UNDERSTANDING WEAR BEHAVIOR OF CoCrMo COATINGS ON Ti6Al4V AND TANTALUM COATING ON TITANIUM FOR LOAD-BEARING IMPLANTS STANLEY ALAN DITTRICK

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1 UNDERSTANDING WEAR BEHAVIOR OF CoCrMo COATINGS ON Ti6Al4V AND TANTALUM COATING ON TITANIUM FOR LOAD-BEARING IMPLANTS By STANLEY ALAN DITTRICK A thesis submitted in partial fulfillment of the requirements for the MASTER OF SCIENCE IN MATERIALS SCIENCE AND ENGINEERING WASHINGTON STATE UNIVERSITY SCHOOL OF MECHANICAL AND MATERIALS ENGINEERING MAY 2011

2 ii SIGNATURE PAGE To the Faculty of Washington State University: The members of the Committee appointed to examine the thesis of Stanley Alan Dittrick find it satisfactory and recommend that it be accepted. Amit Bandyopadhyay, Ph.D., Chair Susmita Bose, Ph.D. Neal Davies, Ph.D.

3 iii ACKNOWLEDGEMENTS I would like to thank Professor Amit Bandyopadhyay for all the support, guidance and encouragement he has given me during my time at Washington State University. I would also like to thank Dr. Vamsi K. Balla, Professor Susmita Bose and Professor Neal M. Davies for guidance along the way. I would also like to thank all of the group members, past and present, for their help and support. I would like to acknowledge the financial support from the W. M. Keck Foundation to establish a Biomedical Materials Research Lab at WSU. I also acknowledge the financial support from the M. J. Murdock Charitable Trust, the National Science Foundation (Grant No. CMMI ) and the National Institutes of Health (Grant No. NIH R01-EB ).

4 iv UNDERSTANDING WEAR BEHAVIOR OF CoCrMo COATINGS ON Ti6Al4V AND TANTALUM COATING ON TITANIUM FOR LOAD-BEARING IMPLANTS By Stanley Alan Dittrick, MSE Washington State University May 2011 Chair: Amit Bandyopadhyay In this investigation negative biological effects of cobalt chrome alloy (CoCrMo) were reduced or eliminated while maintaining a low wear rate. This was accomplished in two ways first by creating a gradient CoCrMo-Ti6Al4V structure and secondly by replacing 100% cobalt chrome with a tantalum coated titanium. Both types of samples were made using a laser engineered net shaping (LENS TM ) process. Novel, unitized structures, to minimize multiple parts for load-bearing implants, were fabricated with porous Ti6Al4V alloy on one side to encourage osseointegration and compositionally graded CoCrMo alloy on the surface, to reduce wear. Gradient structures with 50%, 70% and 86% CoCrMo alloy on the top surface showed high hardness between 615 and 957 HV. These gradient structures were evaluated for their in vitro wear rate and Co ion release in a simulated body fluid (SBF). It was determined the wear rate of ultrahigh molecular weight polyethylene and 100% CoCrMo alloy substrates depends on the hardness and microstructural features of the counter surface. In general, the wear rates of both substrates increased with a decrease in CoCrMo alloy concentration of gradient pins. However, high hardness gradient pins had a lower wear rate than 100% CoCrMo alloy pins. Gradient pins having 86% CoCrMo alloy had the lowest wear rates of 5.07 x 10-8 to 7.99 x 10-8 mm 3 /Nm. During in vitro wear testing the

5 v amount of Co released from gradient structures was, in the range of 0.38 and 0.91 ppm, whereas from 100% CoCrMo alloy it was 0.25 and 0.77 ppm. The second part of the study was focused on in vitro tribological performance of Ta coatings on Ti for load-bearing implant applications. Linear reciprocating wear tests in simulated body fluid showed a wear rate one order of magnitude lower, of the order of 10-4 mm 3 (N.m) -1, for Ta coatings compared to control Ti plate. This result demonstrated that laser processed Ta coatings can minimize early-stage wear debris. In summary, results from both studies indicated that intelligent coatings can be designed and fabricated using agile manufacturing for loadbearing implants to enhance bioactivity and reduce wear using unitized structures.

6 vi LIST OF TABLES Table 1 Co ion concentration (ppm) in the SBF after wear testing (*below the detection limit of the instrument) Table 2 Grain size (µm), surface roughness (µm) and hardness (HV 0.2) of Ta, CoCrMo and CP Ti. Laser energy density (J mm -2 ) used to fabricate Ta and CoCrMo alloy is also included

7 vii LIST OF FIGURES Figure 1 Influence of CoCrMo alloy concentration, in the top surface of metallic pins, on the wear rate of UHMWPE Figure 2 Top surface microstructures of the pins with varying concentrations of CoCrMo alloy in the top layer: (a) GP50, (b) GP86 and (c) P Figure 3 Typical morphologies of wear tracks on UHMWPE when tested against (a) GP50 pin, 1000 m (b) GP50 pin, 3000 m and (c) GP86 pin, 3000 m Figure 4 Wear rate of laser processed P100 and CoCrMo-Ti6Al4V alloy gradient pins Figure 5 Wear rate of laser processed S100 substrate tested for 3000 m against various gradient pins Figure 6 Typical morphologies of wear tracks observed during metal-on-metal wear testing (a) on GP50 pin, 1000 m (b) on GP86 pin, 1000 m and (c) on S100 substrate tested with GP Figure 7 Typical microstructures of (a) laser processed Ta coatings, (b) laser processed CoCrMo and (c) as-received CP Ti substrate Figure 8 Wear rate of CP Ti, laser processed Ta and CoCrMo alloy against hardened 100Cr6 steel ball Figure 11 Continued High magnification SEM images of worn areas for (a) CP Ti substrate (b) laser processed Ta and (c) laser processed CoCrMo alloy Figure A1 Tantalum pentoxide coated surface... 3 Figure A2 Potential nanotubes on tantalum pentoxide surface..3 Figure A3 Top view of nanotube covered surface..4 Figure A4 Tantalum nanotube surfaces dabbed dry...4 Figure A5 Tantalum nanotube surface with a water rinse..5 Figure A6 Water rinsed tantalum pentoxide surface showing nanotube openings below the precipitate.6 Figure A7 Acid then water rinsed tantalum pentoxide nanotubes..7

8 viii Figure A8 Close up photograph of acid then water rinsed tantalum pentoxide nanotubes..7 Figure A9 Tantalum pentoxide nanotube dimensions.8 Figure A10 Cross section of wear track..8

9 ix TABLE OF CONTENTS SIGNATURE PAGE... ii ACKNOWLEDGEMENTS... iii ABSTRACT.. iv LIST OF TABLES... vi INTRODUCTION... 1 PART A: In vitro Wear Rate Graded CoCrMo-Ti6Al4V Structures... 4 CHAPTER 1.Introduction Material and Methods Gradient Sample Fabrication In Vitro Metal Ion Release Results and Discussion Summary PART B: Wear Performance of Laser Processed Tantalum Coatings CHAPTER 1.Introduction Materials and Methods Results and Discussion Summary References.37 Appendix 1 In Situ Growth of Tantalum Nanotubes on Tantalum... A1 CHAPTER 1.Introduction... A1 2.Experimental... A1 3.Results and Discussion... A2 4.Summary... A8 Appendix 2 Experimental Protocol for Metal Ion Release Using AAS... A10 Appendix 3 Wear Testing of Metalic Biomaterials... A13

10 x Raw Nanotube Measurement Data... A16 REFERENCES... A9

11 INTRODUCTION Natural synovial joints, e.g., hip, knee or shoulder joints, are complex and delicate structures capable of functioning under critical conditions. Their performance is due to the optimized combination of articular cartilage, a load-bearing connective tissue covering the bones involved in the joint, and synovial fluid, a nutrient fluid secreted within the joint area. Unfortunately, human joints are prone to degenerative and inflammatory diseases that result in pain and joint stiffness. Primary or secondary osteoarthritis (osteoarthrosis), and to a lesser extent rheumatoid arthritis (inflammation of the synovial membrane) and condromalacia (softening of cartilage), are, apart from normal aging of articular cartilage, the most common degenerative processes affecting synovial joints. In fact, 90% of the population over the age of 40 suffers from some degree of degenerative joint diseases. Premature joint degeneration may arise from deficiencies in joint biomaterial properties, from excessive loading conditions, or from failure of normal repair processes. The explicit degenerative processes are not yet completely understood. Though minor surgical treatments are done to provide temporary relief to ailing patients, the ultimate need is to replace the dysfunctional natural joints by ceramic, metal or polymer based artificial materials by what is known as the total joint replacement (TJR) surgery. The choice of material for each component of such an implant depends on the design, size and required strength of the system. For total hip (THR) and total knee (TKR) joint replacements surgeries, metals are considered as the best candidate due to their higher load-bearing capabilities and higher fatigue resistance. The requirements for modern day metallic implants can be broadly categorized as follows: (i) superior biocompatibility between the material and surrounding environment with no adverse cytotoxicity and tissue reaction; and (ii) the mechanical and physical properties necessary to achieve the desired function. Some desired properties are low

12 2 modulus, high strength, good ductility, excellent corrosion resistance in body fluid medium, and high fatigue strength. The listed criteria are met by only a handful of metals and alloys. In the past, only stainless steel (316 and 316L) and cobalt based alloys (CoCrMo) were considered suitable for metallic implants. However titanium and its alloys started gaining popularity as implant materials in the early 1970 s due to their lower modulus, superior tissue compatibility and better corrosion resistance. Commercially pure (cp) titanium was the first to be used. Though cp-ti exhibited better corrosion resistance and tissue tolerance as compared to stainless steel, cp- Ti s rather limited strength confined its applicability to specific parts such as hip cup shells, dental crown and bridges, endosseous dental implants, pacemaker cases and heart valve cages. In order to improve the strength for load-bearing applications such as total joint replacements, the alloy Ti-6Al-4V ELI (i.e., with extra low interstitial impurity content) was chosen as a candidate biomaterial for surgical implants in the late 1970 s. Ti-6Al-4V is one of the most widely used Ti alloys which exhibits excellent corrosion resistance, low density, good biocompatibility, and excellent mechanical properties, including high strength and low modulus. Ti-6Al-4V has an elastic modulus of ~ 110 GPa which is only about half that of 316L stainless steel (~ 200 GPa) and CoCrMo alloys (~ 210 GPa). The mechanical properties of this alloy are critically dependent on its microstructure and can consequently be tailored by thermo-mechanical processing. While titanium is a better metal than steel it can be improved upon further. To this end researchers are modifying the surface properties of titanium by a number of different methods, which include laser pulses, ion implantation, diamond like carbon coatings and hydroxyapatite coatings (1; 2; 3; 4). The first modification discussed in this thesis is a gradient structure of

13 3 CoCrMo with a titanium base and the second is a laser engineered net shaping (LENS TM ) processed tantalum coating on titanium. CoCrMo alloy has high wear resistance (1) and is used on implant surfaces in high wear applications. Unfortunately CoCrMo alloy has low bioactivity (2), higher density and higher modulus than titanium. Biologically, CoCrMo alloy accumulates near the implant bone interface leading to osteolysis and potential loosening of the implant. Furthermore implant density is higher than bone, which causes patient discomfort due to an imbalance in weight distribution. Finally, higher modulus leads to stress shielding, in which bones remodel to minimize the weight required to support the loads the bone is experiencing. Stress shielding leads to reduced bone strength and possibly structural failure of the surrounding bone. This thesis is divided into 3 main parts. Part A is focused on the wear properties of gradient CoCrMo Ti-6Al-4V materials. Part B is focused on the wear properties of tantalum and how it compares to the wear properties of CoCrMo and Ti-6Al-4V. Finally the appendix focuses on preliminary work undertaken with tantalum nanotubes.

14 4 PART A: IN VITRO WEAR RATE GRADED COCRMO-TI6AL4V STRUCTURES CHAPTER 1. INTRODUCTION Osteolysis and aseptic loosening have been identified as major factors limiting the life of hip prostheses. Fine UHMWPE wear debris (3; 4), generated primarily at the interface between the femoral head and the acetabular cup, promote this degradation. Such wear-particle-induced bone loss is considered one of the key factors affecting the long-term stability of UHMWPE liner-based total hip replacements and other load-bearing implants (5; 6). Several attempts have been made to minimize the wear-induced osteolysis including the use of design modifications (7), UHMWPE property modification (8) and alternative bearing couples, such as metal-onmetal and ceramic-on-ceramic, eliminating the use of UHMWPE (9; 10). Among these approaches, metal-on-metal is a popular one and was used in this study. While metals have much lower wear rates than polyethylene they still produce wear particles. These particles lead to a biological response. Current multi-component load-bearing metal-on-metal hip implants suffer from low dynamic stability, interlocking mechanism problems, and possible malpositioning of the components during surgery including complicated/lengthy surgical procedures. Despite numerous literature documenting lower dislocation rates in association with enhanced soft-tissue repair techniques and the use of large diameter femoral heads (11; 12; 13; 14), hip instability/dislocation was the primary indication reported for a large percentage (22.5%) of revision total hip arthroplasty procedures (15). Therefore, improved stability for enhanced longevity is still an important performance criterion in the development of orthopedic prostheses necessitating the development of innovative designs and use of unitized/mono-block implants.

15 5 Another issue with implant design is patient fit. Currently only a few standard sizes are available and custom parts are expensive and rarely used. Poor patient fit can lead to a number of problems including incorrect axial orientation. Standard implant designs are not round they are oval as is the opening on the femur. This leads to one possible axial position which may not be ideal for a normal gait. Another issue with standard implants is poor patient fit. In standard implants the diameter slowly increases along the length of the shaft. This leads to a limited interface between bone and implant around the circumference of the implant at the top of the shaft. With a custom implant contact occurs along the entire length of the shaft and there is no chance of axial misorientation. Custom implants using conventional manufacturing methods are expensive and tooling must be reset for each implant. A free form fabrication process such as LENS TM does not require these manual tooling changes. In LENS TM processing a cat scan is taken of the leg to form a three dimensional image. This image is transformed into a CAD file where the internal shape of the femur can be used to design the stem. Once designed on the computer the implant part can be appropriately loaded to determine weak points in the implant part and the design adjusted. Then the part can be tested in the software to insure that it slides smoothly into the femur. Upon completion of design the part is digitally sliced and part production on the LENS TM can continue. Considering developments in design and manufacturing, there is a growing interest in the fabrication of novel structures to mimic multiple tissues and tissue interfaces on the same implant, such as implants with gradients in materials, porosity and pore sizes that will allow on one side of the implant high vascularization and direct osteogenesis, while promoting osteochondral ossification on the other-side. In this direction, innovative designs such as

16 6 functionally graded acetabular shells with open porosity on one side in contact with the bone to improve cell material interactions, and a hard ceramic/alloy coating on the other side in contact with UHMWPE liner to reduce liner wear can significantly improve an implant s in vivo life. The benefit of using such unitized structures for metal-on-metal configuration is obvious such as the ability to use the larger femoral head diameter for a given acetabular diameter, which not only increases the range of motion and reduces the risk of dislocation, but also optimizes the tribology for decreased wear for this bearing couple (16; 17). We have successfully fabricated and evaluated in vitro biocompatibility of compositionally and structurally graded CoCrMo- Ti6Al4V (18), TiO2-Ti (19) and ZrO2-Zr-Ti (20) unitized structures using laser engineered net shaping (LENS ) to minimize wear induced osteolysis. Other load-bearing implant design concepts that can potentially be implemented via LENS processing can be found in work by F.A. España et al. (21). 2. MATERIALS AND METHODS Gradient Sample Fabrication Ti6Al4V alloy powder (Advanced Specialty Metals, Inc., Nashua, NH) with particle size between 50 and 150 μm and CoCrMo alloy powder (Stellite Coatings, Goshen, IN) with 50 to 100 μm particle size was used in this study. Gradient samples of 10 mm diameter were fabricated on a substrate of 3 mm thick rolled, commercially pure Ti plates using LENS -750 (Optomec Inc. Albuquerque, NM) equipped with a 500W Nd:YAG laser and a double powder feeder system. Samples were fabricated in a glove box containing argon atmosphere with O 2 content less than 10 ppm to limit oxidation of alloys during processing. Gradient pins with 50%, 70% and 86% CoCrMo alloy on the top surface were made at 450 W laser power and 22.5 mm/s scan

17 7 speed. These compositionally graded structures consisted of porous (~ 30% porosity) 100% Ti6Al4V alloy in the first six layers. The composition in the transition region of the gradient structure was varied from 100% Ti6Al4V alloy at the first layer to various concentrations of CoCrMo alloy at the top layer over 5 6 layers. The porous region of first six layers of Ti6Al4V alloy were made at 200 W (to create 30 % porosity) and remaining layers were made at 450 W to achieve fully dense structures in the transition region and top surface of the gradient structures. In the transition zone, the powder feed rate for CoCrMo alloy increased from 0 to 46.5 g/min over the number of transition layers. For Ti6Al4V alloy, it was 17.5 g/min to start with and reduced to 0 g/min over the same number of layers. Complete process details and optimization details have been provided previously (18). For comparison, 100% CoCrMo alloy pins (ϕ 10 mm, 15 mm height) and 18 mm square blocks with 30 mm height were also made at 450 W laser power, 22.5 mm/s scan speed and 46.5 g/min feed rate. CoCrMo and Ti6Al4V gradient structures were etched with Krolls reagent. 100% CoCrMo was electrolytically etched with at 2 volts with 5% HCL for 30 seconds. Top surface microstructures of the samples were examined using both optical and scanning electron microscopy (SEM). The average composition of the top surface of gradient structures was measures using a field emission SEM (FEI-SIRION) equipped with a Genesis EDAX energy-dispersive spectrometer (EDS). The hardness of gradient pins and laser processed 100% CoCrMo alloy was measured using Vickers micro hardness tester (Leco, M-400G3) at 200 g load applied for 15 s In Vitro Wear Testing In vitro wear testing was performed using 100% CoCrMo alloy pins (here after referred as P100) and gradient pins with 50, 70 and 86% CoCrMo alloy on top surface (here after referred

18 8 as GP50, GP70 and GP86, respectively) rubbing against 3 mm thick square discs of UHMWPE and 100% CoCrMo alloy (here after referred as S100). The UHMWPE discs had dimensions of 25 mm x 25 mm x 3 mm, and 3 mm thick discs were cut from 18 x 18 mm square block of 100% CoCrMo alloy using a high pressure water jet. The top surface of the pins were machined to a radius of 40 mm and further ground using series of SiC grinding papers with various sizes up to 1200 grit. Samples were then polished on velvet cloth using a series of Al 2 O 3 powder up to 1 μm suspended in distilled water a similar procedure was used to prepare the surfaces of UHMWPE and 100% CoCrMo alloy discs. Finally, just before wear testing, all the samples were ultrasonically cleaned in an alcohol bath. The surface roughness of each test samples was measured using a surface profilometer (SPN Technology, Goleta, CA). Linear reciprocating spherical tip pin-on-disc wear testing, according to ASTM G133, was performed using a tribometer (NANOVEA, Microphotonics Inc., CA, USA) with pins having 40 mm radius of curvature rubbing against UHMWPE (metal-on-pe) and 100% CoCrMo alloy discs (metal-on metal). A linear oscillatory motion 10 mm in length (the full cycle represents 20 mm of travel) with a speed 1000 mm/min was used. Wear tests were carried out in simulated body fluid at 37±1 C to represent the biological environment in the body. The wear tests were performed with a normal load of 5 N for sliding distances of 1 and 3 km (equivalent to 150,000 cycles). The width of the wear track was measured using SEM images of worn pins and substrates, and from the known curvature of the pin and linear oscillatory stroke length the wear track volume was estimated. The wear rates of UHMWPE, 100% CoCrMo alloy and gradient sample pins were then calculated in mm3/nm.

19 In Vitro Metal Ion Release At the conclusion of each wear test run, the remaining SBF was collected analyzed for Co ion concentration. The resulting sample volumes varied from 50 ml to 200 ml. Each sample was treated with 5 ml 1 N hydrochloric acid and 1 ml concentrated nitric acid. The liquid volumes were reduced to below 30 ml on a hot plate. Samples were removed after the solution turned yellow or when they were reduced to 15 ml. Then, deionized water was added to the samples to increase the volume to 100 ml. This vigorous digestion process was used to ensure that all the Co was fully ionized prior to analyzing the samples by flame atomic absorption spectrometry (AAS). The concentration of Co in the samples was analyzed on a Shimadzu 6300 flame AAS with a cobalt lamp. Each sample was analyzed twice with an average taken unless they differed by more than 5% than a third aliquot was analyzed and the closest two were averaged. Each measurement was taken for 5 seconds with a 20 sec stabilization period prior to analysis. A linear 5 point standard curve at 0, 0.1, 0.2, 0.5 and 1.0 ppm of cobalt was used. A method spike standard curve was also prepared to ensure the SBF matrix did not cause interference with determination of cobalt concentrations. A blank was also analyzed in AAS for every 10 samples to ensure baseline stability. 3. Results and Discussion In vitro wear rate or wear coefficient, in mm 3 /Nm, of pin and substrate materials was calculated from wear track dimensions measured using SEM images of worn pins and substrates. Figure 1 shows the wear coefficient of UHMWPE discs tested against gradient metal pins with varying concentration of CoCrMo alloy on top surface. Under present experimental conditions, the average wear coefficient of UHMWPE against metallic pins varied between 1.70 x 10-6

20 10 N=2 Figure 1 Influence of CoCrMo alloy concentration, in the top surface of metallic pins, on the wear rate of UHMWPE. mm 3 /Nm and 9.35 x 10-6 mm 3 /Nm. No detectable wear or wear tracks were observed on the metallic pins. However, occasional scratches were found on the pins surface. For all metal-on-pe combinations, the wear coefficient decreased with an increase in the sliding distance. The high wear coefficients during initial 1000 m sliding distance is due to higher initial running-in wear With an increase in the sliding distance, the system enters in to steady-state wear regime, during this period the wear rate gradually decreases and reaches a constant value. A similar decrease in the wear coefficient with an increase in sliding distance has been reported earlier (22; 23). No significant difference in the wear rate of UHMWPE was observed when tested with P100 pin, GP70 and GP86 pins. However, tests with GP50 pins always resulted in higher wear rates of UHMWPE. Since the test conditions are the same for all metal-on-pe couples, the

21 Figure 2 Top surface microstructures of the pins with varying concentrations of CoCrMo alloy in the top layer: (a) GP50, (b) GP86 and (c) P

22 12 Figure 2 Continued Top surface microstructures of the pins with varying concentrations of CoCrMo alloy in the top layer: (c) P100. observed differences in wear coefficients of UHMWPE can be explained on the basis of microstructural features present in the top surface of different pins. Top surface microstructural variation as a function of CoCrMo alloy concentration in graded pins is shown in Figure 2. Ti6Al4V and CoCrMo were the only two materials in the gradient structure, previous work showed they do not form intermetallics therefore the material which had a greater volume in 86% CoCrMo, as identified in Figure 2, was CoCrMo. This is confirmed by the observation that Ti6Al4V is recessed which makes sense because it is a softer material. The 100% CoCrMo alloy contains carbon so the elevated material is carbide since it is harder. As the CoCrMo alloy content increased, the most significant change in the micrographs was the decrease in the volume fraction of the Ti6Al4V alloy phase and the consequent increase in the volume fraction of

23 13 CoCrMo alloy. Notable changes were also observed in the morphology and scale of the CoCrMo alloy phase as the amount was increased. As shown in Figure 2a, the GP50 pins exhibited a bimodal distribution of CoCrMo phase with large chunks of CoCrMo phase and some thin, continuous phase predominantly along Ti6Al4V alloy grain boundaries. With an increase in the concentration of CoCrMo alloy the grain boundary precipitation/phase was completely replaced by a two-phase microstructure. The laser processed 100% CoCrMo alloy (P100 and S100) showed very fine carbide precipitates in the interdendritic regions (Figure 2c). We hypothesize that the microstructure of GP50 pins with large chunks of hard CoCrMo alloy phase acted as isolated abrasive particles leading to local grooving and high wear rate of UHMWPE. Also, the local grooving resulted in large variations in the measured wear rates with GP50. Absence of chunks of isolated CoCrMo alloy phase in the pins with high concentration of CoCrMo eliminated the grooving due to a large area of contact between the pin and the discs. These microstructural features decreased the wear rate of UHMWPE with increasing concentration of CoCrMo alloy on the top surface of the gradient pins. Further, the observed microstructural differences between the gradient pins resulted in some variation in the surface roughness of the as-polished gradient pins before wear testing. The average surface roughness of P100/S100 was found to be 5.0 ± 0.4 μm and the gradient pins showed an average surface roughness of 5.4 ± 0.3 μm, 6.2 ± 0.4 μm and 7 ± 0.9 μm for GP86, GP70 and GP50, respectively. Therefore, the large chunks of hard CoCrMo alloy phase in GP50 pins resulted in high wear rate of UHMWPE due to its high average surface roughness compared to other gradient pins. Typical wear track morphologies of UHMWPE discs are shown in Figure 3. It can be seen from Figure 3a that specimens tested with a sliding distance of 1000 m show the presence

24 Figure 3 Typical morphologies of wear tracks on UHMWPE when tested against (a) GP50 pin, 1000 m (b) GP50 pin, 3000 m and (c) GP86 pin, 3000 m. 14

25 15 Figure 3 Continued Typical morphologies of wear tracks on UHMWPE when tested against (c) GP86 pin, 3000 m. of heavy wear tracks/grooves on their worn surfaces and at 3000 m the wear tracks were shallow and smoother (Figure 3b) indicating lower wear rates at 3000 m. It is also important to note that the no significant grooving was observed on the worn specimens when tested with P100 or GP86 pins (Figure 3c). The observed results suggest that the wear rate of UHMWPE during articulation against metal primarily depends on microstructural features such as morphology and distribution of hard phase within the matrix, which controls the surface roughness of gradient pins during polishing. Also, the wear rate of UHMWPE does not change significantly, compared to P100 (100% CoCrMo alloy), when tested against the new articulating surfaces developed in the present work, especially with CoCrMo alloy concentrations ³ 70%.

26 16 Although, the wear performance of present gradient structures, with ³ 70% CoCrMo on top, against UHMWPE was comparable with that of P100 (100% CoCrMo alloy), the newly developed articulating surfaces might show superior wear characteristics during metal-on-metal articulation. As stated earlier, the real benefits of present gradient/unitized structures comes from its use as metal-on-metal configuration, which include the ability to use a larger femoral head diameter to increase the range of motion and reduce the risk of dislocation (16; 17). To evaluate this hypothesis, the pins with varying concentrations of CoCrMo on top were tested against laser processed S100 (100% CoCrMo alloy) discs. The wear rate of gradient pins and P100 pins under metal-on-metal articulating condition i.e., pins rubbing against laser processed 100% CoCrMo alloy discs, is shown in Figure 4. For both the sliding distances of 1000 m and 3000 m, P100 pins always showed highest wear rate between x 10-8 mm 3 /Nm and 1.68 x 10-7 mm 3 /Nm. Lowest wear rate in the range of 5.07 to 7.99 x 10-8 mm 3 /Nm was observed for GP86 pins. The wear rate was found to increase with a decrease in the top surface CoCrMo alloy concentration; results correlate well with the observed top surface hardness of various pins. The average top surface hardness of gradient pins increased from 615 ± 38 HV (GP50) to 739 ± 25 HV (GP70) to 957 ± 32HV (GP86) with increase in the CoCrMo concentration. Therefore, the wear rate decreased with an increase in the CoCrMo alloy concentration of the pin s surface. Also, the hardness of these gradient pins was significantly higher than the average hardness of laser processed 100% CoCrMo alloy discs/pins (S100 and P100), which was 361±11 HV. Similar hardness of P100 pins with that of S100 disc resulted in its higher wear of P100 pins compared to that of gradients pins with harder top surface. These wear test results show that present laser processed gradient structures have superior wear resistance up to 1000 m and comparable, if not better, wear resistance up to 3000 m than that of 100% CoCrMo alloy. Figure 5 shows the wear

27 17 N=2 Figure 4 Wear rate of laser processed P100 and CoCrMo-Ti6Al4V alloy gradient pins. N=2 Figure 5 Wear rate of laser processed S100 substrate tested for 3000 m against various gradient pins.

28 18 rate of S100 discs tested for 3000 m against P100 pins and gradient pins. The wear rate of the substrate was high when tested against gradient pins and the wear rate increased with a decrease in the CoCrMo alloy concentration in the pin s top surface. This trend is similar to that observed for UHMWPE and can be directly correlated with the microstructural features and consequent surface roughness of the pins. From the experimental data it appears that the wear rate of discs (100% CoCrMo alloy) depends on the hardness and microstructural features of articulating surfaces. With P100 and S100 combination the similar hardness and microstructural features of both the articulating surfaces resulted in lower wear rates of S100. However, the wear rate of S100 discs during articulation against gradient pins was found to strongly depend on the microstructural features of gradient pins. Although the hardness of GP50 is less than GP86/GP70 the large chunks of hard CoCrMo alloy phase in GP50 pins resulted in high wear rate of S100 due to its high average surface roughness compared to other gradient pins. Figure 6 shows typical wear track morphologies of CoCrMo-Ti6Al4V alloy gradient pins and S100 discs during metal-on-metal wear testing. The wear tracks of GP50 pins (Figure a) showed deep scratches with large chunks of hard CoCrMo alloy phase projecting out of soft Ti6Al4V alloy phase, which was worn during wear testing. However, GP86 pins, due to their relatively uniform microstructural features, showed no deep scratches and the wear track was smoother than observed on GP50 pins. The deeper wear tracks of GP50 pins compared to GP86 pins supports its higher wear rate than that of GP86 pins (Figure 4). The wear track morphologies of S100, shown in Figure 4c and 6d, are similar to those observed on UHMWPE - heavy wear tracks/groves when tested with GP50 pins, and no significant grooving when tested

29 Figure 6 Typical morphologies of wear tracks observed during metal-on-metal wear testing (a) on GP50 pin, 1000 m (b) on GP86 pin, 1000 m and (c) on S100 substrate tested with GP50 19

30 Figure 6 Continued Typical morphologies of wear tracks observed during metal-on-metal wear testing (c) on S100 substrate tested with GP50 20

31 21 Table 1 Co ion concentration (ppm) in the SBF after wear testing (*below the detection limit of the instrument). Metal-on-Metal CoCrMo Concentration 100% 86% 70% 50% 1000 m m Metal-on-PE 1000 m * * * 3000 m * * with GP86 pins. The metal-on-metal wear morphologies also suggest that the wear rate of counter articulating surface depends on morphology and distribution of hard phase within the matrix. Table 1 shows the concentration of Co in SBF for metal-on-metal and metal-on- UHMWPE couple after wear testing. As expected the Co release was high during metal-on-metal wear testing compared to that of metal-on-uhmwpe. The Co concentration was found to be in the range of 0.25 and 0.91 ppm for metal-on-metal articulation. After 3000 m of sliding the Co concentration was slightly higher for gradient pins than observed with laser processed 100% CoCrMo alloy pins. This is intuitive, as the wear of S100 discs was high when tested with gradient pins due to their relatively high hardness than the discs. The higher concentration of Co in the SBF for gradients pins with GP50 pins is also attributed to the high wear rate of these pins

32 22 and respective S100 discs. It is also important to note that the present Co concentrations in SBF were significantly lower than observed during previous study on the pin-on-disc wear testing of CoCrMo alloys (24), where the reported Co concentration was between 2 and 13 ppm. Therefore, the present CoCrMo-Ti6Al4V alloy gradients structures showed enhanced in vitro wear resistance with a decrease in Co ion release. 4. Summary Compositionally graded hard and wear-resistance CoCrMo alloy coatings on porous Ti6Al4V alloy have been prepared using LENS process. In vitro wear testing of these gradient structures against UHMWPE discs showed comparable wear rates for both the gradient coatings and 100% CoCrMo alloy. The wear of UHMWPE was found to depend on microstructural features such as CoCrMo alloy phase morphology and distribution at the contact surface. For metal-on-metal articulation, the wear rates of gradient structures were low compared to that of 100% CoCrMo alloy due to their relatively high hardness. The wear rate of 100% CoCrMo alloy discs was found to be a combined effect of surface hardness and microstructural feature present in the top surface of gradient structures. Material combinations that resulted in high amount of total wear also lead to higher Co concentration in SBF after wear testing. These laser-processed gradient structures with CoCrMo alloy concentrations 70%, having low wear rates of the order of 10-8 mm 3 /Nm, show their potential to reduce wear induced osteolysis when used in mono-block structures.

33 23 PART B: WEAR PERFORMANCE OF LASER PROCESSED TANTALUM COATINGS CHAPTER 1. INTRODUCTION The first part of this work (18) showed that functionally graded, hard CoCrMo alloy coatings on porous Ti6Al4V alloy can significantly increase the surface hardness without deleterious effects on in vitro biocompatibility. During microstructural analysis, notable changes were observed in the morphology and scale of the CoCrMo alloy phase as the concentration of CoCrMo alloy in the top surface of the gradient structures was increased (18). It is generally accepted that the wear resistance of any alloy is primarily governed by microstructural features, such as the amount, morphology and distribution of hard particles in the matrix. Therefore, in the present investigation, in vitro wear performance and metal ion release rate of CoCrMo-Ti6Al4V gradient structures with different CoCrMo alloy concentrations in the top surface were evaluated. The wear rate and Co metal ion release was determined for these gradient structures rubbing against UHMWPE and 100% CoCrMo to simulate different articulating conditions. Over the last few decades researchers have attempted to find a bioactive metal with high mechanical strength that can simultaneously bond chemically with surrounding bone on one side and provide hard/wear resistance surface on the other side for orthopedic applications. As it is quite difficult to achieve a high degree of bioactivity and wear resistant properties in one material, a popular approach is to fabricate implants with adequate biological properties followed by a special treatment to enhance their surface properties such as wear resistance. Furthermore, the surface properties can be selectively modified to enhance site-specific biological and/or

34 24 tribological performance of the implants for a variety of orthopedic applications. Because of these reasons, research on surface modification of metallic biomaterials has attracted much attention to improve multifunctionality, tribological and mechanical properties, as well as biocompatibility of artificial metal implants while obviating the need for large expenses and long time to develop new metallic biomaterials. A number of different surface modification techniques have been employed for the preparation of metallic implants to improve biological (25; 26; 19; 27) and tribological properties (20; 8; 28; 29). Among these techniques laser surface modification can offer high degree of process controllability and flexibility. There is a growing amount of published work that testifies to the potential of lasers for altering the surface properties of metallic biomaterials in order to improve their biological and tribological properties (30; 31). Recently there has been renewed interest in tantalum due to its excellent ductility, toughness, corrosion resistance, and bioactivity (32; 33). However, a relatively high cost of manufacture and inability to produce a modular all tantalum implant limited its wide spread applicability. Further, high affinity for oxygen and extremely high melting temperature (3017 C) make it difficult to process tantalum-based coatings or implant structures via conventional processing routes. Nevertheless, recently we have demonstrated fabrication of dense Ta coatings on Ti (34) and bulk porous Ta structures (35), with total porosities between 27% and 55%, using high-power lasers in LENS a solid freeform fabrication technology. Excellent cell survivability, cellular attachment and spreading in terms of high living cell density, intensity and distribution of vinculin expression, respectively, demonstrated superior biocompatibility of laser processed Ta than Ti (34; 35). In other studies (36; 37) it has been shown that laser processed dense Ta coatings provide comparable in vitro and in vivo cell-materials interactions to that of bioactive hydroxyapatite (HA) coatings,

35 25 however, with superior interfacial characteristics and mechanical stability. These results demonstrate that Ta structures offer a favorable biological environment for adhesion, growth and differentiation of human osteoblasts, which can promote biological fixation. Although, HA coatings perform equally well in terms of biological fixation, the low fracture toughness, wear resistance and rough surface morphology of these coatings can generate large amounts of wear debris as a result of micro motions at the bone-implant interface during early stages of surgery. These particulate wear debris have been recognized as one of the major cause of load-bearing implants failures (38; 39). Wear debris can be alleviated if HA coatings are replaced with dense bioactive Ta coatings, as the metallic Ta coatings have high fracture toughness and can be polished to achieve a smooth surface. Therefore, in chapter 3 of this study the in vitro wear performance of laser processed Ta coatings on Ti was evaluated and compared with commercially pure (CP) Ti and CoCrMo alloy. 2. MATERIALS AND METHODS A Ta metal powder (Grandview Materials Inc., Columbus, OH) of 99.5% purity and with particles size between 45 and 75 μm was used. Ta coatings of mm were deposited on a substrate of 3 mm thick rolled CP Ti plates using a LENS -750 (Optomec Inc., Albuquerque, NM) equipped with a 500 W continuous wave Nd:YAG laser. Detailed descriptions of LENS operation and capabilities can be found elsewhere (40; 41). All coatings were fabricated in a glove box containing an argon atmosphere with O 2 content of less than 10 ppm to limit oxidation of materials during processing. Dense Ta coatings were prepared using a laser power of 450 W, a scan speed of 7 mm s 1 and a powder feed rate of 106 g min 1 (34). For comparison, CoCrMo alloy powder (Stellite Coatings, Goshen, IN) with μm particle size

36 26 was used to make 15 mm square blocks with 30 mm height at 450 W laser power, 22.5 mm s -1 scan speed and 46.5 g min -1 feed rate (8). These blocks were removed from the substrate and disc samples of 2 mm thick were then cut using a water jet for further testing and characterization. The Ta coatings, CP Ti substrate and CoCrMo disc samples were abraded using successive grades of 400, 600, and 1200 grit silicon carbide paper and then ultrasonically cleaned in distilled water and dried at room temperature. Final polishing was done using a cotton polishing cloth with a 0.05 µm alumina suspension and the samples were cleaned and used for wear testing. Finally, just before wear testing, all the samples were ultrasonically cleaned in an alcohol bath. The surface roughness of each test samples was measured using a surface profilometer (SPN Technology, Goleta, CA). Linear reciprocating ball-on-disc wear testing, according to ASTM G133, was performed using a tribometer (NANOVEA, Microphotonics Inc., CA, USA) with 3 mm hardened chrome steel ball (100Cr6, HRC) rubbing against test samples. A linear oscillatory motion 10 mm in length (the full cycle represents 20 mm of travel) with a speed 1200 mm/min was used. The wear tests were performed with a normal load of 5 N for sliding distance of 1 km. The width of the wear track was measured using SEM images of test samples, and from the known curvature of the ball and linear oscillatory stroke length the wear track volume was calculated. The wear rate of each test samples was reported in mm 3 Nm 1. All tests were carried out in aseptic conditions in freshly prepared simulated body fluid (SBF) at 37 ± 1 C to represent the biological environment in the body. The ionic concentration of the SBF used in the present study is as follows: 2.5 mm of Ca 2+, 1.5 mm of Mg 2+, mm of Na +, 5.0 mm of K +, mm of Cl, 4.2 mm of HCO 3, 1.0 mm of HPO 2 4, 0.5 mm of SO 2 4. The top surface hardness of the test samples was measured using a Vickers microhardness tester

37 27 (Shimadzu, HMV-2T) at 200 g load for 30 s and the average value of 10 measurements was reported. Top surface microstructures of the polished and etched samples were examined using optical microscope and wear track morphology was observed using scanning electron microscopy (SEM). 3. RESULTS AND DISCUSSION In metal-on-metal implants, wear of articulating surfaces is a combined effect of adhesive and abrasive wear mechanisms. Fatigue wear is also possible in which repetitive loading of localized regions causes surface and subsurface cracks to propagate thus producing wear debris. Under such conditions, wear resistance of any alloy is primarily governed by microstructural features, such as grain size, the amount, morphology and distribution of second phase particles in the matrix. Therefore, the initial microstructures of laser processed Ta coatings, CoCrMo alloy and as-received CP Ti substrate was observed using SEM. Figure 7 shows typical microstructural features of three materials used in the present work. The grain size of the samples was measured using linear intercept method. All the materials showed equiaxed grains and the laser processed CoCrMo alloy showed very fine carbide precipitates (not shown here) in the interdendritic regions. The grain size of laser processed CoCrMo alloy was significantly finer than laser processed Ta coatings. As shown in Table 2, laser processed Ta coatings exhibited an average grain size of µm, which is significantly larger than the grain size of laser processed CoCrMo alloy with an average grain size of µm. The CP Ti substrate showed an average grain size of µm, comparable to that of Ta coatings. Laser processing is characterized by extremely high cooling rates, of the order of 10 3 to 10 5 K s -1, which influences

38 28 (a) (b) Figure 7 Typical microstructures of (a) laser processed Ta coatings, (b) laser processed CoCrMo and (c) as-received CP Ti substrate

39 29 Figure 7 Continued Typical microstructures of (c) as-received CP Ti substrate. (c) Table 2 Grain size (µm), surface roughness (µm) and hardness (HV 0.2) of Ta, CoCrMo and CP Ti. Laser energy density (J mm -2 ) used to fabricate Ta and CoCrMo alloy is also included. Material Energy Density Grain Size Hardness Roughness Laser processed Ta Laser processed CoCrMo As-received CP Ti substrate ± several aspects of liquid metal solidification. It is well known that the scale of dendritic/grain structure is inversely proportional to the solidification cooling rates (23). In general, lower heat/energy input increases the thermal gradients near the melt zone; consequently high cooling rates can be achieved. Therefore, the observed differences in grain size between laser processed Ta and CoCrMo alloy can be directly correlated to the heat/energy input used during deposition.

40 30 The finer grain size of CoCrMo alloy compared to Ta is attributed to the lower energy put, 17 J mm -2, used during deposition, while the Ta was deposited at an energy input of 54 J mm -2. Although the local cooling rate is not directly measured in the present work, within the melt zone the cooling rate, dt/dt (K s -1 ), can be expressed as the product of solidification velocity R (mm s - 1 ) and the local temperature gradient G (K mm -1 ) (42). Experimentally it has been shown that R is on the order of laser scan speed (v) (24) and G is on the order of ~ 100 K mm -1 (43). Thus under present experimental conditions, the cooling rate in the melt zone is expected to be around ~ 700 K s -1 and 2250 K s -1, for Ta and CoCrMo alloy deposition, respectively. Therefore, the finer grain size of CoCrMo alloy is a direct consequence of high cooling rates achieved at lower energy input of 17 J mm -2. Figure 8 shows experimentally determined wear rate of CP Ti substrate and laser processed Ta and CoCrMo alloy for a sliding distance of 1 km at a normal load of 5 N. The average wear rate of CP Ti, Ta and CoCrMo alloy was found to be mm 3 (N.m) -1, mm 3 (N.m) -1 and mm 3 (N.m) -1, respectively. The experimental data clearly indicate that laser processed CoCrMo alloy has superior wear resistance compared to CP Ti and laser processed Ta. It can be seen from the data that after 1 km sliding, CoCrMo alloy exhibited roughly two orders of magnitude lower wear rate than CP Ti substrate and approximately one order of magnitude less wear than laser processed Ta. Similarly, the laser processed Ta showed one order of magnitude less wear rate than CP Ti substrate. The wear rate of these materials is further substantiate by their coefficient of friction values, which were found to be 1.42, 0.97 and 0.65 for CP Ti, Ta and CoCrMo alloy, respectively. Under present experimental conditions, with a normal load of 5 N, the initial maximum Hertzian

41 31 N=2 Figure 8: Wear rate of CP Ti, laser processed Ta and CoCrMo alloy against hardened 100Cr6 steel ball. contact pressure was found to be 1.40 GPa, 1.74 GPa and 1.89 GPa for CP Ti, Ta and CoCrMo alloy, respectively. In spite of lowest contact pressure the CP Ti substrate showed a higher wear rate. The superior wear resistance of CoCrMo alloy is attributed to its high hardness, fine grain size and presence of fine carbide precipitates (44; 45). The observed wear rates of these materials correlate well with their respective top surface hardness values. As shown in Table 2, the average top surface hardness of CoCrMo alloy was HV, which is significantly higher than the hardness of Ta ( HV) and CP Ti substrate (140 4 HV). Generally the wear resistance of a material is directly related to its hardness, i.e., higher the hardness higher the wear resistance. Therefore, the wear rate decreased in the order of CP Ti Ta CoCrMo alloy. However, it is important to note that the wear rate of laser processed Ta is considerably higher

42 32 (a) Figure 9 Wear track morphology after 1 km of sliding distance (a) CP Ti substrate (b) laser processed Ta and (c) laser processed CoCrMo alloy. (b)

43 33 Figure 9 Continued Wear track morphology after 1 km of sliding distance (c) laser processed CoCrMo alloy. (c) than CP Ti substrate, suggesting that the Ta coatings on Ti can provide better wear protection during implant service. Figure 9 presents the worn surfaces of CP Ti, Ta and CoCrMo alloy samples after 1 km sliding. As shown in Figure 9a, deep worn tracks were observed on CP Ti substrate compared to relatively shallow and smoother worn tracks on laser processed Ta surface (Figure 9b). The worn surface of CoCrMo alloy, shown in Figure 9c, was very smooth with minimum amount of

44 34 (a) (b) Figure 10 Continued High magnification SEM images of worn areas for (a) CP Ti substrate (b) laser processed Ta and (c) laser processed CoCrMo alloy.

45 35 Figure 10 Continued High magnification SEM images of worn areas for (c) laser processed CoCrMo alloy. (c) grooves, which are barely visible. From the high magnification SEM images of the worn tracks, shown in Figure 10, parallel arrays of grooves oriented along the sliding direction, were observed on all samples. However, the grooves are very deep on both CP Ti and Ta samples surfaces indicating abrasive type wear. On the other hand, the smooth nature of CoCrMo alloy worn surfaces suggests that this material has undergone an adhesive type of wear. The degree of formation of wear debris, which caused third body wear, appeared to be greater on Ta and CP Ti samples. Present results show that laser processed Ta coatings exhibit superior in vitro wear resistance compared to CP Ti substrate. Further, present Ta coatings can potentially minimize the early-stage bone-implant interface micro motion induced wear debris generation as a result of their excellent bioactivity, high wear resistance and toughness compared to popular HA coatings.

46 36 4. SUMMARY For the first time, laser processed Ta coatings have been evaluated for their in vitro wear resistance to explore their potential to be used in contact surfaces for load-bearing implant applications. The average in vitro wear rate of Ta coatings was found to be mm 3 (N.m) -1, which is one order of magnitude less than CP Ti substrate with a wear rate of mm 3 (N.m) -1. Although the wear performance of Ta coatings is inferior to that of widely used CoCrMo alloy, its superior in vitro wear resistance than Ti demonstrates their potential to replace current HA coatings on Ti implants with better long-term in vivo stability.

47 REFERENCES 1. Z. Jin, P. Firkins, R. Farrar and J. Fisher. 2000, Proc. Instn. Mech. Engrs., Vol. 214 Part H, p S Jakobsen, J Baas, T Jakobsen, K Soballe. 2010, Journal of Biomedical Materials Research Part A, Vol. 94A, p R. Cooper, C. McAllister, L. Borden, T. Bauer. 1992, J Arthroplasty, Vol. 7, p D. Howie, S. Rogers, D. Haynes, M. Pearcy. 1994, Trans. Orthop. Res. Soc., Vol. 19, p T. Schmalzried, L. Kwong, M. Jasry, R. Sedlacek, T. Haire, D. Connor. 1992, Clin. Orthop. Rel. Res., Vol. 274, p A. Edidin, C. Rimnac, V. Goldberg, S. Kurtz. 2001, Wear, Vol. 250, p D. Bartel, V. Bicknell, T. Wright. 1986, J Bone Joint Surg., Vol. 68A, p S. Dittrick, V. Balla, N. Davies, S. Bose, A. Bandyopadhyay. 2010, Mater. Sci. Eng. C, Vol. in press doi: /j.msec C. Rieker, P. Kottig. 2002, Hip Int., Vol. 12, p N. Douglas, P. Christopher, P. Javad, J. Daniel, E. Stefan, B. Andre. 7, Suppl. 2, 2004, J Arthroplasty, Vol. 19, p D. Chivas, K. Smith, M. Tanzer. 2006, Clin. Orthop. Relat. Res., Vol. 453, p P. Lachiewicz, E. Soileau. 2006, Clin. Orthop. Relat. Res., Vol. 453, p R. Sierra, J. Raposo, R. Trousdale, M. Cabanela. 2005, Clin. Orthop. Relat. Res., Vol. 441, p R. White, T. Forness, J. Allman, D. Junick. 2001, Clin. Orthop. Relat. Res., Vol. 393, p K. Bozic, S. Kurtz, E. Lau, K. Ong, T. Vail, D. Berry. 2009, J Bone Joint Surg. Am., Vol. 91, p J. Cuckler, K. Moore, A. Lombardi Jr, E. McPherson, R. Emerson. 8 suppl 3, 2004, J Arthroplasty, Vol. 19, p Cuckler, J. 2005, Clin. Orthop. Relat. Res., Vol. 441, p B. Krishna, W. Xue, S. Bose, A. Bandypadhyah. 2008, Acta biomater., Vol. 4, p. 697.

48 V. Balla, P. DeVasConCellos, W. Xue, S. Bose, A. Bandyopadhyay. 2009, Acta Biomater., Vol. 5, p V. Balla, W. Xue, S. Bose, A. Bandyopadhyay. 2009, Acta biomater., Vol. 5, p F. España, V. Balla, S. Bose, A. Bandyopadhyay. 2010, Mater. Sci. Eng. C, Vol. 30, p C. Ramesh, A. Khan, N. Ravikumar, P. Savanprabhu. 2005, Wear, Vol. 259, p L. Yang. 2005, Wear, Vol. 259, p K. John, L. Zardiackas, R. Poggie. 2004, J Biomed. Mater. Res. Part B: Appl. Biomater., Vol. 68B, p V. Balla, S. Banerjee, S. Bose, A. Banyopadhyay. in press 2010 doi:101016/j.msexc , Material Science Engineering C. 26. M. Roy, V. Balla, A. Bandyopadhyay, S. Bose. 2008, Acta Biomater, Vol. 4, p R. Narayanan, S. Seshadri, T. Kwon, K. Kim,. 2008, J Biomed. Mater. Res. Part B Appl. Biomater., Vol. 85B, p F. Platon, P. Fournier, S. Rouxel, Wear 250 (2001) , Wear, Vol. 250, p C Liu, Q Bi, A Matthews. 2003, Surf. Coat. Technol., Vols , p A Kurella, B Narendra. 2005, J Biomater. Appl., Vol. 20, p L. Hao, J. Lawrence. Laser Surface Treatment of Bio-Implant Materials. England : John Wiley & Sons, M. Bermudez, F. Carrion, G. Nicolas, R. Lopez. 2005, Wear, Vol. 258, p H. Kato, T. Nakamura, S. Nishiguchi, Y. Matsusue, M. Kobayashi, T. Miyazaki, H.M. Kim, T. Kokubo. 2000, J Biomed. Mater. Res., Vol. 53, p V. Balla, S Banerjee, S Bose, A. Bandyopadhyay. 2010, Acta Biomater., Vol. 6, p V. Balla, S. Bodhak, S. Bose, A. Bandyopadhyay. 2010, Acta Biomater., Vol. 6, p M. Roy, V. Balla, A. Bandyopadhyay, S. Bose. 2010, Adv. Eng. Mater. (Adv. Biomater.), Vol. 12, p. B V. Balla, S. Bose, N. Davies, A. Bandyopadhya. 7, 2010, JOM, Vol. 62, p A. Fukunishi. 1995, Nippon Seikeigeka Gakkai Zasshi, Vol. 69, p H. Weinans, R. Huiskes, H. Grootenboer. 1993, J Biomech., Vol. 26, p

49 V. Balla, S. Bose, A. Bandyopadhyay. 2009, J Biomed. Mater. Res. B, Vol. 89B, p V. Balla, W. Xue, S. Bose, A. Bandyopadhyay. 5, 2008, JOM, Vol. 60, p Conf. Proc. on Trends in Welding Research. David, S. [ed.] S.A. David and J.M. Vitek. Gatlinburg, TN : ASM International, Materials Park, Ohio, USA, June 1 5, W. Hofmeister, M. Griffith, M. Ensz, J. Smugeresky. 9, 2001, JOM, Vol. 53, p C. Akihiko, K. Kazushige, N. Naoyuki, M. Satoru. 2007, Acta Mater., Vol. 55, p R. Kenneth, D. Lyle, A. Robert. 2004, J Biomed. Mater. Res. Part B: Appl. Biomater., Vol. 68B, p. 1.

50 A1 APPENDIX 1 IN SITU GROWTH OF TANTALUM NANOTUBES ON TANTALUM 1. INTRODUCTION As discussed in part b tantalum has some important biological activities which make it a desirable metal for implants. Tantalum is more wear resistant material than titanium. However the wear rate of tantalum is still two orders of magnitude larger than that of CoCrMo alloy. One possible method of improving the wear rate and potentially improving bone ingrowth is the formation of nanotubes on the surface by anodization. Converting the surface to tantalum pentoxide may increase the hardness because it is a ceramic. Another potential improvement is the creation of nanopores. The purpose of this study was to investigate processing of Tananotubes on LENS TM fabricated nanotubes. 2. EXPERIMENTAL Pure tantalum samples were prepared from 1mm thick tantalum plate. Lens processed tantalum samples of 10 mm diameter were fabricated using LENS -750 (Optomec Inc. Albuquerque, NM) equipped with a 500W Nd:YAG laser and a double powder feeder system. Samples were fabricated in a glove box containing an argon atmosphere with O 2 content less than 10 ppm to limit oxidation of tantalum during processing. The laser produced 400 watts with a scan rate of 7 mm s -1 and a powder feet rate of 106 g min -1 both sample types were polished with one micron alumina. Tantalum nanotubes were produced by anodization in a simple anodizing cell in which a platinum cathode was used to eliminate potential interaction of the cathode with the electrolyte.

51 A2 The acid solution was selected to dissolve the metal fast enough to anodize the surface in a reasonable amount of time and not too fast so the oxide had time to precipitate on the surface. This solution also needs to be thick enough to allow a ph gradient to develop between the bottom of the nanotube and the top surface of the forming nanotube. The actual acid solution selected was 50 ml concentrated H 2 SO 4, 2 ml H 2 O and 0.5 ml 49% HF. This solution may not be ideal due to how quickly the oxide forms on the surface. The anodized surface was rinsed with concentrated H 2 SO 4 to remove any tantalum in solution on the sample surface. Acid was used because rinsing with water causes a rapid change in ph which resulted in a layer of tantalum precipitating on the surface covering the nanotubes. After the acid rinse a water rinse was used to remove the acid to prevent further dissolution of the surface. Tantalum samples were platinum coated with a plasma sputter coater. This platinum coating was used to eliminate charging at high magnifications. Examination of samples was conducted utilizing a field emission Scanning Electron Microscope (SEM). Nanotube measurements were made with ImageJ software. 3. RESULTS AND DISCUSSION Initial attempts at forming nanotubes on LENS TM processed tantalum failed. The anodization proceeded quickly for all samples. The amperage dropped from as high as 2 amps to below 0.05 amps in less than a minute, usually around 5 seconds, and dropped to 0.01 amps in less than 8 minutes for all samples. After a few failed attempts the decision was made to eliminate some of the variables in order discover the source of the problem. There is no literature at this time on the formation of nanotubes on LENS TM processed materials. Since LENS TM

52 A3 Figure A1 Tantalum pentoxide coated surface Figure A2 Potential nanotubes on tantalum pentoxide surface

53 A4 Figure A3 Top view of nanotube covered surface Figure A4 Tantalum nanotube surfaces dabbed dry

54 A5 processed tantalum using the described methods has porosity between 27 % and 55% and other inhomogeneities that make nanotube formation difficult given that nanotubes require a very ordered and uniform structure. Plate tantalum was used to determine the source of the problem. Early attempts using tantalum plate also seemed to fail as shown in Figure A1. An additional attempt, the image in Figure A2, showed a structure that might be nanotubes but since neither end seemed to be open it might have been nanorods. Notice also that the surface in Figure A3 is similar to the surfaces in Figure A2 though the significance of this was not immediately realized. The images shown in Figure A1 through A3 along with the understanding that the Figure A5: Tantalum nanotube surface with a water rinse

55 A6 amperage started high and dropped to zero seem to indicate that the surfaces were covered in nanotubes but tantalum pentoxide was precipitating on the surface. In the next attempt a paper towel was used to dry the surface immediately after anodization. As shown in Figure A4 this resulted in a highly damaged surface. The next attempt, shown in Figure A5, shows the tantalum pentoxide forming on the nanotube surface after a water rinse. It did form nanotubes first as shown in a higher magnification image in Figure A6. Rinsing with acid and then with water gave the image in Figure A7. It is significant to note that the surface on top of the nanotube plateau and the lower surface are a similar pattern of holes indicating the lower surface is covered in nanotubes. It should also be noted that the top surface is not the top of the nanotubes but a thin surface coating with openings. A close up of this image, shown in Figure A8, shows that there Figure A6: Water rinsed tantalum pentoxide surface showing nanotube openings below the precipitate

56 A7 Figure A7: Acid then water rinsed tantalum pentoxide nanotubes Figure A8: Close up of acid then water rinsed tantalum pentoxide nanotubes

57 A8 140 Tantalum Nanotube Diameter Nanotube Length nm nm Outside Diameter Inside Diameter V p< n= V p<.0001 n=10 Figure A9: Tantalum pentoxide nanotube dimensions are openings on the end making them nanotubes and not nanorods. Images of tantalum pentoxide nanotubes were taken and their dimensions measured. Between 20 and 40 volts the dimensions, as shown in Figure A9 (raw data in ) varied linearly with voltage. These results are consistent with a previous study (A1) which, measured the outside diameter only, and showed the linear relationship continued up to 90 volts. The outside diameters were consistent in both studies. Despite changes in diameter wall thickness remained unchanged at 25 nm. 4. SUMMARY At this time this study is not complete however results up to this point indicate a few options for proceeding through these challenges. It was determined that nanotubes may have been forming on tantalum plate but precipitation of tantalum pentoxide would have covered the

58 A9 surface of any nanotubes that were formed. The first option would be to anodize LENS TM processed tantalum with a proper post anodization cleaning process. Even with this cleaning procedure a surface coating still formed. Therefore another option would be to continue to work on eliminating the tantalum pentoxide coating. One way to eliminate this coating would be to stabilize the precursors of tantalum pentoxide in solution because it precipitates too easily. If the surface coating can be eliminated nanotube dimensions should be optimized for bone tissue interactions. A previous study indicated that bone ingrowth occurs at 200 nm. This correlates to an anodization voltage of 80 volts. The final option would be to drop the pursuit of tantalum nanotubes and just anodize the LENS TM processed material to form the oxide. While this would not exploit the potential advantage of cell integration inside the nanotubes, tantalum pentoxide is probably more wear resistant than tantalum. Even without the added benefit of potential cell ingrowth tantalum pentoxide is probably more bioactive than tantalum and therefore may be a route worth pursuing. REFERENCES A1. J. Barton, C. Stender, P. Lia and T. Odom. June 18, 2009, Journal of Materials Chemistry published as advance article on web DOI: /b904964a.

59 A10 APPENDIX 2 EXPERIMENTAL PROTOCOL FOR METAL ION RELEASE USING ATOMIC ABSORPTION SPECTROSCOPY There are a number of concerns when using atomic absorption spectroscopy (AAS) in either flame or furnace mode including which mode to use. There are a few issues to consider, the first is the amount of sample needed. In flame mode the auto sampler uses a lot more sample though the amount is highly dependent on the settings. Each sample should be measured at least twice with a third measurement taken if they differ by more than 5%. This is not uncommon since there are a number of factors that lead to variations in measurements. Rinsing and sampling times needs to be set long enough to allow the signal to stabilize after switching between the rinse, sample and back to rinse. The previous sample needs time to completely flush from the system. The default parameters will usually work for each sample however they are designed to handle a wide range of samples and are usually take longer than necessary. It is not worth the time to optimize the temperature and time parameters if only a few runs will be performed on a given type of solution. Solution is withdrawn continuously during the sampling process. Typical parameters will take a minute or more and use a few milliliters of sample. Furnace mode does not continuously draw sampling and only uses a few microliters per test however they both take about as long to perform.

60 A11 Another consideration is contamination. In flame mode most of the sample tube is immersed in solution and remains there until the sampling stops. Then the sample tube is immersed in a container of rinse water while the tube is flushed with that water. The main problem is if there is much analyte in the sample it will contaminate the rinse container and result in baseline creep. In order to ensure this is not an issue a blank sample should be analyzed every ten samples and any readings less than blanks must be reported as below detection limits. In addition if the sample has significant salt or other content it will readily contaminate and plug the atomizer tube so all furnace mode samples should be clean water or highly diluted samples. Another option is to use furnace mode. In furnace mode a small tip is used instead of a tube. This tip never touches the rinse water container instead rinse water is pumped out of the tip to clean it so no contamination occurs. Also in furnace mode dilutions of the sample and standards can be automated. The result is only one standard needs to be made and the labor to do manual dilutions if the samples are out of range can be eliminated. The last consideration is the detection limit required. Generally flame mode is used for ppm determinations and furnace mode is used down to the ppb level though with dilutions it can be used for higher concentrations. With automatic dilutions it is easier to use flame mode for most samples. Once the appropriate mode is selected the next consideration is quality control samples. EPA guidelines and other standards require a five point calibration curve for AAS. The range is determined by best guess than trial and error but it should always span the range of

61 A12 values sampled. This is critical since the values may exceed the linear range and result in reporting errors. The first point is always a blank and the next highest standard is the method detection limit (MDL). Any values below the MDL or the blank if it is higher is reported as below detection limits not the value reported by the machine. The MDL should be repeated ever twenty samples. To do otherwise is to imply lower detection limits than the data indicates. A matrix spike (MS) should also be prepared. A MS is the same fluid as the sample and prepared in the same way and then spiked with the same concentrations as the standard. This demonstrates if there is any interference, either positive or negative, in the signal of the desired analyte by anything else in solution. If there is interference matrix modifiers are used to eliminate them. These modifiers are listed with time and temperature profiles for each element. All samples should be bracketed by blanks and MDLs. In this study flame mode was used and no dilutions were used. This is not recommended since contamination of the rinse water occurred and the atomizer tube plugged repeatedly. The rinse water contamination was indicated because the blank concentrations which were used as a baseline increased with time. Deviation of the MS values from the calibration curve indicated a matrix modifier was needed. The interfering element was probably sodium as evidenced by the yellow flame color when analyzing anything with the SBF matrix. SBF is made with sodium and is one of the interfering elements for cobalt. SrCl matrix modifier was added and was effective in eliminating the interference of sodium.

62 A13 APPENDIX 3 WEAR TESTING OF METALLIC BIOMATERIALS A Nanovea tribometer was used with a linear sliding attachment for reciprocating wear. The sliding arm was bolted to the 10 mm sliding position on the disc attached to the motor. The sample tray is bolted to the linear sliding mechanism. Pins were glued to a sample holder which attaches to the vertical bar on the sliding arm. A different vertical bar is designed to hold a wear ball in place. The base sample was bolted inside the sample tray with sample holder. The tribometer must be on a solid surface and adjusted to level. The screws holding the vertical bar in place were loosened and the sliding arm was adjusted up and down until it was level with the sample in the vertical bar touching the sample in the tray. All the weights were removed from the vertical bar but the collar and rubber vibration gaskets were put back on the vertical bar. The sample tray was filled with SBF solution. At this point the sliding arm needed to be balanced. This was accomplished by adjusting the counter weights near the fulcrum were adjusted until the top and bottom samples were touching but a slight breath would cause the bar to rise. This indicated the sliding arm was balanced. The desired weight was added to the vertical bar to provide the force indicated by the procedure. The splash guard was attached and the sample heater was set to 37 C. The sample name and parameters were entered into the computer. The sliding speed was set to 1000 mm/min for part A and 1100 mm/min for part B. At this point the machine was started and ran until completion. Once the test was over each 10mm wear track was measured ever 0.6 mm until 15 measurements were obtained. The widths were averaged and the standard deviation determined. The first sample had a wear track width of 0.75 mm and a standard deviation of mm. The

63 A14 different measurements are shown in figure A10. The cross sectional area was calculated by using half the track width a (0.385 mm) and the radius of curvature of the pin (40 mm) in equation 1. k=0.5*r 2 [2*arcsine(a/r)-sin(2*arcsin(a/r))] (1) The cross sectional area k was 9.70 x 10-4 mm 2. To get the wear volume the cross sectional area is multiplied by the 10 mm track length to give a wear volume of 9.70 x 10-3 mm 3 Equation 2. Wear volume = k*path length (2) Then this wear volume is normalized by dividing by the force in newtons ( 3 N) and the sliding distance in meters ( 3 km) as shown in equation 3. This normalized wear rate was 1.08 x Normalized wear rate = (k*path length)/(force *sliding distance) (3) Figure A10 Cross section of wear track