblast adhesion on nanophase ceramics. J Biomed Mater Res 2000; 51: 475–483. Hench LL. The story of Bioglass. J Mater Sci Mater Med 2006; 17: 967–978.
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Effect of TiO2 addition on surface microstructure and bioactivity of fluorapatite coatings deposited using Nd:YAG laser
Proc IMechE Part H: J Engineering in Medicine 2014, Vol. 228(4) 379–387 Ó IMechE 2014 Reprints and permissions: sagepub.co.uk/journalsPermissions.nav DOI: 10.1177/0954411914528307 pih.sagepub.com
Chi-Sheng Chien1,2, Yu-Sheng Ko3, Tsung-Yuan Kuo3, Tze-Yuan Liao4, Tzer-Min Lee5 and Ting-Fu Hong6
Abstract To study the effect of titania (TiO2) addition on the surface microstructure and bioactivity of fluorapatite coatings, fluorapatite was mixed with TiO2 in 1:0.5 (FA + 0.5TiO2), 1:0.8 (FA + 0.8TiO2), and 1:1 (FA + TiO2) ratios (wt%) and clad on Ti-6Al-4V substrates using an Nd:YAG laser system. The experimental results show that the penetration depth of the weld decreases with increasing TiO2 content. Moreover, the subgrain structure of the coating layer changes from a fine cellular-like structure to a cellular-dendrite-like structure as the amount of TiO2 increases. Consequently, as the proportion of TiO2 decreases (increase in fluorapatite content), the Ca/P ratio of the coating layer also decreases. The immersion of specimens into simulated body fluid resulted in the formation of individual apatite. With a lower Ca/P ratio before immersion, the growth of the apatite was faster and then the coating layer provided a better bioactivity. X-ray diffraction analysis results show that prior to simulated body fluid immersion, the coating layer in all three specimens was composed mainly of fluorapatite, CaTiO3, and Al2O3 phases. Following simulated body fluid immersion, a peak corresponding to hydroxycarbonated apatite appeared after 2 days in the FA + 0.5TiO2 and FA + 0.8TiO2 specimens and after 7 days in the FA + TiO2 specimen. Overall, the results show that although the bioactivity of the coating layer tended to decrease with increasing TiO2 content, in accordance with the above-mentioned ratios, the bioactivity of all three specimens remained generally good.
Keywords Laser cladding, fluorapatite, titania, Ti-6Al-4V, simulated body fluid
Date received: 9 February 2013; accepted: 18 February 2014
Introduction Hydroxyapatite (HA, Ca5(PO4)3OH) and other calcium phosphates are the main inorganic phases of human hard tissue.1,2 HA is frequently used as a coating material for surgical implants since its chemical, structural, and biological similarities to human bone prompt the formation of chemical bonds with the body tissue.3,4 However, when pure HA coating by sol-gel method is exposed to human body fluid, its relatively high dissolution rate presents long-term performance problems, resulting in the disintegration and detachment of the coating layer (CL) from the underlying implant.5,6 Moreover, HA coatings tend to crack and become detached from the metallic substrates during the deposition process due to the large mismatch between the thermal expansion coefficients of the two materials.7 Some weaknesses such as poor toughness
and low bonding strength have limited the usage of HA in the orthopedic and dentistry fields.8,9 It has been shown that these problems can be resolved by replacing 1
Chimei Foundation Hospital, Tainan, Taiwan, R.O.C. Department of Electrical Engineering, Southern Taiwan University of Science and Technology, Tainan, Taiwan, R.O.C. 3 Department of Mechanical Engineering, Southern Taiwan University of Science and Technology, Tainan, Taiwan, R.O.C. 4 Department of Materials Engineering, National Cheng Kung University, Tainan, Taiwan, R.O.C. 5 Institute of Oral Medicine, Medical College, National Cheng Kung University, Tainan, Taiwan, R.O.C. 6 Graduate Institute of Materials Engineering, National Pingtung University of Science and Technology, Pingtung, Taiwan, R.O.C. 2
Corresponding author: TY Kuo, Department of Mechanical Engineering, Southern Taiwan University of Science and Technology, Tainan 710, Taiwan, R.O.C. Email: [email protected]
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the OH21 ions in the HA structure with F21 ions to form fluorapatite (FA, Ca10(PO4)6F2). FA not only provides better protein adsorption and cell attachment capability than those of HA, but also improves alkaline phosphatase activity.10 FA has good phase stability even at higher temperatures.11,12 Its high chemical stability13 and low fluoride solubility could improve bone formation with comparable bioactivity and biocompatibility with that of HA.14 Finally, the F2 ions in the FA structure improve the mineralization and crystallization of calcium phosphate during the boneforming process.15,16 FA has thus gradually replaced HA as the preferred coating material for metal implants. However, pure FA coatings have high brittleness and low wear resistance and are therefore unsuitable for most artificial hard tissue applications.17 The mechanical properties of coating materials can be improved by adding bioinert ceramics as reinforcement within the coating.18,19 Coatings on a Ti alloy substrate with pure TiO2 have been shown to have better bonding strength than that of other biomaterial coatings while retaining good biocompatibility.20 Coating a metal implant with a thin intermediate layer of TiO2 before the deposition of the HA layer not only increases the bonding strength between the coating and the substrate, but also enhances the osteoblast adhesion of the medical implant and prompts more vigorous cell growth following implantation.21,22 Furthermore, it has been shown that the corrosion resistance of HA coatings increases with increasing thickness of the TiO2 intermediate layer.23 Accordingly, many studies have attempted to improve the mechanical properties, bioactivity, and corrosion resistance of surgical implant coatings by combining the advantages of both HA and TiO2 by adding TiO2 powder to HA coatings.15,20,24–26 However, the addition of TiO2 powder to FA coatings has received little attention in the literature. Moreover, in the studies mentioned above, the coating was generally deposited using sol-gel or plasma spraying techniques, which produce a relatively weak mechanical bond between the coating and substrate. Laser beams have coherence and directionality and generate a strong metallurgical bond between the coating and substrate.27–30 Accordingly, in this study, FA coatings with various levels of TiO2 are prepared and deposited on Ti-6Al-4V substrates using an Nd:YAG laser system. The coatings are prepared using anatase phase TiO2 since anatase has a high photocatalytic reaction efficiency31 and is highly conducive to the formation of apatite.32 Moreover, the Nd:YAG laser cladding method is chosen in preference to sol-gel or plasma spraying techniques due to the higher strength of the resulting metallurgical bond. Laser-clad specimens were immersed in simulated body fluid (SBF) for 2, 7, or 14 days. The surface properties and chemical compositions of the specimens were examined using optical microscopy (OM), scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), and X-ray diffraction (XRD) in
Figure 1. XRD pattern of FA powder. FA: fluorapatite.
order to investigate the effect of TiO2 addition on the microstructure and bioactivity of FA coatings.
Experiments TiO2 powder used in this study was supplied by Ishihara Inc., Japan. FA powder (3Ca3(PO4)2 + CaF2 ! Ca10(PO4)6F233) was prepared by mixing b-tricalcium phosphate (TCP) and CaF2 in a stoichiometric ratio of 3:1 followed by annealing in air for 3 h at a temperature of 1000 °C. The atomic structure of the asprepared FA powder was examined using XRD (Figure 1). Coating materials were prepared by mixing FA and TiO2 in 1:0.5, 1:0.8, and 1:1 ratios (wt%). The resulting powders, designated as FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2, respectively, were mixed with a polyvinyl alcohol binder material ((C2H4O)n) in a 1:1 ratio (wt%) and then thoroughly stirred until a slurry-like consistency was obtained. Substrates for the laser cladding process were prepared by machining Ti-6Al-4V alloy with the chemical composition shown in Table 1 into thin plates with dimensions of 100 mm 3 60 mm 3 3.8 mm. To remove any traces of surface oxidation, the plates were ground using 120-grit SiC paper, cleaned in an acetone solution, and dried in an electric oven. One slot (80 mm 3 40 mm 3 0.8 mm) was milled on the Ti-6Al-4V plate and refilled with the FA/TiO2 + binder slurry, and then the slurry was leveled by a stainless steel scraper.34 All the pre-placed specimens were dried in an oven under atmospheric conditions at 100 °C for 40 min. Once this pre-placed layer was dry, the specimens were then laser clad using an Nd:YAG laser (ROFIN CW025, 2500 W; Rofin Sinar Technologies Inc., Germany) set to continuous-wave mode with an output power of 850 W. The laser beam was guided to the workstation by an optical fiber (core diameter: 600 mm) and focused by a lens with a 120 mm focal length. The cladding process was performed in an Ar-shielded atmosphere (Ar flow rate: 25 L/min)
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Table 1. Chemical composition (wt%) of Ti-6Al-4V. Al
V
O
Fe
C
N
H
Ti
6.1
4.24
0.152
0.16
0.017
0.008
0.0006
Balance
Table 2. Inorganic ion concentrations of SBF and human blood plasma.
SBF (mM) Blood (mM)
Na +
K+
Ca2 +
Mg2 +
Cl2
HCO 3
HPO2 4
142.0 142.0
5.0 5.0
2.5 2.5
1.5 1.5
148.8 103.0
4.2 27.0
1.0 1.0
SBF: simulated body fluid.
Results and discussion Morphology and microstructure of weld beads
Figure 2. Experimental setup for Nd-YAG laser cladding process. CNC: Computer Numerical Control.
with a 5° laser incident angle, a 15-mm plus defocus length, and a 600-mm/min cladding travel speed. The laser spot size was about 3 mm in diameter. The experimental setup is shown schematically in Figure 2. The microstructures of the clad specimens were observed using OM and SEM. The Ca/P ratios of the various coatings were calculated by measuring the atomic percentages of Ca and P within the structure using a SEM (JEOL JSM-6390LV, JEOL Ltd., Japan) equipped with an EDS detector. After the microstructures of the various coatings had been examined, the specimens were cleaned, sterilized, and immersed in a standard SBF solution prepared in accordance with Kokubo’s protocol.35 The ion concentrations in the SBF solution and human blood plasma, respectively, are presented in Table 2. The specimens were soaked in SBF for 2, 7, and 14 days at a temperature of 37 6 0.1 °C and then examined for the nucleation and growth of apatite using SEM with EDS. In addition, the phases within the coatings were identified using XRD (Cu Ka radiation; Rigaku D/Max III.V, Rigaku Ltd., Japan). The XRD patterns were collected in a 2u range of 20°– 70° with a scanning rate of 2 °min21.
Figure 3 shows the morphologies and microstructures of the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 weld beads. In each case, the weld bead comprises a CL at the top of the bead and a transition layer (TL) at the bottom of the bead. Moreover, in each specimen, an obvious heat-affected zone (HAZ) formed in the substrate beneath the TL. A comparison of the three OM images indicates that the weld penetration depth decreases with increasing TiO2 content. The melting point of TiO2 (1855 °C) is higher than that of FA (1650 °C). Thus, with a higher proportion of TiO2, the average melting point of the FA/TiO2 coating material increases. As a result, a greater amount of laser energy is required to melt the coating material. Consequently, for a given laser output power, the laser energy transmitted to the substrate decreases with increasing TiO2 content and thus the weld depth decreases. The SEM images in Figure 3 show that a distinct interface formed between the CL and TL regions of the three weld beads. In each case, the morphologies tend to become coarser toward the surface of the bead. In the FA + 0.5TiO2 specimen, the microstructure in the middle of the CL has a fine cellular-like appearance. In contrast, in the FA + 0.8TiO2 and FA + TiO2 specimens, the microstructures are predominantly coarse cellular-like and cellular-dendrite-like structures, respectively. In other words, as the TiO2 content in the coating material increases, the subgrain structure changes progressively from a fine cellular-like structure to a cellular-dendrite-like structure. The microstructural change induced by TiO2 addition is the result of increased constitutional supercooling.36
In vitro bioactivity of laser-clad coatings Figure 4 shows the surface morphologies of the CL regions of the three laser-clad coatings before and after SBF immersion. It can be seen that after 2 days of SBF immersion, the surface of the FA + 0.5TiO2 specimen was covered with a dense arrangement of fine
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Figure 3. OM and SEM metallographs of weld bead, CL/TL interface, and mid-point position of CL in specimens with various levels of TiO2 addition. FA: fluorapatite; CL: coating layer; TL: transition layer; HAZ: heat-affected zone; OM: optical microscopy; SEM: scanning electron microscopy.
precipitates. Precipitates also formed on the surface of the FA + 0.8TiO2 specimen, but with a lower density than that for the FA + 0.5TiO2 specimen. The surface of the FA + TiO2 specimen contained only a very small number of precipitates. Following SBF immersion for 7 days, precipitates grew significantly with increased density. The size and density of the precipitates on the FA + 0.8TiO2 specimen after 7 days of immersion were similar to those for the FA + 0.5TiO2 specimen after 2 days of immersion. Similarly, the size and density of the precipitates on the FA + TiO2 specimen after 7 days of immersion resembled those for the FA + 0.8TiO2 specimen after 2 days of immersion. After 14 days of SBF immersion, the size of the precipitates on all of the specimen surfaces further increased slightly. The EDS spectrum of the CLs in the three specimens before immersion, as shown in Figure 5, indicates Ca, P, O, F, Ti, and Al as the main components. In the succeeding section, the XRD analysis shown in section ‘‘XRD patterns of CL surface before and after SBF immersion’’ comprises the FA, CaTiO3, and Al2O3 phases. Figure 6 presents the EDS analysis results for the precipitates formed on the surfaces of the three specimens following 14 days of immersion in SBF. In each case, dominant peaks corresponding to Ca, P, and O
appear. These elements are the main components of apatite.37–39 The XRD analysis discussed in section ‘‘XRD patterns of CL surface before and after SBF immersion’’ confirms that the precipitates are main forms of hydroxycarbonated apatite (HCA, Ca10(PO4)3(CO3)3(OH)2), which is an essential characteristic of bioactivity. The C peaks in the EDS spectra are most likely the result of adsorbed gas molecules (contaminants) on the specimen surface.40 In addition, the small quantity of Mg within the apatite phase is the result of ionic migration from the SBF solution.39 In general, the results presented in Figures 4 and 6 show that the amount of apatite phase formed, and thus, bioactivity on the FA/TiO2 surface decreases with increasing TiO2 content. The reduction in bioactivity is due to the addition of TiO2 to pure FA increasing the Ca/P ratio of the CL (see section ‘‘Ca/P ratio before and after SBF immersion’’), prompting a greater formation of the CaTiO3 phase during the laser cladding process (see section ‘‘XRD patterns of CL surface before and after SBF immersion’’).
Ca/P ratio before and after SBF immersion Figure 7 shows the Ca/P ratios of the CLs in the three specimens before and after SBF immersion. Prior to
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Figure 4. Surface morphologies of three coating layers following immersion in SBF for 0–14 days. FA: fluorapatite.
Figure 5. EDS analysis results of CLs for three specimens before SBF immersion: (a) FA + 0.5TiO2, (b) FA + 0.8TiO2, and (c) FA + TiO2. FA: fluorapatite.
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Figure 6. EDS mapping analysis results of CLs for three specimens following immersion in SBF for 14 days: (a) FA + 0.5TiO2, (b) FA + 0.8TiO2, and (c) FA + TiO2. FA: fluorapatite.
Figure 7. Ca/P ratios of three coating layers before and after SBF immersion. FA: fluorapatite.
immersion, the Ca/P ratio had a relatively high value in all three specimens. Furthermore, the Ca/P ratio increased with increasing TiO2 content. However, after 2 days of immersion, the Ca/P ratio fell markedly in all three specimens. Specifically, the Ca/P ratio for the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens decreased from 6.32 to 1.95, from 8.38 to 2.13, and from 9.02 to 4.43, respectively. In other words, the higher the amount of FA in the coating (i.e. the lower the amount of TiO2), the more rapidly the Ca/P ratio dropped toward the optimal bioactivity value of 1.67. After 7 days of SBF immersion, the Ca/P ratio fell to \ 1.67 in all three specimens. However, in each case, the Ca/P ratio increased slightly when the immersion time was extended from 7 to 14 days. Specifically, the final Ca/P ratios for the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens were 1.71, 1.83, and 2.17, respectively. The SBF immersion test is primarily used to explore apatite growth conditions for a variety of coating
material (layer) surfaces. It is also the determining factor for bioactivity. In this study, a correlation between the Ca/P ratio (6.32–9.02) of the specimen CL before soaking and the growth of apatite after soaking was found. During SBF soaking, it was observed that for a lower Ca/P ratio of the CL, apatite growth was faster, demonstrating better bioactivity. This is due to the superiority of the FA bioactivity to that of TiO2. Higher FA content in the coating material leads to a better environment for bone growth on the CL after laser cladding. The Ca and P in the coating material come from FA. Therefore, the Ca/P ratio on the CL before SBF immersion depends on the FA (Ca/P = 1.67) content, in addition to the P evaporation due to laser irradiation (this effect significantly increases the Ca/P ratio). The higher the FA content, the closer the Ca/P ratio in the CL is to that of FA, which improves the bioactivity during SBF soaking. The results presented in Figure 7 show that for all specimens, the Ca/P ratio increased slightly when the immersion time was increased beyond that required to obtain the ideal value of 1.67. Previous investigations20,39,41 into the mechanism of apatite formation on the surface of NaOH-treated TiO2 or HA/TiO2 coatings following immersion in SBF reported similar findings. In the immersion process, an amorphous calcium phosphate (ACP) phase with a low Ca/P ratio (\ 1.67) initially forms on the coating surface. The ACP phase grows gradually as the soaking time is increased and develops a granular morphology. Since apatite is the stable phase in the body environment, the ACP phase spontaneously transforms into bone-like apatite as the soaking time is further increased due to the consumption of calcium and phosphate ions from the SBF. Thus, the Ca/P ratio increases following a prolonged immersion time. In addition, when the NaOH-treated
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substrate coatings mentioned above are soaked in SBF, Ti–OH groups form on their surface. These Ti–OH groups can induce ACP precipitation on the surface of the NaOH-treated TiO2 or HA/TiO2 coatings. However, after the substrate in this study underwent laser cladding, the TiO2 completely dissolved, leaving no residue. Furthermore, no ACP was detected in the XRD patterns, as shown in section ‘‘XRD patterns of CL surface before and after SBF immersion.’’ The Ca/ P ratio thus increased after it had decreased. This phenomenon requires further research.
XRD patterns of CL surface before and after SBF immersion Figures 8–10 show the XRD analysis results for the CL surfaces of the FA + 0.5TiO2, FA + 0.8TiO2, and FA + TiO2 specimens before and after SBF immersion. Prior to SBF immersion, the three coatings consisted mainly of FA, CaTiO3, and Al2O3. In other words, following the laser cladding process, some of the FA in the original coating material remained. The presence of CaTiO3 can be attributed to the interaction between FA (Ca-rich phase) and TiO2.42 The amount of CaTiO3 formed in the laser cladding process increases with increasing TiO2 content in the original coating material. Thus, the amount of residual FA decreases as the TiO2 content increases. Although some studies43,44 have reported that CaTiO3 is conducive to the growth of apatite, its effectiveness in prompting apatite precipitation and growth is not as great as that of FA. As a result, the bioactivity of the present specimens decreases with increasing TiO2 content due to the corresponding reduction in FA. Nonetheless, all three specimens still have good bioactivity (i.e. Ca/P ratios of 1.71, 1.83, and 2.17 wt%, respectively, following SBF immersion for 14 days, see Figure 7). The Al2O3 peak observed in the three specimens prior to SBF immersion is the result of a reaction between the O atoms from the coating materials and the Al atoms released from the Ti-6Al-4V substrate during the laser cladding process. A comparison of the XRD patterns of the three specimens indicates that the intensity of the Al2O3 peak increases with decreasing TiO2 content. The average melting point of the FA/ TiO2 coating material decreases as the proportion of TiO2 decreases. Thus, a greater amount of the laser input energy is transmitted to the Ti-6Al-4V substrate. Consequently, a greater number of Al atoms are released by the substrate during the cladding process, and thus, the amount of Al2O3 increases. The XRD patterns for the FA + 0.5TiO2 and FA + 0.8TiO2 specimens show an obvious peak corresponding to HCA (Joint Committee on Powder Diffraction Standards (JCPDS) card no. 19-272) following 2 days of immersion in SBF. A similar peak appears in the XRD pattern of the FA + TiO2 specimen after 7 days. As the immersion time was increased,
Figure 8. XRD analysis results for FA + 0.5TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
Figure 9. XRD analysis results for FA + 0.8TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
the intensity of the HCA peak increased notably in all three specimens. The strengthening of the HCA peak was accompanied by a weakening of the FA, CaTiO3, and Al2O3 peaks. This suggests that the CL becomes progressively covered by HCA as the SBF immersion time increases. Thus, after 14 days of immersion, Al2O3 peaks were no longer detected in any of the specimens. In general, the XRD results presented in Figures 8– 10 show that the specimens with a greater amount of FA in the coating material had improved bioactivity. F2 has a high affinity to oxygen ions.3 Thus, as the proportion of FA in the coating material increases, the number of F2 ions available to react with the OH2 ions to form negatively charged F–OH and Ti–OH groups on the surface increases. The 2OH groups promote the attachment and adhesion of cells.45 Thus, the HCA layer formed on the surface of coatings with a
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Proc IMechE Part H: J Engineering in Medicine 228(4) FA, CaTiO3, and Al2O3 peaks became weaker and were replaced by a dominant peak corresponding to HCA. The HCA peak appeared following a shorter immersion time in the specimens with lower TiO2 content. Thus, based on the FA/TiO2 ratios used in this study, it can be inferred that the bioactivity of the FA/TiO2 coatings increases with decreasing TiO2 content. Declaration of conflicting interests The authors declare that there is no conflict of interest. Funding
Figure 10. XRD analysis results for FA + TiO2 specimen following various SBF immersion times. FA: fluorapatite; HCA: hydroxycarbonated apatite.
greater proportion of FA is more conducive to the attachment of proteins such as collagen, fibronectin, and vitronectin.46 This in turn enhances the osteoblast adhesion of the coating47 and improves the bonding strength between the implant and the surrounding tissue.48
Conclusion 1.
2.
3.
4.
TiO2 has a higher melting point than that of FA. Thus, as the amount of TiO2 in the coating material increases, the amount of laser energy transmitted to the substrate decreases. As a result, the penetration depth of the weld bead decreases. Furthermore, a higher TiO2 content increases the supercooling effect. Consequently, the subgrain structure of the CL changes from a fine cellularlike structure to a cellular-dendrite-like structure as the amount of TiO2 in the coating material increases. With SBF immersion, the rate at which apatite precipitated on the coating surface increased with decreasing TiO2 content. However, following prolonged SBF immersion (14 days), all three specimens were covered with a dense arrangement of HCA precipitates, indicating good bioactivity. The Ca/P ratios of the three specimens dropped rapidly following SBF immersion. The rate at which the Ca/P ratio falls increased with decreasing TiO2 content. However, following 7 days of SBF immersion, the Ca/P ratios of all three specimens were close to the ideal bioactivity value of 1.67. After 14 days, the Ca/P ratios further increased slightly due to the spontaneous transformation of the ACP phase into bone-like apatite. The XRD analysis results show that the laser-clad coatings consisted mainly of FA, CaTiO3, and Al2O3 phases. Following immersion in SBF, the
This study was financially supported by the Chimei Foundation Hospital, Taiwan, R.O.C., under grant 110990223, and the National Science Council, Taiwan, R.O.C., under grant NSC 100-2622-E-218-002-CC3. References 1. Peru L and Daculsi G. Synthetic calcium phosphates: models for biological crystals? Clin Mater 1994; 15: 267–272. 2. Jarcho M. Calcium phosphate ceramics as hard tissue prosthetics. Clin Orthop Relat Res 1981; 157: 259–278. 3. Sivakumar M and Manjubala I. Preparation of hydroxyapatite/fluoroapatite–zirconia composites using Indian corals for biomedical applications. Mater Lett 2001; 50: 199–205. 4. Kumar RR and Wang M. Functionally graded bioactive coatings of hydroxyapatite/titanium oxide composite system. Mater Lett 2002; 55: 133–137. 5. Gineste L, Gineste M, Ranz X, et al. Degradation of hydroxylapatite, fluorapatite, and fluorhydroxyapatite coatings of dental implants in dogs. J Biomed Mater Res 1999; 48: 224–234. 6. Overgaard S, Lind M, Josephsen K, et al. Resorption of hydroxyapatite and fluorapatite ceramic coatings on weight-bearing implants: a quantitative and morphological study in dogs. J Biomed Mater Res 1998; 39: 141–152. 7. Liu DM, Chou HM and Wu JD. Plasma-sprayed hydroxyapatite coatings: effect of different calcium phosphate ceramics. J Mater Sci Mater Med 1994; 5: 147–153. 8. Sergo V, Sbaizero O and Clarke DR. Mechanical and chemical consequences of the residual stresses in plasma sprayed hydroxyapatite coatings. Biomaterials 1997; 18: 477–482. 9. Yang CW, Lee TM, Lui TS, et al. A comparison of the microstructural feature and bonding strength of plasmasprayed hydroxyapatite coatings with hydrothermal and vacuum post-heat treatment. Mater Trans 2005; 46: 709– 715. 10. Fathi MH and Mohammadi Zahrani E. Mechanical alloying synthesis and bioactivity evaluation of nanocrystalline fluoridated hydroxyapatite. J Cryst Growth 2009; 311: 1392–1403. 11. Ben Ayed F, Bouaziz J and Bouzouita K. Pressureless sintering of fluorapatite under oxygen atmosphere. J Eur Ceram Soc 2000; 20: 1069–1076. 12. Ben Ayed F, Bouaziz J and Bouzouita K. Calcination and sintering of fluorapatite under argon atmosphere. J Alloy Compd 2001; 322: 238–245.
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