Metal Oxide Nanostructures and their Applications (ISBN 1-58883-170-1) Edited by Ahmad Umar, Najran University, Saudi Arabia Yoon-Bong Hahn, Chonbuk National University, South Korea. Published by American Scientific Publishers, CA, USA, (http://www.aspbs.com/mona/)

October 2009 5-Volume Set, 4,000 pages, Hardcover

Volume 4 Applications (Part-2) Chapter 13 Bioactive Bioceramic Coatings. Part II: Coatings on Metallic Biomaterials (Vol.4, pp 479-509)

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Reference [97]

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Figure 6b

Correction 39. A. K. Lynn and D. L. DuQuesnay, Biomaterials 23, 1947 (2002). 97. M. Wei, A. J. Ruys, B. K. Milthorpe, C. C. Sorrell, and J. H. Evans, J. Sol–Gel Sci. Techn. 21, 39 (2001). ⊕ PO4

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CHAPTER 13

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials Jin-Ming Wu1 , Min Wang2 1

Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, P. R. China 2 Department of Mechanical Engineering, The University of Hong Kong, Pokfulam Road, Hong Kong CONTENTS 1. 2.

3.

4.

5. 6.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Bioactive Bioceramic Coatings . . . . . . . . . . . . . . . . . . . . . . . 2 2.1. Physical Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 2.2. Sol–Gel Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5 2.3. Electrocrystallization, Electrophoretic Deposition, and Electrostatic Spray Deposition . . . . . . . . . . . . . . . . 6 2.4. Biomimetic Process . . . . . . . . . . . . . . . . . . . . . . . . . . . 7 2.5. Other Coating Techniques . . . . . . . . . . . . . . . . . . . . . 10 Nanostructured Titania for In Vitro Apatite Deposition . . . . 10 3.1. Coating Approaches . . . . . . . . . . . . . . . . . . . . . . . . . 11 3.2. Chemical Modification Approaches . . . . . . . . . . . . . . 14 Composite Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 20 4.1. Ceramic/Metal Composite Coatings . . . . . . . . . . . . . 20 4.2. Ceramic/Ceramic Composite Coatings . . . . . . . . . . . 21 4.3. Ceramic/Polymer Composite Coatings . . . . . . . . . . . 22 Bioceramic Coatings Incorporated with Biomolecules . . . . . 24 Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25 References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

ISBN: 1-58883-170-1 Copyright © 2010 by American Scientific Publishers All rights of reproduction in any form reserved.

Metal Oxide Nanostructures and Their Applications Edited by Ahmad Umar and Yoon-Bong Hahn Volume 4: Pages 1–31

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Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

1. INTRODUCTION The concept of bioactive bonding, which is defined as the direct chemical bonding of biomaterials to the surrounding tissue through the formation of fine hydroxylapatite nanocrystals at the tissue-implant interface, opened a new era of design, manufacture, and application of biomaterials, when Hench et al. published their epoch-making paper in 1971 [1]. Although various bioactive glasses, ceramics, and glass-ceramics have been developed since the 1970s, most of these so-called bioactive bioceramics are not suitable for the load-bearing conditions for the replacement of bone or teeth, due to their poor mechanical properties. Compared to bioceramics and biomedical polymers, metals used in the medical field possess high mechanical properties, long fatigue life, good wear resistance, and excellent formability. Some of these metals exhibit additional properties of good corrosion resistance and acceptable biocompatibility and hence are good candidates for implants for high load-bearing applications. Metallic biomaterials, such as commercially available pure Ti (CPTi), Ta, Nb, 316L stainless steel, Co–Cr, Co–Cr–Mo, and Co–Ni–Cr alloys, have been widely used in restorative surgery [2, 3]. However, none of the metallic biomaterials is bioactive. When they are implanted in the body, a fibrous capsule is formed to separate them from the host tissue [4]. Therefore, numerous efforts have been made to introduce a bioactive surface on the metallic biomaterials so that implants of these surface-modified metals can be bounded to the surrounding tissue shortly after implantation. In the first chapter of this two-chapter series on bioactive bioceramic coatings, we reviewed the concept of bioactivity, in vivo and in vitro evaluation methods for bioactivity, mechanisms of apatite formation on biomaterials, and bioactive bioceramic coatings on non-metallic biomaterials. This chapter concentrates on various approaches used to introduce a bioactive surface on metallic biomaterials, with special attentions paid to the bonebonding ability of nanostructured titania thin films. For most systems, another major effect of coating metals with a bioceramic is the improvement in corrosion resistance and wear resistance of metal implants [5–7].

2. BIOACTIVE BIOCERAMIC COATINGS Hench et al. in 1970 found that the Na2 O–CaO–SiO2 –P3 O5 system glasses of certain compositions formed a direct chemical bonding to the surrounded bone tissue without being separated by a fibrous tissue and hence started the development of bioactive glasses [8]. In the following decades, bioactive ceramics, such as hydroxylapatite (HA) and tricalcium phosphate (TCP), and bioactive glass-ceramics, such as Ceravital@ and A–W glassceramics, were invented [9]. Perhaps with the exception of A–W glass-ceramics, bioactive glasses, ceramics or glass-ceramics themselves are not strong enough to be used for loadbearing applications. However, these materials can be deposited on the surface of strong, stiff, and tough metallic substrates in order for the metals to have bioactivity, which provides a solution to the clinical needs for strong and bioactive implants. The mid-1980s witnessed rapid industrial development of coating orthopedic and dental prostheses with HA using the plasma spraying technique, which has become an industrial gold standard. Since then, various techniques have been investigated to coat bioactive ceramics, mostly HA, on metallic substrates.

2.1. Physical Deposition 2.1.1. Thermal Spraying Plasma spraying [10–24] is the most widely used coating technique to form bioactive bioceramic coatings on metallic biomaterials. A typical plasma spraying operation employs plasma, or ionized gas, to partially melt and carry the ceramic particulates onto the surface of the substrate. During the coating process, the substrate is maintained at a relatively low temperature (generally lower than 300
C) so that the mechanical properties of the metallic implants will not be compromised. HA coatings with a typical thickness

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of 40∼60
m can be obtained. HA feed stock powders are generally synthesized using different chemical reactions [25, 26]. However, bovine- or human bone-derived HA has also been used [22, 23]. Using flame-spheroidized HA powders as feedstock was found to decrease coating porosity and increase deposition efficiency [18]. In vivo evaluation confirmed that plasma-sprayed HA coatings led to an earlier and better physiological integration of the implant with the bone [21, 27, 28]. However, it was reported that plasma-sprayed HA coatings on titanium did not contribute to the boneimplant interfacial strength at 12 weeks after implantation and one year after loading, although the bone apposition ability improved significantly as compared to the plasmasprayed titanium coating and uncoated titanium implants [29]. The temperature generated in the plasma exceeds 10,000
C, which makes it difficult to control the composition, crystal structure, and crystallinity of the resultant coating. HA coatings produced by the plasma spraying often contain considerable amounts of amorphous calcium phosphate (ACP) and small amounts of crystalline phases other than HA, such as - and -TCP [30] and tetracalcium phosphate (TTCP) [31]. At the coating/titanium interface deposited apatite may react with titanium dioxide that resulted from oxidation of titanium with plasma gas, or it may be thermally decomposed due to catalysis of titanium [32]. This adds to the complexity of controlling the plasma-spraying of HA coatings. The interfacial strength between the HA coating and the substrate is also insufficient due to the existence of micro-cracks as a result of the high residual stress generated due to the large difference between the linear thermal expansion coefficients (CTE) of ceramics (14.5 × 10−6 K−1 for HA [33]) and the metallic substrates (10.1 × 10−6 K−1 for CPTi [33]), especially for thick HA coatings [34–36]. A study on the HA coating on Ti-6Al-4V substrates revealed clearly a decreased bonding strength with increasing residual stress [35, 36]. In bonding tests, the fracture occurred mainly inside the HA coating under a low residual stress. Under a high stress, the fracture tended to occur along the coatingsubstrate interface. A plasma-sprayed HA coating with a thickness of 200
m exhibited higher residual stress than that having a 50
m thickness [35]. Introducing an intermediate layer as a bond coat was suggested to effectively improve the adhesion strength between the HA coatings and the substrates [22, 23, 37, 38]. However, careful selection of the intermediate layer is important. It should match with both the top coating and the substrate. Otherwise, an adverse effect will occur. The intermediate layers of plasmasprayed titanium and zirconia were found to reduce the interfacial strength between the plasma-sprayed HA and CPTi substrates, due to the reduction in the contact area between the intermediate layers and the HA coatings [30]. A post-spray heat treatment is generally applied to transform soluble phases in the plasma-sprayed calcium phosphate (Ca–P) coating to stable crystalline HA in order to improve the long-term stability of the coating. However, the post-spray heat treatment could lead to reduced fatigue life [39]. Lower post-spray heat-treatment temperatures may result in an integrated coating without compromising the mechanical properties [39, 40]. A study focusing on the concomitant influence of implant surface chemistry and roughness on bone/implant fixation revealed that plasma-sprayed fluorhydroxyapatite (FHA) coatings on a complex system inhibited bone mineralization, due to the loss of the coating adhesion to the substrate and higher dissolution rates, as well as ion release from the underlying metal [41]. In another study, the plasma-sprayed FHA coating was found to induce no apatite deposition in simulated body fluid (SBF) for 30 d [42]. A lower surface roughness could induce a good osteointegration process while a higher roughness did not induce the same osteointegration level and stimulated a conspicuous fibrous tissue formation at the interface, due to either an excessive irregularity of the surface or to an increased ion release [41]. In such cases, the presence of the plasma-sprayed HA coating could be less advantageous than no coating at all. Therefore, the biologic effects due to surface topography and chemistry must be simultaneously considered [41]. Bioactive glass coatings were deposited on Ti-6Al-4V substrates through plasma spraying [12]. The coating induced apatite deposition in SBF after 1 d and was covered thoroughly after 2 d soaking. Thus, the original bioactive property of glasses was preserved

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Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

during plasma spraying [12]. Plasma-sprayed wollastonite (CaSiO3  and dicalcium silicate (-Ca2 SiO4 ) coatings on Ti-6Al-4V substrates were also found to stimulate apatite formation in vitro [16, 17] and enhance the short-term osteointegration property in vivo, due to the surface Si-rich layer formed while immersed in physiological fluids [24]. The in vitro bioactivity decreased with the increasing crystallinity of wollastonite due to its decreased solubility in SBF [43]. Zirconia coatings were also deposited on titanium and CoCrMo alloy substrates using plasma spraying. The adhesive strength of the ZrO2 (4% CeO2  coating to the substrates was higher than 68 MPa [44]. Apart from plasma spraying, other thermal spraying techniques have also been used for producing bioceramic coatings. In flame spraying, the carrier gasses are not ionized, and the temperatures generated are much lower than in plasma spraying. The high velocity oxy-fuel (HVOF) technique involves a much higher velocity of ceramic particles and a much lower temperature than plasma spraying, thus maintained the initially high crystallinity of the HA particles used [11, 45]. Also, the very high kinetic energies led to high coating quality characteristic of high adhesion to the substrate, high cohesion of the coating, and less porosity [11]. The HVOF-produced HA coating on Ti-6Al-4V substrates induced a much thicker apatite layer than the conventional air plasma-sprayed HA coatings after 7 d soaking in SBF [45].

2.1.2. Sputtering Deposition Magnetron sputtering [46–48], radio frequency (RF) magnetron sputtering [49–53], and ion-beam sputtering [54–58] use an ion beam to bombard a target material in a vacuum chamber. The atomic-sized fragments of sputtered material form coatings on suitably placed substrates in the chamber. The main advantage of these methods is that physicochemically better defined Ca–P coatings could be produced [52]. A thermal treatment is needed to induce crystallization of the sputtered coatings after coating production. For RF-sputtering HA coatings, a temperature of 500
C was sufficient to provide thermal energy to achieve a predominantly crystallized HA coating [52]. The crystallization of the Ca–P coating is a hydroxyl-diffusion controlled process. Through the introduction of water vapor into the ion-beam sputtering process and post-deposition treatment, the crystallization temperature of the Ca–P coatings could be decreased to 400
C, which in turn avoided damages that were caused by the high temperature treatment, such as reduced adhesive strength and decreased purity of the HA phase [56]. The presence of water vapor at 450
C in the post-deposition heat treatment significantly improved the crystallinity of RF magnetron sputtered Ca–P coatings [51]. However, it did not have a significant effect on the crystallinity when the coatings were subjected to heating beyond 450
C. A study on osteoblast response to as-deposited and heat-treated Ca–P coatings made by RF sputtering showed higher specific alkaline phosphatase (ALP) activity of cells on as-deposited coatings that exhibited relatively poor crystallinity and lower contact angles [49]. The as-deposited coatings also exhibited higher ultimate interfacial strength than the heat-treated crystalline coating and the control of uncoated metal after implantation in the mandibles of dogs for 3 weeks [29]. At 12 weeks of implantation, no statistical differences in the mean ultimate interfacial strengths were observed among the three samples, which was similar to the result obtained on plasma-sprayed HA coating [29]. However, histomorphometric evaluation indicated a greater percent of bone contact for as-deposited Ca–P coated implants than for heat-treated Ca–P coated implants or control Ti implants 3 and 12 weeks after implantation [57]. The RF magnetron sputtered and heat-treated Ca–P coating as thin as 0.1
m was sufficient to simulate carbonate apatite deposition under in vivo conditions [52]. Thin bioceramic coatings, typically with a thickness less than 1
m, with controlled porosity and stoichiometry in the nanoscale region can be obtained by ion-beam sputtering [55]. The improved bone response to ion-beam sputtering deposited thin Ca–P coating was also confirmed in vivo using a rabbit animal model [58]. For an ion-beam sputtered HA coating on CPTi substrates, diffusion of Ca and P into the natural oxide film of CPTi was observed by the depth profile of Auger electron spectroscopy (AES) [59].

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This suggested that chemical bonding between HA and the substrate could be formed, which guaranteed an adequate interfacial strength.

2.1.3. Electron-Beam Evaporation The electron-beam deposition method, which uses an electron beam to evaporate the materials to be coated (evaporants), has been employed to produce a series of Ca–P coatings with various Ca/P ratios on CPTi and Ti-6Al-4V substrates [60–62]. The as-deposited coatings were amorphous and a thermal treatment of 500
C was required to induce crystallization of the coating. The stability of Ca–P coatings was improved after fluorine incorporation [62]. The ALP activity of the cells on the fluoridated HA (FHA) coating was not significantly different from that on the HA coating. However, the osteoblast-like cells proliferated on the FHA coating to a lower extent than on the HA coating. The HA coating on CPTi prepared through the electron-beam evaporation resulted in higher bone-to-implant contact as well as higher removal torque than the control sample [61].

2.1.4. Laser Deposition Laser ablation [63–65], or pulsed laser deposition (PLD) [66–70], allows better control of composition and crystallinity of coatings by varying the related parameters. Coatings with different crystalline structures, ranging from amorphous and mixed crystalline phases to pure crystalline HA, could be deposited under different conditions thus providing coatings with different solubilities in physiological solution [65]. The substrate temperature readily affected the crystallinity, the composition, and the Ca/P ratio of the HA coating [68]. Thin coatings of octacalcium phosphate (OCP), which is involved in the early biomineralization process and cannot be deposited using plasma spraying [71], were grown on Ti substrates heated at 20∼180
C, by PLD [69]. In addition, PLD provides strong bonding between coating and substrates and is able to deposit very thin coatings, to control surface roughness, to ablate any materials and to fabricate coatings on any substrates. A thin Ca–P layer with a highly disordered, amorphous-like structure, which resulted from the decomposition and melting of the initial HA or FHA materials in the laser plasma, could be deposited on CPTi when the deposition was carried out at a high laser beam fluence of about 12 J/cm2 [70]. Such a highly distorted Ca–P layer with a thickness of 2.7∼2.9
m gave no significant peaks in the X-ray diffraction (XRD) pattern and exhibited an extremely high hardness of about 18 GPa [70]. The failure load decreased with increasing film thickness due to the increased residual stress [64]. A bioactive pseudo-wollastonite (-CaSiO3  layer was deposited on a titanium substrate by PLD, followed by a soft laser treatment [67]. The PLD-derived coating was composed of pseudo-wollastonite and amorphous materials, which had a porous structure of gathered grains and poor cohesion. The subsequent soft laser treatment improved the crystallinity and cohesion, making the coating dense and well adhered to the substrate.

2.2. Sol–Gel Process Most of the above-mentioned physical deposition techniques are line-of-sight techniques. They are not applicable to implants with complicated shapes. Therefore, other coating techniques that can be used for implants with complex surface morphologies have been developed. HA has been deposited on CPTi [72–80] and stainless steel [81] through sol–gel processes. A typical route to prepare HA sol is to use triethyl phosphite and calcium nitrate, or phosphorus and calcium precursors, with water or anhydrous ethanol as the solvents [75, 82]. A subsequent thermal treatment at a temperature of 375∼500
C was used to improve the crystallinity of HA and also the adhesion strength [81]. For the water-based sol–gel process, a critical aging time was required to complete the reaction between the Ca and P molecular precursors to form a desired intermediate complex that permitted a further transformation to apatite under an appropriate thermal treatment [82]. A period greater than 24 h was also required to obtain monophasic HA for the anhydrous ethanol procedure [83]. The sol–gel approach provides a much milder condition for the synthesis

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of the HA coating and is able to coat implants with complicate shapes. However, the high reactivity of the precursors makes sol–gel transformation very fast and thus difficult to control. A thin HA layer was coated on a microarc-oxidized titanium substrate through the sol–gel method [76]. The microarc oxidation enhanced the biocompatibility of the Ti, and the bioactivity was further improved by the sol–gel HA layer. A prior coating with titania [77, 78] or calcium titanate [77] using the sol–gel technique has been proven to enhance the adhesion of sol–gel derived HA coatings to the Ti substrates and also the corrosion resistance of CPTi. To reduce the difference between CTEs of the coatings and the substrates and hence the residual stress, 8 wt% Mn was added to the pure Ti to develop a Ti–Mn alloy with a linear CTE of 13.1 × 10−6 K−1 , a value much closer to that of HA. A significantly higher bonding strength of sol–gel HA coatings with the Ti–Mn alloy than with the CPTi substrate has been obtained [33]. However, the addition of Mn decreased the ductility of the metal. The addition of citric acid into the dipping solution resulted in a strong gelation, and in turn an improved wetablility of the solution, which improved the adhesive strength of the sol–gel rough and porous HA coating to the substrate [80]. Addition of ammonium hydroxide into the sol also enhanced the phase and structural stability and morphological integrity of the sol–gel HA coating on Ti substrates because of the improved gelation that shortened the aging time prior to the heat treatment needed for crystallization of the apatite coating [79]. Sol–gel calcium titanate (CaTiO3 ) coatings on CPTi substrates were prepared using a precursor solution of calcium nitrate and Ti-isopropoxide dissolved in 2-methoxyethanol [84]. Crystallization of single-phase CaTiO3 coatings on titanium started at 500
C. However, the sol–gel CaTiO3 coating failed to induce apatite coverage after soaking in SBF for 12 weeks. Although an energy dispersive spectra (EDS) detected some Ca and P elements on the coated surface, no peaks corresponding to apatite could be found on the XRD spectrum [85].

2.3. Electrocrystallization, Electrophoretic Deposition, and Electrostatic Spray Deposition Using an electrolyte of Ca(NO3 2 and NH4 H2 PO4 , a composite coating of apatite and brushite (CaHPO4 · 2H2 O, DCPD) was deposited on the cathodic titanium substrate [86, 87]. Under the application of a DC voltage, the following reaction occurred on the cathode, 2H2 O + 2e− = H2 + 2OH−

(1)

This reaction increased the pH value in the vicinity of the cathode surface and hence increased the supersaturation corresponding to Ca–P. The positive Ca2+ ions migrated to − the cathodic Ti substrate to react with the PO3− 4 and OH ions there to form a Ca–P layer on the surface [86, 87]. The bath temperature, the voltage, and the current density affected the composition and crystal structure of Ca–P coatings [87]. At an ambient temperature of 25
C for coating formation, DCPD was the main component of the coating deposited at lower current densities. The HA structure was obtained at a current density above 10 mA/cm2 [87]. The H2 evolution on the cathode inhibited the crystallization process and damaged as well the homogeneity of the coating. The addition of ethanol in the electrolyte inhibited the gas evolution and hence improved the coating quality. The electrolytic deposition at 80 Torr improved bubble removal in the vicinity of the cathode surface and enhanced deposition of calcium phosphates [88]. A brushite coating was obtained on the Ti-6Al-4V substrate using a similar process [89]. The coating transformed to pure apatite during the subsequent hydrothermal treatment. Monetite (CaHPO4 ) was also deposited on the surface of the Ti cathode, using a solution of Ca(OH)2 , H3 PO4 , and lactic acid [90, 91]. A subsequent immersion of the coated sample in a 0.1 M NaOH solution at 60
C for 48 h converted the monetite coating to thermodynamically stable HA. OCP and carbonated HA (CHA) layers could also

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be deposited on metallic substrates using an electrolyte of 137.8 mM NaCl, 1.67 mM K2 HPO4 , and 2.5 mM CaCl2 · 2H2 O aqueous solution [92]. Both the crystals and their morphologies could be modified by controlling the current density, the electrolyte temperature, and the coating time. An electrochemical deposition of a flake-like OCP coating was applied to porous titanium, using an electrolyte with concentrations of calcium and phosphor ions 1.5 times higher than those in human blood plasma [93]. Pure nanophase HA was directly deposited on the cathodes using very low concentrations of calcium (0.61 mM) and phosphate (0.36 mM) ions and at a biological pH value (pH = 6.0) [94]. On the other hand, direct precipitation of HA on the anodes could also be achieved from a basic electrolyte (pH = 9.1) with high concentrations of calcium (0.78 M) and phosphate (0.24 M) ions, under an anodic voltage of DC 2∼4 V [95]. The electrostatic attraction of HA nuclei at the interface contributed to the direct deposition. A hydrothermal-electrochemical deposition method was used to deposit needle-like HA with various morphologies on CPTi [96]. In this approach, the electrolyte was heated in an autoclave assembled with two electrodes. Both the size and the shape of the HA needle could be regulated accurately by systematic control of the electrolyte temperature, current density, and current loading time, which in turn modified the nucleation and crystal growth of HA. Using HA particulate suspension in an alcohol or other suitable solution and then subjecting the suspension to an electric field, HA could be deposited on Ti [7, 97–100], Ti-6Al-4V [97] and 316L stainless steel [97, 101–103] substrates through electrophoretic deposition (ED). A subsequent high temperature sintering was often required to achieve a strong bonding between the coating and the substrates. HA coatings were developed on 316L stainless steel through ED at an optimized potential of 60 V for 3 min from an HA powder suspension in isopropyl alcohol, followed by vacuum sintering at 800
C for 1 h to enhance the bonding strength [101–103]. Prior to the ED of HA on the CPTi substrates, an intermediate sol–gel dip-coating layer of silica or calcium–silica improved the adhesion strength of the HA coating to the metal substrate [7]. A prior alkali-treatment of CPTi, which formed a porous sodium titanate intermediate layer, contributed to a much denser and uniform apatite coating on titanium substrates [99]. A highly ordered macroporous apatite coating was also deposited on titanium substrates by an ED procedure followed by a subsequent heating at 900
C to remove the organics from the coating [98]. Although the influence of the HA coating produced via ED on the long-term stability of the implants has not been reported, a short-term advantage of the HA-coated implants has been confirmed [100]. In the electrostatic spray deposition (ESD) approach, a spray of charged, micro-sized droplets is generated and then directed toward grounded and heated substrates under the function of an applied potential. The droplet is accomplished by means of electrostatic atomization of precursor solutions containing inorganic precursor salts. Ca–P coatings with various morphologies could be deposited by adjusting various deposition parameters, such as the nozzle-to-substrate distance, the precursor liquid flow rate, and the deposition temperature [104, 105].

2.4. Biomimetic Process Mimicking the inorganic mineral formation process in natural organisms, apatite can be deposited from calcium phosphate supersaturated solutions, after functional groups have been introduced on metallic material surfaces, at the human body temperature of 37
C. Apatite film with a thickness of about 1
m can be obtained by simply immersing CPTi in SBF at 37
C for 4 weeks [106]. Bone-like apatite layers with thickness of ca. 5
m were also deposited on polished Ti-6Al-4V and Ti-Al-2.5Fe surfaces after soaking in Hank’s balanced salt solution (HBSS) at 37
C for 2 weeks [107]. The reaction between titanium and SBF was supposed to create a large number of Ti–OH groups that were essential for apatite nucleation [106]. However, the soaking time to induce apatite formation on natural surfaces of these metals is considered to be too long. Thus, an aqueous solution containing all major inorganic components present in the body, mainly HCO− 3,

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2+ ions, was developed for depositing the apatite coating [108]. Ca2+ , HPO2− 4 , and Mg Compared to SBF, the concentrations of Ca2+ (6.0 mM vs. 2.5 mM) and HPO2− 4 (2.4 mM vs. 1.0 mM) in the new solution were much higher. An apatite coating nearly 2∼3
m thick was obtained through repeatedly soaking CPTi in the solution at 45
C for 3 d for 3 times. The removal of HCO− 3 ions from the solution in the form of CO2 at the tem− perature of 45
C, through the reaction of HCO− 3 → OH + CO2 , resulted in an increase of solution pH, and hence initiated the apatite deposition process [108]. A uniform Ca–P coating was also deposited on Ti-6Al-4V substrates using a 5 SBF solution within 24 h [109]. A two-step chemical deposition method was developed to deposit an adherent apatite coating on titanium substrate [110]. First, titanium substrates were immersed in an acidic Ca–P solution of 0.2 M CaCO3 + 0.1 M NaH2 PO4 · H2 O + 2.9% H3 PO4 at 75
C for 24 h, resulting in the deposition of a monetite (CaHPO4 ) coating. Second, the monetite crystals were transformed to apatite by hydrolysis in an 0.2 M NaOH solution at 75
C for 24 h. Formation of sodium titanate during the transformation of monetite to apatite also favored apatite deposition and adhesion. A saturated Ca–P solution was developed recently for the quick deposition of HA on titanium [111]. The calcifying solution contained 25.5 Na+ , 2.5 Ca2+ , 5.0 Cl− , 18.0 HCO− 3, (in mM). A uniform HA coating was deposited on Ti-6Al-4V substrates by and 2.5 PO3− 4 simply soaking in the solution at 37
C for only several hours. The HA coating exhibited a finer lamellar structure than that formed from SBF. It was demonstrated that biomimetic nano-apatite was capable of conducting bone formation and promoting direct bone apposition [108, 112]. Such an osseointegration effect was significant even at the early stage of the implantation. The apatite-coated group exhibited a 21-fold greater fixation strength than the control group after implantation in a rat for 1 week [112]. The above-mentioned biomimetic process utilized the natural titania film existing on Ti or Ti-6Al-4V to induce Ca–P deposition. More typically, a biomimetic procedure involves introducing surface functional films or groups through seeding, chemical modifications, electrochemical deposition, or self-assembling, followed by immersion in a saturated Ca–P solution (typically SBF, which was developed by Kokubo’s group, or a variant such as 1.5 SBF or 5 SBF).

2.4.1. Seeding A two-step biomimetic approach has been used to deposit Ca–P coatings on titanium and its alloys [71, 106, 113–119]. First, the metallic substrates were soaked in a supersaturated Ca–P solution (5 SBF) to deposit a thin amorphous carbonated Ca–P film. This film acted as a seed surface for the subsequent growth of crystalline biomimetic Ca–P coatings, either CHA or OCP, from a second supersaturated Ca–P solution. The surface roughness of the substrates did not affect the heterogeneous nucleation of Ca–P in 5 SBF. However, the further growth and mechanical attachment of the final amorphous carbonated Ca–P coating depended strongly on the surface, for which a rough topography was beneficial [117]. This biomimetic approach could also be applied to porous tantalum [113, 114] and porous Ti-6Al-4V [115] implants. In vivo evaluation revealed that the OCP coating had a stronger osteogenic potential than CHA [113] and that the CHA coating enhanced bone integration when compared with the uncoated implants [114]. Rat bone marrow cell culture experiment suggested that a crystalline OCP coating was more suitable for cell attachment than an amorphous carbonated apatite [118]. However, the two-step biomimetic-deposited CHA represented the best substrate for goat bone marrow cells attachment, followed by the biomimetic deposited OCP and hence an electrochemical deposited CHA, owing to a higher dissolution rate and the relative rougher surface [116]. Soaking titanium in a solution of sodium silicate buffered at pH 7.25 and containing 100 ppm Si could also induce the subsequent apatite formation in 1.5 SBF [120]. In vivo tests confirmed biocompatibility and bioactivity of the biomimetic-deposited apatite coating.

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2.4.2. Chemical Modifications Treating CPTi and Ti-6Al-4V substrates with an HCl and H2 SO4 mixed acidic solution followed by a boiling 0.2 M NaOH solution at 140
C for 5 h under the pressure of 3 bars, an apatite layer formed on the metal surface after soaking in Ca–P solutions [121–125]. A Ca–P coating with a thickness of 20
m was formed from the supersaturated Ca–P solution, which consisted of two layers: an outside loose OCP crystal layer and an inside dense CHA layer [122]. Coating Ti-6Al-4V substrates with porous Ti and then treating them with a 5 M NaOH solution to form a sodium titanate layer also induced biomimetic apatite deposition in supersaturated Ca–P solutions [126]. A similar biomimetic procedure was used to obtain a uniform OCP layer on the inner pores of porous titanium substrates [93]. The crystal structure and morphology of the OCP coating was similar to that obtained by electrochemical deposition from the same Ca–P solution. A CHA layer with a thickness of a few micrometers was deposited on a CPTi surface by simply soaking in a boiling saturated Ca(OH)2 solution for 40 min and then in a supersaturated Ca–P solution at 37
C for 3 d [127]. Fine particles with sizes of 1∼2
m, which were determined to be Ca(OH)2 , CaTiO3 , and CaCO3 , uniformly covered the CPTi surface after precalcification treatment. During subsequent soaking in the saturated Ca–P solution, the dissolution of calcium ions and adsorption of phosphate ions occurred. The adsorbed phosphate ions reacted with surface calcium compounds, leading to the formation of large nuclei of Ca–P, which grew spontaneously to form the CHA layer. Alkali and heat treatment of 316L stainless steel, i.e., soaking in a 10 M NaOH solution at 60
C for 24 h followed by heating at 600
C for 1 h, resulted in nothing on the metal surface [42, 128, 129]. However, some researchers detected a thin sodium chromium oxide layer on the surface [130]. After the alkali-treated 316L stainless steel was soaked in 1.5 SBF at a temperature of 80
C, a dense and uniform bone-like apatite layer was formed on the surface [130].

2.4.3. Electrochemical Deposition Under the illumination of ultraviolet light, the porous nanocrystal titania film consisted mainly of anatase, which was produced by micro-arc oxidation of a CPTi plate at voltages of 250∼400 V, could induce quick deposition of bone-like apatite in SBF within 2 h, due to the beneficial Ti3+ and ·OH radicals as a result of photo-generated electron–hole pairs [131]. Thin films of amorphous Ca–P, which also contained some OCP crystals, were obtained by electrochemical deposition in an electrolyte of supersaturated calcium and phosphate solution [132]. After a subsequent soaking in another supersaturated Ca–P solution for several hours, a continuous OCP layer formed on the surface.

2.4.4. Self-Assembled Monolayer (SAM) The self-assembly (SAM) technique refers to the spontaneous formation of an ordered monolayer of organic molecules on a surface. The molecules can assemble onto a substrate in a self-limiting manner so that only one monolayer of the molecules is deposited at each step. The molecules forming SAMs all contain a head group that bonds to the surface, the body of the molecule, and an end functional group [133]. Functional groups of –COOH, –SO3 H, –PO4 H2 , –CH3 , and –NH2 , etc. were introduced on a Ti-6Al-4V substrate using a self-assembly technique [134]. A crystallized Ca–P layer was deposited on all SAM surfaces after subjecting the Ti-6Al-4V substrate to an immersion in various supersaturated Ca–P solutions. Various functional groups possess different apatite-forming abilities [135]. SAMs of alkanethiols having CH3 , PO4 H2 , COOH, CONH2 , OH, and NH2 terminal groups were formed on a gold surface via sulfur attachment. In SBF, the growth rate of apatite decreased on the order of PO4 H2 > COOH ≥ CONH2 OH > NH2 ≥ CH3 0. The fact that negatively charged groups strongly induced apatite formation, but the positively

10

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

charged group did not, suggests that apatite formation was initiated via calcium ionadsorption upon complexes with a surface negatively charged group. Poorly crystallized HA was deposited from a saturated Ca–P solution on the SAM surfaces with –PO4 H2 and –COOH functional groups, but not on the SAM surfaces with –CH CH2 and –OH groups [136]. The –PO4 H2 group exhibited a stronger nucleating ability than that of –COOH. Regardless of its amount, the pre-deposited HA rapidly induced biomimetic apatite layer formation after immersion in 1.5 SBF for 18 h.

2.5. Other Coating Techniques In a so-called rapid immersion coating approach, the metallic substrate is oxidized, usually simply by heating it in air at appropriate temperatures, before being immersed in a molten bath of glass [137]. Through the interaction between the oxide layer and the molten glass, a chemical bonding can be achieved at the metal–glass interface. The outer layer of bioactive glass on the metallic substrate is not contaminated by metal oxides due to the short immersion time, which is only a few seconds, and hence the bioactivity of the coating is not damaged. The rapid immersion method has been applied to 316L stainless steel, Co–Cr alloy and Ti alloy. HA spheres were implanted into surfaces of CPTi [138] and titanium alloys possessing super plastic deformation ability [139, 140] by hot pressing. After implantation, the dissolution of HA particles was supposed to leave cavities at implant surfaces for the ingrowth of new bone and hence functioned as micro-anchors for implants. Ion implantation of Ti surfaces with amino groups induced higher concentration of calcium and phosphorus precipitation and more mineralization [141], as well as enhanced osteoblast-like cells attachment [142], when compared to Ti surfaces not treated with ion implantation. Surface ion implantation of calcium also improved the bone conductivity of titanium [143]. By mixing ethanol with a bioactive glass powder and then coating it on titanium and its alloys, a well-bonded bioactive glass coating was obtained after a subsequent heat treatment [144, 145]. For the SiO2 -CaO-MgO-Na2 O-K2 O-P2 O5 system, glasses with silica content higher than 55 wt% could be used to prepare crack-free coatings with good adhesion through such an enameling approach. Increasing the silica content in the glass increased the thermal expansion to a value close to that of titanium. But in vitro bioactivity decreased with the increasing silica content beyond 60 wt% [145]. Using a similar method, a bioactive glass-ceramics coating containing -Ca3 (PO4 2 crystals was deposited on the surface of titanium alloys [146]. Various coatings of HA, rutile, corundum, -TCP, and their combinations were deposited on CPTi substrates through simple sandblasting at ambient temperatures using corresponding powders [147–149]. The bonding strength of the sandblasted HA layer was much higher than that achieved with other room temperature coating techniques, such as dipping, electrophoretic deposition, and electrochemical deposition [148]. The titanium implant with the sandblasted HA coating showed strong bone response and much better osteointegration in vivo, compared with the uncoated titanium [149].

3. NANOSTRUCTURED TITANIA FOR IN VITRO APATITE DEPOSITION Without any special surface treatments, CPTi exhibited excellent in vivo biocompatibility [150] and could induce apatite deposition in vitro [106]. However, on the naturally formed oxide layer, the apatite-forming process for Ti or Ti-6Al-4V takes an extremely long time [106]. Titania film prepared simply by thermal oxidation was found to induce apatite formation [151, 152]. Unfortunately, the induction time was also too long. To accelerate the apatite-forming process, a titania layer with various nanostructures can be specifically introduced on the CPTi surface through various methods other than simply heating titanium in air.

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

11

3.1. Coating Approaches 3.1.1. Sol–Gel Coating Sol–gel coating is a widely used process to prepare various nanostructured oxide films. The process involves preparations of sols through hydrolyzation of metal-organic complexes and polymerization of the sols to form an amorphous oxide film during the subsequent coating processes of dip-coating or spin-coating. Generally, to crystallize the amorphous oxide film, a heat treatment follows. This simple process has the advantage of being able to prepare oxide films with precisely controlled film thickness, surface morphology, porosity, specific surface area, composition, crystallinity, etc. Due to also the wide spectrum of applications of titania films in photocatalyst, optical devices, solar cells, and gas sensors [153, 154], sol–gel preparation of titania films has been widely studied. sol–gel titania films have been successfully coated on titanium and its alloys to induce bioactivity [155–165]. The mechanical properties of the titania layer, such as the hardness and the Young’s modulus, were strongly correlated to the polycondensation, densification, and crystallization behaviors [166]. It was therefore possible to tailor their responses under stress by varying the heat treatment temperatures. The implantation of the sol–gel prepared titania coating on a Ti-6A1-4V core into the femurs of goats showed an accumulation of Ca–P within the titania coating 12 weeks postoperatively, which led to the connection of the titania coating to the bone [155]. Homogeneous titania coatings were deposited on CPTi, which had been subjected to different pre-treatments of machining, plasma cleaning, titanium nitride coating, and sodium hydroxide corroded etching, all providing bonding strength sufficiently high for implants [155, 157]. Influence of the sol, its composition, the number of coating layers and surface morphology on the bioactivity was studied [156, 158, 160, 161]. Increasing the coating layers favored apatite deposition in SBF, due to the increased surface area [156]. The apatite-forming ability was the highest for the sol–gel coatings heated at 450∼550
C, or the coating heated at 600
C with the addition of valeric acid to sol [156]. It was found that only surface topographies giving peak distances of 15∼50 nm, as observed by atomic force microscopy (AFM), favored apatite deposition [158, 159, 161]. Doping the sol–gel titania coating with Ca and P ions decreased the apatite-forming ability, although the dissolution of Ca and P into SBF may have increased the supersaturation with respect to the apatite. This observation indicates the importance of surface roughness and Ti–OH functional groups for inducing apatite deposition in SBF: the Ca-dopant decreased the surface roughness, and the P-dopant reduced the Ti–OH groups [160]. In addition, as compared to the controlled CPTi, enhanced soft tissue attachment on the sol–gel titania coatings was identified, contributing to their ability to initiate calcium phosphate nucleation and growth on the surface [164]. It was also possible to adjust the bioactivity of the as-coated sol–gel titania coatings, which can be important for various surfaces of the same implant for both hard and soft tissue contacts, by surface treatment with a CO2 laser to induce crystallization of the amorphous as-coated sol–gel titania within selected areas [162, 163].

3.1.2. Slurry Coating Dipping in anatase or rutile gelatin slurry followed by heating at 750
C for 30 min to solidify the titania and to improve the adhesion strength provides another approach to coating titania on CPTi substrates [151]. After immersion in a supersaturated Ca–P solution at 37
C for 2 weeks, the titania layer formed by dipping in the anatase slurry showed greater apatite coverage on the surface than the titania formed by dipping in the rutile slurry. The apatite coverage on both titania was higher than on the titania film with rutile as the predominant phase, which was derived by oxidizing CPTi with 30 mass% H2 O2 at 70
C for 72 h followed by heating at 750
C for 30 min [151].

3.1.3. Anodic Oxidation Titania films produced by anodic oxidation of CPTi, which uses spark discharges at a high electrolytic voltage to form rough porous titania surfaces, also exhibited in vitro

12

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

bioactivity [76, 167–175]. The electrolytes used could be H2 SO4 , H3 PO4 , acetic acid, etc. A subsequent thermal treatment at about 600
C was required to crystallize the titania film. The apatite induction time in SBF decreased with increasing anatase or rutile [169, 170]. In vivo examination revealed that the porous titania film produced by the anodic oxidation method possessed high bone-bonding ability at the early stage of implantation (4 and 8 weeks). However, at the later stages of implantation (16 and 24 weeks), the improvement of the bone-bonding ability was not as significant due probably to the low porosity and to the superficial ingrowths of apatite-like deposits into the pores of the titania layer [169]. The bone-bonding ability of the anodic oxidized CPTi was higher than that of alkali- and heat-treated CPTi and comparable to that of sodium-free alkali- and heat-treated CPTi [169]. In vivo tests revealed significantly stronger bone anchorage and higher removal torque values for the anodic oxidized implants with oxide thickness of 600, 800, or 1000 nm than the implants with an oxide thickness of 17 or 200 nm. The different bone-bonding ability may be ascribed to oxide thickness, porosity, pore size distribution, and crystallinity of the oxides [173].

3.1.4. Electrodeposition A layer of hydrated titania can be deposited on an NiTi substrate serving as a cathode in an electrolyte consisting of TiCl4 in a mixture of water, methyl alcohol, and H2 O2 at a temperature of 2
C [176]. The electrodeposition procedure is supposed to take place according to the following reactions [176–179]: (a) Dissociation of TiCl4 TiCl4 → Ti4+ + 4Cl−

(2)

(b) Formation of the titanium peroxo complex 4−n+

Ti4+ + H2 O2 + n − 2H2 O → TiO2 OHn−2

+ nH+

(3)

(c) Generation of a basic environment at the cathode 2H2 O + 2e− → H2 + 2OH−

(4)

(d) Hydrolysis of the peroxo complex to form the hydrated titania deposit 4−n+

TiO2 OHn−2

+ 4 − nOH− + kH2 O → TiO3 H2 Ok+1

(5)

The electrodeposited hydrated titania transformed to anatase after hydrothermal treating in steam at a temperature of 180
C: 2TiO3 H2 Ox → 2TiO2 + O2 + 2xH2 O

(6)

The crystallinity of such low-temperature derived anatase was comparable to that derived by heating the hydrated titania at a high temperature of 500
C. The lowtemperature crystallized anatase layer increased not only the bioactivity, but also the corrosion resistance of the NiTi alloy [176]. To carry out the electrodeposition at room temperature, TiCl4 can be replaced by TiOSO4 to react with H2 O2 to form the titanium peroxo complex, in the presence of nitrate ions [180]. The cathodic electrodeposition procedure has been applied to modify metallic biomaterials of CPTi [178], Ti-6Al-4V [179], NiTi [176], and stainless steel [179, 180].

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

(a)

13

(b)

Figure 1. Surface morphology of CPTi after soaking in the TiF4 solution at 60
C for (a) 1 h and (b) 24 h, followed by ultrasonic cleaning for 5 min. Reprinted with permission from [184], J. M. Wu et al., Surf. Coat. Techn. 201, 3181 (2006). © 2006, Elsvier.

3.1.5. Low-Temperature Chemical Precipitation To prepare titania films with a high specific surface area, an acidic titanium tetrafluoride (TiF4 ) solution was used to deposit titania film with an anatase crystal structure on various substrates [181–184]. It was reported that, after removing the incorporated fluorine atoms from the anatase film by heating in air at temperatures beyond 300
C, CPTi coated with such films also exhibited in vitro bioactivity [182]. Figure 1 shows the anatase film derived by soaking CPTi in the TiF4 solution at 60
C for 1 h and 24 h, respectively, followed by ultrasonic cleaning for 5 min to remove the loosely attached nanoparticles. The film thickness increased gradually with increasing soaking durations, which affected readily the apatite formation ability of the anatase film [184]; however, the surface morphology of the anatase film with various thickness was similar, which was homogeneous aggregates of fine particles with sizes of ca. 20 nm. This method was applicable to biomaterial substrates that could undergo heating beyond 300
C. As can be discerned from the XRD patterns shown in Figure 2, various metallic substrates of Ti, NiTi, Ta, and SUS 316L were coated with a layer of apatite after introducing the bioactive anatase intermediate layer [183]. Similarly, well-crystallized rutile thin films with various thickness can be deposited on CPTi substrates by hydrolysis of aqueous titanium tetrafluoride (TiCl4 ) solution at a low temperature of 60
C [185]. Unlike fluorine atoms, the chlorine atoms incorporated in the rutile film do not affect the apatite formation, which is not surprising because SBF itself contains lots of chlorine ions; therefore, such an approach might be applied to polymer substrates in cases in which satisfying the interfacial strength between the substrate and the rutile layer can be guaranteed, because no subsequent thermal treatment is required. A mixed anatase and rutile film was also deposited directly on the surface after soaking CPTi in a TiOSO4 /H2 O2 solution at 80
C for 24 h, which induced apatite deposition effectively while soaked in SBF [186]. A subsequent hot water treatment promoted the crystallization and hence the apatite-forming ability [186].

3.1.6. Other Methods Plasma immersion ion implantation (PIII) has been used to form titania films on CPTi surface [187, 188]. However, high voltage (−30 KV), high vacuum (2 × 10−6 mbar), relatively high temperature (265–550
C), and sophisticated equipment were used in PIII [187]. Krupa et al. obtained an amorphous Ca- and P-rich layer with a thickness of ca. 100 nm on the CPTi surface by ion implantation of Ca and P in sequence, both at a dose of 1017 ions/cm2 at a beam energy of 25 KeV and a vacuum of 10−6 Pa [189]. After immersion in SBF, Ca–P precipitated but did not form a continuous layer due to the amorphous structure of the film. Anatase thin films were deposited on 316L stainless steel through the plasma assisted chemical vapor deposition (PACVD) method [190]. The film significantly improved the pit corrosion resistance and also the general corrosion resistance of the substrate. Any

14

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

Anatase

Apatite

SUS 316L

Ta

NiTi

Ti 20

30

40

50

60

2-theta/degree Figure 2. The TF-XRD patterns of the Ti, NiTi, Ta, and SUS 316L substrates, treated with the TiF4 solution, heated at 300
C for 1 h, and subsequently soaked in SBF for 5 d. The diffractions without any symbols were from the substrates. Reprinted with permission from [183], J. M. Wu et al., J. Mater. Sci. Mater. Med. 14, 1017 (2003). © 2003, Springer.

damage or partial removal of the titania coating did not cause an increased galvanic corrosion of the substrate. Recently, Liu et al. developed a new method to provide CPTi with bioactivity [191]. A slurry prepared by mixing pulverized glass of 3CaO · 4B2 O3 · 3TiO2 with ethanol was first coated on the CPTi substrate and then heated at temperatures of 600–750
C for 30 min to ensure a strong bonding. A subsequent hot water soaking at 80
C for 7 d thoroughly dissolved the glass and resulted in a rod-like rutile layer. The rutile layer then induced apatite deposition with a 3 d soaking in SBF. The highly oriented nanoscale potassium titanate (K2 Ti6 O13 ) rod arrays on CPTi substrates prepared through a similar procedure also induced thorough apatite coverage after soaking in SBF for 3 d [192].

3.2. Chemical Modification Approaches 3.2.1. Alkali Treatment A series of papers published by Kokubo and Nakamura’s group revealed that a bioactive surface could be obtained by treating CPTi with a 5 M or 10 M NaOH aqueous solution at 60
C for 24 h. A porous amorphous sodium titanate layer with gradual composition change from the Ti substrate to the film was formed on the titanium surface, which induced apatite deposition with a mechanism described in the first chapter of this review series. A subsequent heat treatment at 600
C caused densification and partial crystallization of the film to form a small amount of crystalline rutile and Na2 Ti5 O11 phases. The induction time for apatite formation in SBF increased from 72 h for the as-treated samples to 168 h [193]; however, a much higher bone/implant interface bonding strength was achieved [128, 194]. The alkali-treated titanium without heat treatment had no bonebonding ability due to the unstable reactive surface layer. Only after a subsequent heat

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

15

treatment, could the alkali-treated CPTi be bounded to bone [195]. The alkali- and heattreated CPTi showed higher a bonding strength of the apatite layer to the substrates than those of the apatite layers formed on Bioglass 45S5-type glass, dense sintered HA, and glass-ceramic A–W [194]. The alkali- and heat-treatment methods have been applied successfully to introduce bioactivity on the surfaces of porous Ti [126, 196–200] and various titanium alloys, such as Ti-6Al-4V [128, 196, 201], Ti-6Al-2Nb-Ta [128, 196], and Ti-15Mo-5Zr-3Al [128, 196, 202]. The alkali-treatment procedure for Ti-6Al-4V was optimized by immersing in 5 M NaOH solution at 80
C for 3 d followed by heat treatment at 600
C for 1 h, which resulted in a strongly adhering film inducing apatite deposition in SBF within 3 d [201]. Ta [196, 197, 203, 204] and Zr [196, 205] could also be alkali-treated to form Ta–OH and Zr–OH groups on the surfaces, respectively, which in turn initiated apatite nucleation. The bioactivity of metallic biomaterials subjected to such alkali and heat treatment has been verified by in vivo animal evaluations [195, 199, 200, 206–214]. However, attempts to apply the alkali-treatment to SUS 316L stainless steel to induce bioactivity were unsuccessful [42, 128, 129]. Sodium removal from the sodium titanate layer produced by the alkali and heat treatments, through hot water soaking, resulted in anatase on the surface, and hence enhanced in vitro bioactivity [213, 215]. However, the interfacial bonding strength between bone and the implant was compromised due to the loss of the graded structure [213]. To avoid such a drawback, the alkali- and heat-treated CPTi was pre-calcified through soaking subsequently in 0.5 M Na2 HPO4 overnight and then in a saturated Ca(OH)2 solution for 5 h, instead of sodium removal, to accelerate the apatite-forming process [216]. The final SBF soaking test showed that in vitro bioactivity was enhanced without compromising the interfacial strength. The degree of supersaturation with respect to apatite increased after the pre-calcination process, which in turn accelerated apatite deposition [216]. It was noticed that a pre-treatment of CPTi before the alkali treatment was important [217]. Without pre-treatments, reproducibility of the sodium titanate hydrogel layer formation after alkali treatment and the subsequent apatite deposition was poor due to differences in the titanium surface structure that depended on the titanium processing history before the alkali treatment. CPTi subjected to the alkali treatment directly, without any pre-treatment, induced only inhomogeneous apatite deposition after soaking in SBF for up to 20 d, due to the varying thickness of the natural oxide film. A pre-treatment of HCl acid etching of CPTi under an inert atmosphere before the alkali treatment led to the formation of a uniform micro-roughened surface that provided improved conditions for apatite formation in SBF. The apatite coating on the acid-etched and alkali-treated CPTi was homogeneous after soaking in SBF for 10 d [217].

3.2.2. Hydrogen Peroxide Treatment The excellent biocompatibility of CPTi has attracted much attention since it was noticed over two decades ago. It was found that oxidative Ti-peroxo intermediates formed through the interaction of Ti and hydrogen peroxide may be important for the biocompatibility of implants [218, 219]. The exposure of titanium in a phosphate-buffered saline (PBS) solution or Hank’s solution resulted in a passive film with dual layers: the inner layer had a structure close to TiO2 whereas the outer layer consisted of hydroxylated compounds [143, 220]. The introduction of H2 O2 in the PBS solution broadened the hydroxylate-rich region, probably due to the formation of a Ti4+ -H2 O2 complex [220]. Based on such findings, CPTi was treated at 60
C with hydrogen peroxide solutions containing various metal chlorides [221]. The addition of SnCl2 and TaCl5 in the solution resulted in a titania gel favoring the in vitro apatite deposition on CPTi in SBF due to the Ti–OH groups attached on the surface. In vivo tests confirmed the enhanced bone-bonding ability of the CPTi subjected to the treatment of hydrogen peroxide solution containing TaCl5 [222]. Fast healing and tight bonding were also observed on the Ti-6Al-4V with a titanium fiber mesh on the surface, which was subjected to the same hydrogen peroxide treatment [223]. This approach combined the osteoinductive ability

16

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

of porous morphology and the osteoconductive ability of the H2 O2 -modified surface to achieve a tight fixation of the implant at an earlier stage of implantation. A subsequent investigation revealed that heating the titania gel at an appropriate temperature of 400
C resulted in an anatase layer, which shortened significantly the induction time for apatite formation in SBF [224]. The ratio of CPTi surface area to the solution volume as well as the solution volume itself affected the in vitro bioactivity [225]. Treating CPTi with 8.8 M H2 O2 containing 0.1 M HCl at 80
C for up to 1 h followed by a thermal treatment at temperatures of 400∼600
C, which developed a predominantly anatase layer on the surface, also effectively induced apatite deposition in SBF [226]. The apatite layer adhered strongly to the surface of the H2 O2 -treated CPTi [227]. Direct oxidation of Ti in pure H2 O2 solution resulted in an amorphous hydrated titania layer with porous structure on the surface, as illustrated in Figure 3 [228–231]. The specific surface area of such derived titania thin film with nanosized pores can be as high as 535 m2 /g [231]. DeRosa et al. studied in detail the oxidation kinetics of Ti in an H2 O2 solution by monitoring the resistance of a thin film of Ti thermally evaporated on glass substrates [230]. The Ti film with a thickness varying from 20 to 200 nm, with various grain sizes, was immersed in 10 wt% H2 O2 solution at 80
C for various durations. Two distinct growing stages of the amorphous hydrated titania gel were recorded. In Stage I, growth of the titania layer occurred at the titania/Ti interface, which is controlled by the out-diffusion of the Ti species from the unreacted Ti film to the titania layer. In Stage II, the mechanism controlling growth was diffusion of the H2 O2 molecules through the titania layer with growth occurring at the titania/Ti interface. As a result, Ti films with a thickness of 50 nm or less exhibited only Stage I oxidation while thicker films exhibited both Stages I and II. The oxidation rate in Stage I was greater than that in Stage II, and the grain size of Ti affected the oxidation rate in Stage I but not that in Stage II. The microstructure of the Ti surface readily affected the formation of the bioactive titania layer. Wen et al. subjected CPTi substrates to a surface mechanical attrition treatment (SMAT), which introduced a surface Ti layer with grain sizes of tens of nanometers as a result of the repeatedly plastic deformation on the surface due to the attrition by stainless steel balls. Compared to the untreated coarsen grain Ti, the H2 O2 -treating temperature could be decreased to room temperature because the Ti surface with nanosized grains stored excess energy contributing the reactivity. Raman spectra combined with SEM examinations confirmed the formation of a nanoporous anatase layer on the Ti substrate subjected to SMAT [232]. Soaking the CPTi after SMAT at 8.8 M H2 O2 /0.1 M HCl at 80
C for 15 min followed by heating at 400
C in air for 1 h transformed the surface fine grain Ti to mesoporous anatase, which induced apatite deposition in SBF within 1 d [233]. Many researchers have emphasized the importance of a crystalline structure of titania to induce apatite deposition, which may be ascribed to a much smaller crystal mismatch between titania and apatite compared to the amorphous surface structure and hence

(a)

(b)

Figure 3. Surface morphology of amorphous titania layer on Ti substrates after immersing Ti in 30 wt% H2 O2 solution at 80
C for (a) 2 h and (b) 8 h.

17

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

(a)

(b)

(c)

(d)

Figure 4. Apatite coating precipitated from SBF after 1 d soaking on the surface of grooved CPTi (a), (b) and titanium mesh (c), (d), both of which have been subjected to treatment of 30 mass% H2 O2 containing 3 mM TaCl5 at 80
C for 3 d.

favors the eptaxial growth of apatite [165, 170, 213, 224, 226, 234–236]. Unfortunately, nearly all the techniques to produce crystalline titania coatings require a thermal treatment for crystallization. The high temperature thermal treatment not only deteriorates the mechanical properties of the metal substrates, but also causes the loss of Ti–OH groups in the titania layer that contributes efficiently to the bioactivity. Wu et al. found that simply prolonging the soaking time of CPTi in the hydrogen peroxide containing TaCl5 at 80
C (a)

(b) ⊕ Ti ⊕

⊕

⊕

72 h

⊕ ⊕ ⊕

8h

⊕ Ti

∇

⊕

Intensity/arb. units

Intensity/arb. units

72 h

∇ Anatase

⊕ ⊕

⊕∇

∇

⊕ ∇

⊕

∇

⊕

⊕ ∇

8h

⊕

⊕

∇ ⊕

2h 2h 20

25

30

35

40

2-theta/degree

45

50

20

∇ 25

⊕ 30

35

∇ 40

45

50

2-theta/degree

Figure 5. XRD patterns of CPTi oxidized by the H2 O2 solution at 80
C for various durations (a), followed by hot water aging at 80
C for 72 h (b). Reprinted with permission from [229], J. M. Wu et al., Surf. Coat. Techn. 201, 755 (2006). © 2006, Elsvier.

18

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials (b) ⊕ Apatite ∇ Anatase ∆ Ti

∆

Intensity/arb. units

∇

⊕

⊕

∇

∇ 24 h

⊕

∇

∆

0h 20

25

30

2-theta /degree

∆

∆

35

∇

∆ ∇

⊕

48 h

∆

∆

48 h

⊕

⊕ PO4 ∇ CO3 Titania

∇∇

Reflectance/arb. units

(a)

⊕

∇ ⊕

⊕

⊕

⊕

⊕

⊕ ∇

24 h ∇

∆

⊕⊕

0h

∇ 40

2000

1600

1200

800

400

Wavelength/cm–1 Figure 6. (a) XRD and (b) FT-IR spectra of CPTi oxidized by the H2 O2 solution at 80
C for 8 h, followed by hot water aging at 80
C for 72 h, and SBF soaking for various durations. Reprinted with permission from [229], J. M. Wu et al., Surf. Coat. Techn. 201, 755 (2006). © 2006, Elsvier.

for up to 3 d could directly achieve a layer of titania of the mixture of anatase and rutile on the surface [228, 237–240]. The crystallization of titania at such a low temperature was suggested to proceed mainly in a dissolution precipitation mechanism, favored by the aqueous environment [228, 237, 239]. The crystalline structure of the low-temperaturederived titania was influenced by the anions in the solution [228, 237, 240, 241]. In vitro bioactivity tests confirmed the excellent ability of such low-temperature-modified CPTi as inducing apatite deposition in SBF within 12 h [237]. Figure 4 shows apatite deposited on the low-temperature H2 O2 -modified titanium mesh and grooved CPTi after soaking in SBF for 1 d. The low-temperature approach is very effective to provide excellent bioactivity on titanium with complex surfaces. Similarly, amorphous titania gel derived from oxidation of CPTi in pure H2 O2 solution crystallized to anatase after a subsequent hot-water treatment, as can be discerned clearly from the XRD patterns shown in Figure 5 [229]. Apatite deposition was significant after soaking CPTi with such anatase film in SBF for 24 and 48 h (Fig. 6) [229]. The hot-water treatment has also been successfully applied to induce crystallization of amorphous titania derived by anodic oxidation of CPTi. The in vitro bioactivity of the hot-water treated titania, as evaluated by SBF soaking, was significantly superior over that crystallized by ordinary thermal treatment in air [242].

(a)

(b)

Figure 7. Surface morphology of the nanostructured titania layer after Fenton oxidation of NiTi substrates at 60
C for 24 h (a), and then followed by a boiling water treatment for 12 h (b). Reprinted with permission from [245], C. L. Chu et al., Acta Biomater. 3, 795 (2007). © 2007, Elsvier.

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

(a)

19

(b)

Figure 8. Rod-like and flower-like titania thin films derived through Ti–H2 O2 reactions.

The bioactivity of sodium titanate gel derived by the alkali treatment of titanium and its alloys depended heavily on ion release from the gel and hence was degradable in the aqueous solution [234]. In addition, the released Na+ caused an increase in external alkalinity that could trigger an inflammatory response and lead to cell death. Through repetitive washing with water, sodium could be reduced to ca. 3.5 at.%, which was the lowest sodium content still able to induce apatite deposition in SBF [243]. Such a problem of high local alkalinity is not encountered in the H2 O2 modification method, because apatite deposition in this route does not rely on the ion-exchange process, and the possible Ti-peroxo complexes formed during the chemical reaction already exits after the implantation of titanium implants in the human body and to the body’s inflammatory response [218]. Titania thin film with a mixture of anatase and rutile is reported to deposit on NiTi substrates through simply soaking in a boiling 30 wt% H2 O2 aqueous solution for 2 h [244]. A modified H2 O2 approach has also been developed to achieve nanostructured titania film on NiTi alloy substrates [245]. In this approach, trace Fe2+ ions were added to the H2 O2 solution to cause a so-called Fenton’s reaction, which has been widely used to remove organic pollutants in wastewater. After soaking the NiTi alloy in the solution at 60
C for several hours, the alloy surface was oxidized by hydroxyl radicals (·OH), which were produced by decomposition of the H2 O2 catalyzed by the Fe2+ ions. Figure 7 illustrates the obtained nanostructured titania film on the NiTi alloy surface. Similar to the results achieved by Wu et al. for depositing crystallized titania film on Ti substrates by low-temperature oxidation of Ti by H2 O2 solution [237], the low-temperature derived dense titania layer was also a mixture of anatase and rutile. A subsequent treatment in boiling water for 12 h further improved the crystallinity. The XPS measurement revealed that such Fenton oxidation produced a nanostructured titania film with depleted Ni in the surface and a graded interface with the substrate, which inhibited out-diffusion of the

(a)

(b)

Figure 9. Titania nanotube arrays derived through anodic oxidation of CPTi in 5 wt% HF solution at 20 V for 30 min.

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Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

(a) (b)

Figure 10. Surface morphology of a titania nanotube array (a) and then after alkali treatment followed by immersing in 1.5 SBF for 14 days (b). Reprinted with permission from [251], X. F. Xiao et al., Mater. Chem. Phys. 106, 27 (2007). © 2007, Elsvier.

toxic Ni in body fluid. In addition, the titania film possessed good mechanical properties and had little effect on the shape memory behavior of the NiTi alloy. Well-crystallized titania thin films with various nanostructures of nanorods and nanoflowers, as illustrated in Figure 8, have been achieved based on such simple Ti–H2 O2 reactions [241, 246–248]. In addition, well-aligned titania nanotube arrays, as shown in Figure 9 as an example, have been developed recently for several applications in gas sensors, dye-sensitized solar cells, and photocatalysis [249]. Liu et al. have demonstrated the unique property of rod-like rutile to induce apatite-deposition in SBF [191, 192]. Also, Oh et al. argued that the unique nanostructure of nanotube arrays accelerated apatite formation on titania after an alkali treatment to induce a bioactive sodium titanate on top of the tube wall [250]. A comparative study reported that the alkali treatment in 5 M NaOH for 30 min applied to Ti or compact titania film induced no apatite deposition in 1.5 SBF for 14 d, while that applied to titania nanotube arrays was capable of depositing a significant apatite layer, as illustrated in Figure 10 [251]. A cell culture result by Srammer et al. revealed that the titania nanotube array interacted more efficiently with primary bovine aorta endothelial cells as compared to a flat Ti surface. The combined nano topography and nano cues caused an enhanced endothelial response in vitro [252]. In addition to the ability to promote osteoblast differentiation, the titania nanotube arrays were also utilized to load antibiotics, thus providing a local delivery of antibiotics off-implant at the site of implantation [253]. Therefore, further studies on the bone-bonding ability of titania with various nanostructures other than nanotube arrays will be interesting.

4. COMPOSITE COATINGS To obtain coatings with enhanced mechanical properties, interfacial strength, and/or bioactivity, composite coatings of ceramic/metal, ceramic/ceramic, and ceramic/polymer have been developed on metallic biomaterial substrates.

4.1. Ceramic/Metal Composite Coatings Functionally graded ceramic/metal composite coatings on metal substrates can effectively minimize the residual stress caused by the different thermal expansion coefficients of ceramic coatings and metallic substrates and hence result in an enhanced adhesion [254–262]. Composite coatings were deposited on metallic substrates layer by layer, with gradually changed compositions, through mainly spraying techniques. The outer layer of the composite coating was Ca–P-rich to ensure in vitro bioactivity [259]. Table 1 lists some functionally graded ceramic/metal composite coatings on metal substrates. There were significant increases in the interfacial strength of the composite coatings. For an RF plasma-sprayed HA/Ti composite coating on Ti substrates, with a total thickness of 175
m, a maximal 50 MPa adhesive strength could be achieved; whereas for the pure

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

21

Table 1. Some functionally graded bioactive ceramic/metal composite coatings on metal substrates. Substrate Ti and its alloys Ti Ti Ti-6Al-4V Ti-6Al-4V

Coating

Preparation method

Reference

HA/Ti Ca–P/Ti HA/Ti HA/Ti-6Al-4V HA/YSZ/Ti-6Al-4V

RF plasma spraying Ion-beam sputtering Hydrothermal-electrochemical technique Plasma spraying Plasma spraying

[254, 255] [256] [257] [258–261] [262]

HA coating with a film thickness of 50
m, the adhesive strength was only 23 MPa [254]. The bonding strength of the composite coating was found to increase with increasing Ti content [256]. The mechanical properties of the ceramic/metal composite coatings are much higher and stable than the single ceramic coating. The hardness, Young’s modulus, and bonding strength of the HA/YSZ/Ti-6Al-4V composite coating were found to be significantly higher than those of pure HA coatings even after immersion in the SBF solution for a long time, which deteriorated the mechanical properties of the coatings [262]. Plasmasprayed and heat-treated HA/Ti-6Al-4V composite coatings on the Ti-6Al-4V substrates were stable, only undergoing small changes in microstructures after soaking in SBF for 1∼10 weeks [261].

4.2. Ceramic/Ceramic Composite Coatings Through adding calcium and phosphate complexes, such as calcium -glycerophosphate (Ca–GP) and calcium acetate monohydrate (CA), into the electrolyte, HA crystals could be precipitated and incorporated into the anodic oxidized titania film, which consisted of anatase and minor rutile, after a subsequent hydrothermal treatment [171, 172, 263–270]. An optimized condition of the anodic oxidation was a concentration of the electrolyte of 0.02 M Ca–GP and 0.15 M CA, current density of 70 A/m2 , and final voltage of ca. 350 V. The anodic oxidized titania film exhibited porosity, intermediate roughness, and high crystallinity [271]. Higher bone-to-implant contact for the composite coating was observed than on the CPTi control [171]. A culture of the rat bone marrow stromal cells on the titania/HA composite coating with thin HA as the outermost layer revealed higher levels of early cell attachment due to the improved hydrophilic wetability than on the untreated CPTi [266]. Early osteoblast attachment was enhanced on the anodic oxidized and hydrothermal-treated surfaces [268, 269]. The removal torque strength was significantly higher for the anodic oxidized implant than for the untreated sample at 6 and 12 weeks after implantation in the cortical bone of rabbits [270]. However, no significant difference was observed for the percentage of bone contact on all implants. Although the HA needles precipitated in the titania film during the hydrothermal treatment resulted in rapid osteointegration [268, 269], it decreased the removal torque because the HA needles bonded only weakly to the surface [270]. It is noted that the same research group reported enhanced bone apposition ability but not improved interfacial strength for the plasma-sprayed HA coating [29] and RF magnetron sputtered Ca–P coating [57]. It seems that there may not be a strong correlation between the bone apposition ability and the interfacial strength. The other electrochemical approach to prepare titania/Ca–P composite coating is anodic plasma-chemical (APC) treatment, which relies on the dielectric breakdown of an insulating oxide film at the surface of a metal anode in contact with a suitable electrolyte and is accompanied by visible plasma-like sparking at the anode surface. A composite coating with fully or partially soluble amorphous Ca–P phases embedded in a matrix of amorphous or anatase titania was deposited on CPTi by the APC treatment [272]. A mixed anatase and rutile film containing CaTiO3 crystals deposited through a plasma electrolytic oxidation was also found to induce apatite deposition in SBF for 14 d [273]. The sol–gel method was also used to deposit Ca–P/TiO2 composite coatings on CPTi [274–277]. The sol mixture was prepared either by mixing a titania sol and a solution

22

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

of calcium nitrate/phosphoric acid esters [274, 275], or by mixing a titania sol with an HA anhydrous ethanol suspension [276, 277]. A porous composite coating consisting of a titania (anatase) network encapsulating a HA particulate phase was found to adhere well to the substrates [277]. The adhesion of the (HA+-TCP)/TiO2 composite coating was better than that of the HA+-TCP coating obtained through a similar route, due to the involvement of a chemical component in the binding [275]. The interfacial strength increased with decreasing coating thickness and increasing roughness of the substrates [275]. A study on the behavior of human MG63 osteoblast-like cells on the sol–gel HA/TiO2 composite coatings confirmed the bioactivity of the composite coating owing to the presence of hydroxyl groups detected on the surface that promoted the calcium and phosphate precipitation and improved the interactions with osteoblastic cells [276]. A much simpler method was developed to prepare Ca-containing titania film through hydrothermal modification of the titanium surface in calcium solutions [277]. Soaking CPTi in a saturated Ca(OH)2 solution at 121
C under a pressure of 2 atm for 2 h in an autoclave resulted in anatase films containing CaTiO3 . The formation of CaTiO3 was favored by the high pH value, and the high temperature and high pressure involved in the hydrothermal process, according to the followed reactions: TiO2 + 2H2 O = TiOH4 aq

(7)

Ti + 2H2 O = TiO2 + 2H2 ↑

(8)

Ca2+ + TiOH4 aq = CaTiO3 s + 2H+ + H2 O

(9)

An apatite layer covered the Ca-containing titania film after soaking in Hank’s solution at 37
C for 30 d. This hydrothermal modification of CPTi in water promoted apatite deposition in Hank’s solution due to the increment in the number of hydroxyl radials. However, the precipitation of apatite was inhibited by the hydrothermal treatment in CaCl2 solution due to the decrease in the surface layer thickness and the lack of calcium ions contained in the surface layer [277]. Boyd fabricated a titania/Ca–P hybrid coating on Ti6Al4V substrates by sputtering a titanium inner layer and HA outer layer, followed by calcination at 700
C. The hybrid coating was achieved through out-diffusion of titania (rutile) into the HA layer. The growth and proliferation of osteoblast-like cells were more readily when compared to the HA coatings or the rutile TiO2 surfaces [278, 279]. Biphasic calcium phosphate (BCP) bioceramics are a group of bone-substituting biomaterials that consist of an intimate mixture of HA and -TCP, of varying HA/-TCP ratios. They are recommended currently for use as an alternative or additive to autogeneous bone for orthopedic and dental applications [280]. A composite coating consisting of a superficial layer of BCP and a deeper layer of HA ceramic (a sandwich coating) exhibited the highest shear strength as compared to BCP or HA coating alone [281]. The in vivo evaluation in beagle dogs confirmed that the sandwich coating prepared through plasma spraying on sandblasted Ti-6Al-4V substrate was effective in promoting massive metaphyseal osseointegration, which ensured mechanical stability for early weight-bearing and should prevent long-term complications. Table 2 lists some ceramic/ceramic composite coatings produced on metallic substrates. These composite coatings also possess higher mechanical properties, improved interfacial strength, and in vitro bioactivity than the corresponding bioceramic coatings of single phases.

4.3. Ceramic/Polymer Composite Coatings A CaO-SiO2 -poly(dimethylsiloxane) (PDMS) organic–inorganic composite coating was developed for Ti-6Al-4V substrates using a sol–gel dip-coating method [293]. After soaking in SBF for 7 d, the ion exchange between calcium ions in the coating and protons in the solution led to an increase in silanol groups and hence a nanocrystalline apatite-like phase formed on the surface.

23

Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials Table 2. Some bioactive ceramic/ceramic composite coatings on metal substrates. Substrates

Coatings

Preparations

Coating characteristics

Ti and its alloys

HA/TiO2

Porous film, needle-like HA precipitations on anatase and minor rutile matrix

[164, 165, 255–262]

Ti-6Al-4V

HA/TiO2

Anodic oxidation with a subsequent hydrothermal treatment Hybrid of microarc oxidation and electrophoresis

[175]

Ti

HA/TiO2

Inner layer: porous anatase Intermediate layer: mixed anatase and crystalline HA Top layer: dense crystalline HA HA-contained anatase and minor rutile film

Ti

Ca/TiO2

Ti

Ca–P/TiO2

Ti

Ca–P/TiO2

Anodic plasmachemical treatment

Ti

CaTiO3 /TiO2

Ti-6Al-4V Ti-6Al-4V

TiO2 /Ca–P TiO2 /HA

Anodic plasmachemical treatment Plasma spraying HVOF spraying

Ti-6Al-4V

HA/glass

Plasma spraying

Ti Ti

Sol–gel Sol–gel

Ti Ti-6Al-4V

HA/TiO2 (HA+TCP)/TiO2 CaTiO3 /TiO2 BCP/HA

Hydrothermal Plasma spraying

Ti alloy

HA/Zirconia

Plasma spraying

Ti, Ti-6Al-4V

HA/glass

Cullet method

Ti alloy

-TCP/glassceramic Ti-Ca–C-O-(N), Ti-Zr-C-O-(N), Ti-Si-Zr-O-(N) Ti-Nb-C-(N)

Sintering

Metals

Micro-arc oxidation followed by hydrothermal treatment Micro-arc oxidation followed by hydrothermal treatment Micro-arc oxidation

Magnetron sputtering

References

[264, 265]

Ca-contained anatase and rutile film

[283]

Macroporous, anatase, or rutile film containing Ca–P Ca–P amorphous phases embedded in a matrix of amorphous or anatase titania Anatase and rutile film containing CaTiO3 crystals Functional graded layer Titania-reinforced HA film, with some reaction production of CaTiO3 Under-layer of HA and a surface layer of CaOP2 O5 glass-HA composite, with 2 or 4 wt% of glass Mixed HA and anatase Mixed anatase, rutile, HA, and TCP Ca-contained anatase A superficial layer of BCP and a deeper layer of HA CaO and ZrO2 reinforced HA composite layer Compositional graded HA/glass composite coating Glass-ceramic layer containing -TCP crystals Amorphous films

[284, 285] [272]

[273] [286] [287] [288]

[272, 273, 285] [275] [277] [280, 281] [289] [290] [291] [292]

Sol–gel SiO2 coatings on 316L stainless steel showed no bioactivity. The inclusion of Ca- and P-alcoxides in the sol composition did not result in bioactivity. Bioactive coatings were obtained from suspensions prepared by adding glass (CaO · SiO2 · P2 O5 ) particles to a composite organic–inorganic SiO2 sol, due to the dissolution of glass particles while soaking in HBSS solution [294].

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Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials

Cooperative organic/inorganic systems self-assembled into ordered structures could be used to coat mesoporous silica films on many substrates of metals or polymers by spin coating of a sol–gel solution. The coated silica film induced apatite deposition in SBF due to both the mesoporous structure and the presence of silanol groups [295].

5. BIOCERAMIC COATINGS INCORPORATED WITH BIOMOLECULES Co-deposition of biologically active molecules, such as osteogenic agents and growth factors, with apatite crystals onto metal implants is of great interest for bone substitution [296–299]. The combination of Ti-mesh with rat bone marrow (RBM) cells could generate bone formation, although very limited new bone was formed [299]. A thin crystalline Ca–P coating on the Ti-mesh had an additional positive effect on the bonegenerating properties of the Ti-mesh scaffold, due to the high affinity of Ca–P ceramics to bone-growth-stimulating proteins [299]. In addition, local delivery of antibiotics, which offers advantages over systemic therapy such as decreasing the systemic toxicities and side effects of parenteral antibiotics, yielding higher drug concentrations in the relevant tissues, and thereby improving efficacy and reducing the necessary duration of treatment, has been advocated in clinical practices to prevent infection of the prosthesis [300]. Incorporation of drugs in bioactive ceramics also has great potential for developing drug-carrier systems in orthopedics as degradation of the carrier in vivo should result in the gradual exposure and release of incorporated molecules [297, 298, 300–302]. The extremely high temperatures involved in the plasma-spraying technique to produce HA coatings render the incorporation of biologically active molecules impossible. The incorporation of drug molecules in bioceramic coatings can be realized only by utilizing the mild solution approach, such as biomimetic process or electrochemical deposition. Plasma polymerization of alkyl amine was used to provide functional groups for immobilization of biomolecules on bioinert metal surfaces, without destroying their activities [303]. The amount of protein weakly and strongly bound to metallic biomaterials could be controlled by the choice of surface treatment and immobilization chemistry. Vancomycin-containing Ca–P coatings on Ti-6Al-4V rods were manufactured [207]. A Ca–P coating was first deposited on the metal substrates through electrophoretic deposition followed by sintering. The coating was then immersed in a simulated physiological solution containing vancomycin and HCl at 37
C, or followed by a second-stage lipid loading. The loading by immersion provided effective release of vancomycin and bacterial inhibition for up to 24 h. The subsequent lipid coating slowed antibiotic elution and hence demonstrated significant release and effective bacterial inhibition up to 72 h. Immersing chemically modified CPTi in a supersaturated calcification solution with the addition of bovine serum albumin (BSA), a bioactive protein, obtained BSA-loaded Ca–P coatings [297]. The Ca–P coating without BSA was a mixture of HA and OCP, while only the HA phase was detected in the BSA-loaded Ca–P coating. A continuous release of BSA was observed while soaking the Ca–P/BSA coating in an acidic phosphate-buffered saline at 37
C. The Ca–P/BSA coating was also produced on Ti-6Al-4V plates [298, 304]. A thin Ca–P layer was first deposited on Ti-6Al-4V by soaking in a 5 SBF solution at 37
C for 24 h, with the help of bubbled gaseous carbon dioxide. The subsequent soaking in a tris-buffered supersaturated Ca–P solution containing BSA at 37
C for 48 h resulted in co-deposition of BSA and Ca–P. The amount of BSA incorporated affected the crystal structure of Ca–P in a way that was different from the one on CPTi [297, 298]. In the presence of BSA, both OCP and carbonate-HA were deposited while only pure OCP was deposited in the absence of BSA. The co-deposited coating lost