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)
Proof-reading Corrections: No. Original Correction 1. Reference [39] 39. A. K. Lynn and D. L. DuQuesnay, Biomaterials 23, 1947 (2002). 2. Reference [97] 97. M. Wei, A. J. Ruys, B. K. Milthorpe, C. C. Sorrell, and J. H. Evans, J. Sol–Gel Sci. Techn. 21, 39 (2001). 3. Figure 6b ⊕ PO ∇ CO Titania 4 3
48 h ⊕ ∇ ⊕ ⊕ ∇∇ ⊕ ⊕
⊕ ⊕ ⊕ ∇ 24 h ∇ ⊕⊕ Reflectance /arb. units /arb. Reflectance 0 h
2000 1600 1200 800 400 Wavelength /cm-1
CHAPTER 13
Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials
Jin-Ming Wu1,Min Wang 2
1Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, P. R. China 2Department of Mechanical Engineering, The University of Hong Kong, Pokfulam Road, Hong Kong
CONTENTS 1. Introduction ...... 2 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 3. Nanostructured Titania for In Vitro Apatite Deposition .... 10 3.1. Coating Approaches ...... 11 3.2. Chemical Modification Approaches ...... 14 4. Composite Coatings ...... 20 4.1. Ceramic/Metal Composite Coatings ...... 20 4.2. Ceramic/Ceramic Composite Coatings ...... 21 4.3. Ceramic/Polymer Composite Coatings ...... 22 5. Bioceramic Coatings Incorporated with Biomolecules ..... 24 6. Concluding Remarks ...... 25 References ...... 25
ISBN: 1-58883-170-1 Metal Oxide Nanostructures and Their Applications Copyright © 2010 by American Scientific Publishers Edited by Ahmad Umar and Yoon-Bong Hahn All rights of reproduction in any form reserved. Volume 4: Pages 1–31 1 2 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 hydroxylap- atite nanocrystals at the tissue-implant interface, opened a new era of design, manufac- ture, 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 biomateri- als, 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]. There- fore, 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 intro- duce a bioactive surface on metallic biomaterials, with special attentions paid to the bone- bonding 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 Na2O–CaO–SiO2–P3O5 system glasses of certain com- positions 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 trical- cium phosphate (TCP), and bioactive glass-ceramics, such as Ceravital@ and A–W glass- ceramics, 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 load- bearing 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 pro- vides 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 bio- ceramic 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 rel- atively 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 Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials 3
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 bone- implant interfacial strength at 12 weeks after implantation and one year after loading, although the bone apposition ability improved significantly as compared to the plasma- sprayed 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 decom- posed 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 ceram- ics (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 coating- substrate 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 interme- diate 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 plasma- sprayed 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 rough- ness on bone/implant fixation revealed that plasma-sprayed fluorhydroxyapatite (FHA) coatings on a complex system inhibited bone mineralization, due to the loss of the coat- ing 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 spray- ing [12]. The coating induced apatite deposition in SBF after 1 d and was covered thor- oughly after 2 d soaking. Thus, the original bioactive property of glasses was preserved 4 Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials
during plasma spraying [12]. Plasma-sprayed wollastonite (CaSiO3 and dicalcium sili- cate ( -Ca2SiO4) 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 ion- ized, 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 sub- strates 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 physico- chemically 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 treat- ment, 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 sput- tering [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]. Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials 5
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 coat- ings 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 crys- tallization 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. Coat- ings with different crystalline structures, ranging from amorphous and mixed crystalline phases to pure crystalline HA, could be deposited under different conditions thus pro- viding 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 pro- vides 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 sub- strate by PLD, followed by a soft laser treatment [67]. The PLD-derived coating was com- posed 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 pro- cesses. 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 6 Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials
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 CTEof 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. How- ever, 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 ele- ments 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 NH4H2PO4, a composite coating of apatite and brushite (CaHPO4 · 2H2O, DCPD) was deposited on the cathodic titanium substrate [86, 87]. Under the application of a DC voltage, the following reaction occurred on the cathode, − − 2H2O + 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 3− − the cathodic Ti substrate to react with the PO4 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 treat- ment. Monetite (CaHPO4) was also deposited on the surface of the Ti cathode, using a solution of Ca(OH)2,H3PO4, 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 coat- ing to thermodynamically stable HA. OCP and carbonated HA (CHA) layers could also Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials 7 be deposited on metallic substrates using an electrolyte of 137.8 mM NaCl, 1.67 mM
K2HPO4, and 2.5 mM CaCl2 · 2H2O aqueous solution [92]. Both the crystals and their morphologies could be modified by controlling the current density, the electrolyte tem- perature, 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 concentra- tions 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 Cfor1hto 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 sub- strates 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 advan- tage 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 param- eters, 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 forma- tion 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 HCO3 , 8 Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials
2+ 2− 2+ Ca , HPO4 , and Mg ions, was developed for depositing the apatite coating [108]. 2+ 2− Compared to SBF, the concentrations of Ca (6.0 mM vs. 2.5 mM) and HPO4 (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 HCO3 ions from the solution in the form of CO2 at the tem- − − perature of 45 C, through the reaction of HCO3 → 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 NaH2PO4 · H2O + 2.9% H3PO4 at 75 C for 24 h, resulting in the deposition of a monetite (CaHPO4) coating. Second, the monetite crys- tals 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 + 2+ − − titanium [111]. The calcifying solution contained 25.5 Na , 2.5 Ca , 5.0 Cl , 18.0 HCO3 , 3− and 2.5 PO4 (in mM). A uniform HA coating was deposited on Ti-6Al-4V substrates by 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 supersat- urated 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 sur- face 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 topog- raphy 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 elec- trochemical 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 contain- ing 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. Bioactive Bioceramic Coatings: Part II. Coatings on Metallic Biomaterials 9
2.4.2. Chemical Modifications
Treating CPTi and Ti-6Al-4V substrates with an HCl and H2SO4 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 supersatu- rated 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 witha5MNaOH solution to form a sodium titanate layer also induced biomimetic apatite deposition in supersaturated Ca–P solutions [126]. A simi- lar 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 sur-
face 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 sat- urated Ca–P solution, the dissolution of calcium ions and adsorption of phosphate ions occurred. The adsorbed phosphate ions reacted with surface calcium compounds, lead- ing 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 sub- strate 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, –SO3H, –PO4H2, –CH3, and –NH2, etc. were introduced on a Ti-6Al-4V sub- strate 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,PO4H2, 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 PO4H2 > 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 ion- adsorption upon complexes with a surface negatively charged group. Poorly crystallized HA was deposited from a saturated Ca–P solution on the SAM
surfaces with –PO4H2 and –COOH functional groups, but not on the SAM surfaces with –CH CH2 and –OH groups [136]. The –PO4H2 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, usu- ally 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 possess- ing 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-Na2O-K2O-P2O5 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 specifi- cally 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 com- plexes 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 mor- phology, 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, densifica- tion, 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 nucle- ation 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% H2O2 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 H2SO4,H3PO4, 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 oxi- dation 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 H2O2 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
4+ − TiCl4 → Ti + 4Cl (2)
(b) Formation of the titanium peroxo complex
4+ 4−n + + Ti + H2O2 + n − 2 H2O → TiO2OH n−2 + nH (3)
(c) Generation of a basic environment at the cathode
− − 2H2O + 2e → H2 + 2OH (4)
(d) Hydrolysis of the peroxo complex to form the hydrated titania deposit
4−n + − TiO2OH n−2 + 4 − n OH + kH2O → TiO3H2O k+1 (5)
The electrodeposited hydrated titania transformed to anatase after hydrothermal treat- ing in steam at a temperature of 180 C:
2TiO3H2O x → 2TiO2 + O2 + 2xH2O (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 low- temperature 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 H2O2 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 13
(a) (b)