HYDROXYAPATITE DEGRADATION AND BIOCOMPATIBILITY

DISSERTATION

Presented in Partial Fulfillment of the Requirements for

the Degree of Philosophy in the Graduate

School of The Ohio State University

By

Haibo Wang, B.S.

* * * * *

The Ohio State University 2004

Dissertation Committee: Approved by Professor John Lannutti, Adviser

Professor Henk Verweij ______Adviser Professor Derek Hansford Materials Science and Engineering

ABSTRACT

Hydroxyapatite (HA) is widely used as a bioactive ceramics since it forms a chemical bonding to bone. The disadvantage of this material is its poor mechanical properties. HA can be degraded in body, which is the reason for its bioactivity, but too fast degradation rate could cause negative effects, such as macrophage present, particle generation, and even implant clinical failure. HA degradation rate will be greatly changed under many conditions: purity, HA form (i.e. bulk form, porous form, coating, or

HA/polymer composites), microstructure, implant site, body conditions, etc. Although much work has been done in HA properties and application areas, the HA degradation behavior and mechanism under these different conditions are still not clear.

In this research, three aspects of HA degradation have been studied: 1) Two very common impurities—Tri-Calcium Phosphate (TCP) and Calcium Oxide and their influences on HA degradation in vitro and in vivo, 2) influence of HA/polymer composite form on HA degradation, 3) HA material particle generation and related mechanism.

From the in vitro and in vivo tests on bulk HA disks with various Ca/P ratios, HA degradation can clearly be found. The degradation level is different in different Ca/P ratio samples as well as in different test environments. In same test environment, non- stoichiometric HA samples have higher degradation rate than stoichiometric HA.

HA/PMMA composite design successfully intensifies HA degradation both in vitro and ii in vivo. Grain boundary damage can be found on in vivo test samples, which has not been clearly seen on bulk HA degraded surface. HA particle generation is found in in vitro and in vivo HA/PMMA composite surface and in vivo bulk HA surface. Sintering temperature and time does affect HA grain size, and this affect HA degradation rate. Intergranular fracture is found in a several micron zone close to the Ca/P ratio 1.62 and 1.67 sample degraded surfaces. At Ca/P ratio greater than 1.667, after HA degradation in water, solution pH increases because of CaO presence.

iii

Dedicated to my wife

iv ACKNOWLEDGMENTS

I want to thank my adviser, Dr. John J. Lannutti, for his intellectual guidance, encouragement and incredible patience which made this thesis possible.

I thank Dr. David Anderson from Veterinary Clinical Sciences, who give us the opportunity to do in vivo tests. I also thank Dr. Amr Moursi and Phillip Winnard for their help on osteoblastic cell culture and the in vitro test.

Special thanks go to Dr. Henk Verweij and Dr. Derek Hansford for serving on my defense committee and providing valuable suggestions to the completion of this work. I am especially grateful to Dr. Jong-Kook Lee-the visiting professor from Chosun

Univerisity, Kwangju, Korea, who joined this research and during the one year on his staying with us, he gave us important contribution on the literature review. I wish to thank all the members of our research group for providing a joyful working site, especially Kathy Lu, Nan Guo, Shiling Ruan, J. Dawson White, Wenxia Li, Taryn

Sproule, Jin Nam, Heather Power, Zhijun Zhao, Alex Tsai, Andrew Kohm, and Kathy

Elias.

I thank my grandparents, my mother for their constant love and support to me.

Finally, my loving appreciation goes to my wife Ying, who encouraged me and unconditionally supported me.

v VITA

1972 ……………………………………… Wanshan, GuiZhou Province, China.

1995 ……………………………………… B.S. Mechanical Engineering, Tsinghua Univ. Beijing, China.

1998 ...... Research Assistant, National Ceramic Lab, Beijing, China.

1999 - present ...... Graduate Teaching and Research Associate, The Ohio State University

PUBLICATIONS

Research Publication

1. Wang, Haibo, Jong-Kook Lee, Amr Moursi, John J. Lannutti. “Microstructural disassembly of calcium phosphates.” Journal of Biomedical Materials Research, Vol. 68A, p61-70, 2004.

2. Wang, Haibo, Jong-Kook Lee, Amr Moursi, David Anderson, Phillip Winnard, Heather Powell, John Lannutti. “Ca/P ratio effects on the degradation of hydroxyapatite in vitro.” Journal of Biomedical Materials Research, Vol. 67A, p599 - 608, 2003.

3. Wang, Haibo, Jiemo Tian. “Improvement of ZrO2 Ceramic Performance by Coating Small Amount CeO2 On 3mol%Y2O3－ZrO2 Powder.” 1st International High-Performance Ceramic Conference in Beijing, China. 1998.

FIELDS OF STUDY

Major Field: Materials Science and Engineering

vi TABLE OF CONTENTS

Page Abstract ………………………………………………………………….ii Dedication ...... iv Acknowledgments ...... v Vita ...... vi List of Tables ...... ix List of Figures ...... x

Chapters:

1. Introduction ...... 1

1.1 The advantages and disadvantages of HA material...... 2 1.2 Preparation of stoichiometric HA powders ...... 5 1.2.1 Syntheses based on theoretical compositions………………….5 1.2.2 Equilibrium syntheses in solution...... 7 1.2.3 Miscellaneous methods...... 8 1.3 HA dissolution behavior and mechanism ...... 8 1.3.1 non-biological HA dissolving behavior in water and acid...... 9 1.3.2 Biology-related HA dissolving ...... 16 1.4 The unresolved issues and importance of study on HA degradation..... 21 1.5 Practical issues which influence HA degradation ...... 23 1.5.1 HA impurity ...... 23 1.5.2 HA form of HA composites...... 30 1.5.3 HA material particle generation...... 36

2. Ca/P ratio effects on the degradation of hydroxyapatite in vitro ...... 48

2.1 Introduction ...... 48 2.2 Experiment design...... 50 2.3 Results ...... 52 2.4 Discussion ...... 54 2.4.1 Weight gain/surface precipitation...... 57 2.4.2 Potential in vivo significance ...... 58 2.5 Conclusions ...... 59

3 In vivo degradation results and microstructural disassembly...... 86

vii 3.1 Introduction ...... 86 3.2 Experiment...... 88 3.2.1 Preparation of pure and biphasic HA...... 88 3.2.2 In vitro sample preparation ...... 90 3.2.3 Exposure to primary osteoblast culture ...... 91 3.2.4 Subcutaneous implantation ...... 93 3.2.5 Bone implantation ...... 93 3.2.6 Specimen Harvest...... 94 3.3 Results ...... 95 3.3.1 In vivo study under skin and in bone on bulk HA disks...... 95 3.3.2 In vivo bone apposition and subcutaneous tests on HA/PMMA composite...... 96 3.3.3 Grain pullout in vitro...... 97 3.3.4 Grain pullout in vivo...... 98 3.3.5 Pure HA degradation in vivo...... 100 3.4 Discussion...... 100 3.5 Conclusions ...... 104

4 Further study on hydroxyapatite degradation ...... 138

4.1 Introduction ...... 138 4.2 Experiment...... 140 4.2.1 Ca/P 1.55~1.78 non-stoichiometric HA sintering ...... 140 4.2.2 Degradation Test ...... 141 4.2.3 HA sintering at different temperatures and times ...... 142 4.3 Results ...... 143 4.3.1 XRD result for Ca/P ratio 1.55~1.78 sintered samples ...... 143 4.3.2 Solution influence on HA degradation ...... 144 4.3.3 HA microstructure vs. degradation rate in water ...... 146 4.4 Discussion...... 147 4.5 Conclusions ...... 148

Bibliography……………………………………………………………………………166

viii LIST OF TABLES

Table Page

1.1 List of chemical reactions in Ca-P-H2O solution system. Each of the reaction is plotted in Figure 1.3 and is marked by the respective letter...... 44

2.1 Phosphate buffered saline solution composition; 0.03wt% sodium azide was also present to prevent bacterial growth ...... 52

ix LIST OF FIGURES

Figure Page

1.1 Hierarchical levels of a human long femur structural organization ...... 41

1.2 Solubility isotherms for various phases in the system CaO-P2O5-H2O at 25oC...... 42

1.3 Equilibrium diagram for hydroxyapatite. Solid lines are the pH-invariant lines on the solubility surface. Broken lines separate the stability domain of various solution species...... 43

1.4 Schematic picture of the calcium ions concentration profile through a boundary layer including the calcium adsorbed layer, Nernst layer, and bulk solution ...... 45

1.5 Phase equilibrium diagram of calcium phosphates in a water atmosphere. In shaded area HA-containing material will be obtained after sintering...... 46

1.6 An example of impure commercial HA products. HA powder is from GFS chemicals. These three XRD results are for HA from GFS, β-TCP and pure HA...... 47

2.1 XRD pattern of the hydroxyapatite used in this study following sintering at 1200°C for 10 h. (.) Hydroxyapatite. No trace of the 100% 2θ peaks of either TCP or CaO are present ...... 61

2.2 Volumetric percentages of TCP and CaO in HA versus the targeted Ca/P ratios. In terms of the volume of second phase added/present, considerably more TCP is required to achieve the same 5% variance away from the stochiometric atomic ratio ...... 62

2.3 Weight versus time and pH for all three Ca/P ratios. (a) pH 6.8, (b) pH 7.0, (c) pH 7.4. After 15 days exposure, consistent weight gains were observed for the 1.72 samples under all conditions ...... 65

2.4 Representative SEM micrographs of polished surfaces before testing in pH-modified PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72...... 68

x 2.5 Representative SEM micrographs of polished surfaces after 3 days exposure to pH 6.8 PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72...... 71

2.6 Representative SEM micrographs of polished surfaces after 3 days exposure to pH 7.4 PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72. In spite of an extensive search, no precipitates were observed on the surface of any 3 day Ca/P ratio 1.72 sample...... 74

2.7 Representative SEM micrographs of polished surfaces after 15 days exposure to pH 6.8. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72. The arrows in (c) point to pits displaying pronounced angularity. The pores in (b) are considerably deeper than those seen in Figure 5(b)...... 77

2.8 SEM micrographs of the precipitation on the Ca/P ratio 1.72 sample surfaces after 15 days of in vitro exposure. (a) following exposure to a pH of 6.8, a precipitate having a platy microstructure having greater connectivity forms; (b) following exposure to pH 7.0, a precipitate again showing a platy microstructure. (c),(d) following exposure to a pH of 7.4 a two-phase precipitate forms. In none of these cases did the precipitate cover the entire sample surface; the images in Figure 2.7 were obtained from the uncovered areas ...... 81

2.9 SEM micrographs of the cross-section of HA with Ca/P ratio of (a) 1.62, (b) 1.67, (c) 1.72. The stoichiometric sample displays primarily transgranular fracture; grain size estimates were based on those areas in which fracture was intergranular (see arrow). In (c) the CaO inclusions are visible as Ca-rich (via EDS) grain boundary precipitates having a distinctive surface morphology...... 84

2.10 High magnification SEM of the interior of a large pore in the surface of the CaO-containing sample after only 3 days exposure to pH 7.4 solution. Although abundant polishing debris partially obscure some features, this pore interior displays abundant evidence of smaller pores along the grain boundaries indicating that, once exposed, the CaO inclusions dissolve leaving distinctive submicron pores...... 85

3.1 Control microstructures of as-polished surfaces via SEM. The Ca/P ratios were (a) 1.62 and (b) 1.67. Polishing debris cover the surface and make direct microstructural observations difficult...... 106

xi 3.2 PMMA disk containing an embedded HA bar. The upper surface was carefully polished to expose a small (20 µm wide) portion of the bar. The HA occupies only 0.12% of the total specimen surface ...... 107

3.3 HA with Ca/P ratio 1.67, 1.62 and 1.72 degradation for day 3 under skin. All photos were in same magnification. Large holes and slots started to show on all Ca/P 1.72 (HA + 1.6v% CaO) samples (c). Not much soft tissue can be found on top of all day 3 samples ...... 110

3.4 HA with Ca/P ratio 1.67, 1.62 and 1.72 degradation for day 15 in bone. Holes on pure HA surface are relative small and abundant (3.3a,3.4a). A few fairly large (5-10 µm) pores are surrounded by large, undamaged areas on Ca/P 1.62 (HA+27v%TCP) sample surface (b). Slots and pores are bigger than day 3 samples on Ca/P 1.72 sample surface (c). All day 15 samples were covered by soft tissues after sample harvest. These tissues were carefully removed using ultrasonic in ethanol before SEM test...... 113

3.5 Oxytetracycline-labeled histological sections, new bone start to form on day 15 mineralized samples. New bone formation line (white shining line) has already been formed on pure HA surface (a). The HA+TCP implants also showed new bone formation (b), but bone-to-HA direct bonding was present only in some areas. On the HA+CaO implants not only new bone was formed, but also a thick collagen fiber layer was found (c). The two particles being evolved by the HA+1.6v%CaO composition are each marked with a ‘*’...... 116

3.6 Comparison of degradation on bulk HA surface and on HA/PMMA composite surface after 15 days. Bulk HA surface in bone apposition (a). HA bar surface of HA-PMMA in bone apposition (b). HA bar surface of HA-PMMA under skin (c)...... 119

3.7 Grain pull out after 3 days in vitro. In (a), grains are extracted from the polished face of the bar. The image in (b) shows that grain extraction also occurs when the ECM comes into contact with the rougher, unpolished sides of the bar. Significantly, the left- and right-hand sides of the image in (c) show that without contact with the ECM (apparently due to bridging of the bar-PMMA gap), no grain extraction is visible ...... 122

3.8 Longer term degradation of HA in vitro. (a) 9 and (b) 15 days. In (a), extraction by the ECM is still visible but the underlying grains have xii undergone considerable degradation. In (b), a higher magnification shows that after 15 days the average diameter of the grains remaining on the surface continues to decrease ...... 124

3.9 Surface morphology following 15 days exposure to media alone. (a) shows the bar after exposure and mechanical extraction from the PMMA. Some surface discoloration is present probably due to contamination by media salts. (b) shows a higher magnification of the surface in which polishing lines (fine vertical scratches) are still visible...... 126

3.10 Micrographs of biphasic HA with a Ca/P ratio of 1.62 after 15 days subcutaneous implantation. At low magnifications, the image in (a) shows the ‘peeling’ of the ECM layer on the left side of the micrograph. A higher magnification image (b) shows demarcation between peeled away and still- adherent ECM. The ECM (the right-hand side of the image) precisely replicates the as-polished surface morphology. Without clear evidence of ECM-based extraction from the surface, we could easily mistake this for the HA surface. A small portion of the hundreds of grains studding the peeled-off layer in (a) is shown in (c). Extraction involving groups of (as opposed to just single) grains can clearly occur. In (d) the layer of ECM is ~1-2 µm thick; HA grains can be seen adhering to the opposite surface...... 130

3.11 Transgranular cracking (visible in the collection of grains at the center of this image) observed infrequently...... 131

3.12 High magnification of superficial degradation of biphasic HA after 15d subcutaneous testing. The step-like cracks on the grain surfaces suggest successive grain boundary degradation...... 133

3.13 Resorption lacunae (a) present in phase pure HA after 15d subcutaneous implantation. Grain extraction was not observed for this material in vivo. The bottom of the lacunae (b) shows some remnants of the original microstructure that have undergone relatively severe degradation. While grain dissolution appears to be governed by the individual characteristics of each grain, it is clear than the generation of “loose” particles at the bottom of the lacunae is far less likely than dissolution ...... 135

3.14 Cross-section of the 15d biphasic HA-tissue interface. The left-hand portion of each image shows a region displaying strictly transgranular fracture. The right-hand portion (near the surface) shows intergranular fracture. In xiii some portions of the fracture surface in (a) the depth of the intergranular zone is highly variable; the region of transgranular fracture marked with an ‘*’ is almost at the surface while intergranular fracture in the same image extends inwards to a depth of 20 microns. In (b), a higher magnification shows that the interface between the intergranular and transgranular regions is very distinct. No mixed-mode fracture is observed ...... 137

4.1 XRD patterns of non-stoichiometric HA. (a) Ca/P<1.67, (b) Ca/P>1.67 ...... 150

4.2 Linear relationship between the impurity peak height and Ca/P ratios. (a) Ca/P<1.67, (b) Ca/P>1.67 ...... 152

4.3 Surface etching after 15 days in α-MEM. Images taken by optical microscrope (50X). Black areas are etching pits. (a) Pure HA surface, (b) CaO-HA surface ...... 154

4.4 Pits area percentage gained from computer analysis. 5ml DI water, PBS, and α-MEM were used here. (a) HA surface, (b) CaO-HA surface...... 156

4.5 pH test results of (a) pH 7.4 water, (b) pH 7.4 PBS and (c) α-MEM...... 159

4.6 Water penetration depth on CaO-HA disk surface calculated from Figure 4.5a. Water penetration is fast at the first three days...... 160

4.7 pH change in 5 ml DI water. The vertical broken line represents Ca/P 1.667... 161

4.8 The sample surfaces after thermal etching show the grain sizes. (a) 1150oC, 90 mins, (b) 1200oC, 90 mins, (c) 1250oC, 50 hrs...... 164

4.9 Degradation results in DI water after three different sintering procedures...... 165

xiv

CHAPTER 1

INTRODUCTION

In the past 30 years, hydroxyapatite (HA) (Ca5(PO4)3(OH)) ceramic implants have attracted much attention as an alternative substance for autogenous free bone grafting [1-

6]. It is the most prominent bioactive ceramics and is widely used and investigated.

Applications include coatings of orthopedic and dental implants, alveolar ridge augmentation, maxillofacial surgery, otolaryngology, and scaffolds for bone growth and as powders in total hip and knee surgery [6-10].

Synthetic HA is classified as polycrystalline ceramics since its material structure is derived from individual crystals that have been fused together by a high temperature sintering. The most popular processes to fabricate dense HA ceramics use uniaxial compression (for preparing small specimens [11]) through which HA powder is pressed in a metal die to produce a compact, which is sintered late at high temperatures in the range of from 1100oC to 1300 oC to gain densification through a typical solid sintering kinetics. Because of the pressure asymmetry in uniaxial copression, isostatic compaction is used to generate denser and more uniform green compact via hydraulic systems. Insead of pressure from one direction in uniaxial pressing, pressure is applied from all directions 1 to a flexible mold, which is filled with ceramic powder [12, 13]. Some researchers even use hot-pressing to gain even better density and fine. The powder is pressed and heated to a relative low temperature simultaneously to generate the high density and fine grains in one step [14, 15]. Slipcasting is a procedure through which HA powder is casted into a mold in wet state. After drying in mold, the HA compact is sintered and densified [12,

16].

1.1 The advantages and disadvantages of HA material

Many biocompatibility studies prove that HA has a very similar chemical composition as the inorganic part of human hard tissue, such as bone and teeth [17,18].

The most important advantage of HA being a bioactive material is that bone will form a direct chemical bonding to HA implant without forming a collagen interface layer which is usually found in many other bioinert materials after implantation [19-23]. From many research papers it can be concluded that the reactions happening at the bone-HA interface follow the following procedure: 1) After implantation, calcium phosphate solid-solution equilibria are established by calcium and phosphate ions which are released from implant and surrounding bone (This means slightly dissolution of HA or bioglass is very important for the so-called bioactivity of these bioactive materials). 2) This process results in calcium and phosphate ions supersaturation in the surrounding body fluid, and then carbonate apatite crystallites epitaxially reprecipitate on the surface of the HA [21].

3) This modified surface are known to accommodate protein adsorption and cell adhesion 2 more rapidly, in particular, cells (osteoblasts) associate with bone bonding [22]. 4) A cellular bone matrix originated from osteoblasts appears on the HA surface, producing a narrow amorphous electron-dense band only 3-5 µm wide. Collagen bundles are seen between this area and the cells. Small bone mineral crystals can be identified in this amorphous area. 5) As the site matures, the bonding zone shrinks to a depth of only 0.05-

0.2 µm. Also normal calcification processes are observed immediately adjacent to the implants, as evidenced by Ca/P ratios rising from 1.50 to 1.62 (by electron microprobe) after six months as well as calcium and phosphorous concentration increasing [23]. At six months, mineralization within the implant sites is comparable to the surrounding bone.

TEM image analysis of dense HA bone interfaces show almost perfect epitaxial alignment of some growing bone crystallites with the apatite crystals in HA implant.

However, the process that this amorphous layer gradually matures and becomes bone structure remains unknown [24]. Due to this chemical bonding interface, the bonding strength of HA and bone is much higher than other materials, such as Al2O3, ZrO2,

Titanium alloy, etc. Thereby the relative micro-movement between the implant and bone is dramatically reduced by this direct bonding, and no fibrous tissue capsule can be found between the implant and bone. This is important for the patient’s recovery in the early period after implantation.

The main disadvantage of HA is the poor mechanical property. Like most ceramics, the low toughness and impact resistance limit the clinical application of these materials. HA can’t be used as a bulk material sustaining tension or impact. The tensile strength and compressive strength of the synthetic dense HA are 10-28x103psi and 30-

130x103psi. The tensile strength and compressive strength of cortical bone are 10x103psi 3 and 20 x103psi [16]. There are two problems with HA mechanical properties which affect

HA clinical use: 1) HA is stiffer than bone. There is a modulus difference between cortical bone and dense HA. The Young’s Modulus of cortical bone is 2x106psi, while 5-

15x106psi for dense HA [16]. 2) Bulk HA implants have low reliability under tensile loads [31]; cracks inside HA could cause catastrophic failure under use. Biologically, the

HA in bone has very unique molecular structure, microstructure, and macrostructure

(Figure. 1.1). These structures give bone a unique ceramic/collagen fiber composite structure, which successfully prevent bone from breaking under mechanical load. Among them the microstructure plays an important role in this protection. The osteons are composed of concentric lamellae. Each lamella is composed of collagen fibers helixes.

The apatite crystals can be found both inter-and intra-fibrillarly within the collagen layer.

In this way, the bone elastic modulus drops but the elasticity increases. Certainly this microstructure is an ideal model for all HA composite, while the synthesis methods already developed are not capable of simulating this structure. The most advanced composite can only simulate the bone macrostructure—the comby HA structure with about millimeter size holes filled with polymer. If HA were used as a bulk implant, the elastic modulus mismatch of bone and implant materials would result in disproportionate load sharing. As a result of bone remodeling originating from this abnormal localized bone stress distribution around the implant may cause bone to undergo stress protection atrophy. This results in bone mass loss and osteoporosis. At normal condition, since living bone keeps remodeling, any crack forming in bone will be repaired by degradation of old osteon and regeneration of new osteon and the bone function can be recovered

[24]. However, bulk synthesize HA can’t remodel by itself, and cracks and flaws can 4 easily be developed during bulk ceramic synthesis, which will lead material to brittle failure.

1.2 Preparation of stoichiometric HA powders

The large number of different synthetic routes and published preparations demonstrate the difficulty of making stoichiometric HA powder. The variability that is found in such products has been reviewed by Young R.A et al [25]. Basically, three kinds of HA synthesis methods have been proposed [29].

1.2.1 Syntheses based on theoretical compositions of HA

This method usually is used for high temperature reactions with solid state precipitation. Raw materials must have the Ca/P=1.67. The starting calcium compound material could be CaCO3, Ca(OH)2 or CaO etc. Gas pretection is required to prevent absorption of CO2, which is in air and very easily incorporated into apatites. The carbonated substitution can occur in two different ways [29]. In type ‘A’ carbonate

2- - In type .”ڤ“ apatite, one CO3 has replaced two OH and created one channel vacancy

2- 3- ‘B’ carbonate apatite, CO3 replaces PO4 , and cation vacancies are created in order to maintain charge balance. Usually high purity H3PO4 solution is titrated against NaOH to conform its concentration before it is used [26]. Sometimes, CaPO4 (DCPA), CaPO4 • 5 2H2O (DCPD) or Ca2P2O7 can be used instead. For quantitative reactions in solution, the reactants can be H3PO4 and Ca(OH)2 [28], or salts of these with Ca and P ions that are unlikely to be incorporated in the apatite lattice. A typical reaction follows:

3H3PO4 + 5Ca(OH)2 ! Ca5(PO4)3(OH) + 9H2O (1)

Nitrates and ammonium salts are used instead of chlorides and sodium salts [27]. The pH must be controlled to above 10 by adding NH4OH or NH3 gas, not NaOH or KOH. The advantage of using ammonium and nitrate salts is that NH4NO3 can readily be decomposed to gas from the separated precipitate on heating. If precipitation has been produced at a high pH, there should be likely no acid phosphate in the HA powder, charge balance dictates that the solid also has the correct proportion of OH- ions. While if precipitation is made near pH=7, in principle, there could be acid phosphate present follows the equation (2) in a direction from right to left:

- 2- 3- OH + HPO4 ↔ H2O + PO4 (2)

2- If the solid is now heated, there is the possibility of transformation of HPO4 to pyrophosphate intermediate [32] (Equation 3). At even higher temperature, calcium pyrophosphate could react with HA and forms TCP [33] (Equation 4), which is an important impurity in HA:

2- ⎯⎯→320-⎯340⎯o C 4− + 2 HPO4 P2O7 H 2O (3)

2 Ca5(PO4)3(OH) + Ca2P2O7 ! 4 Ca3(PO4)2 + H2O (4)

Usually, by this method, very tiny submicron HA crystals are formed with considerable surface area in the order of 10-150 m2.g-1.

It is important to ensure complete reaction to a single phase. The choice of starting materials is not important. Several starting materials could be used including 6 CaCO3, CaO, Ca(OH)2, DCPD, DCPA and Ca2P2O7. A sample reaction uses mole ratio

3:4 of Ca2P2O7 and CaCO3 [34]:

. 3 Ca2P2O7 + 4 CaCO3 ! 10 CaO 3P2O5 + 4 CO2 (5)

The raw materials must be ground and pelleted to ensure completion of the reaction.

Water vapor is supplied at high temperature to ensure reaction product in the form HA finally.

1.2.2 Equilibrium syntheses in solution

From Figure 1.2, stoichiometric HA is assumed to be the most stable phase in solution at a pH above the HA/DCPD singular point (pH 4.3 at 25oC). If this assumption is correct, any CaP left in an aqueous solution for long enough time should transform to stoichiometric HA, but the pH is not allowed to fall below the singular point. In practice, it is usual to hydrolyse DCPD or DCPA. Usually this reaction is carried at boiling temperature solution to speed the reaction, also NH4OH or NH3 gas is added to adjust solution pH [35]. Usually the HA precipitate gained in this way is non-stoichiometric with Ca/P ratio lower than 1.67, an adjustment of pH to higher values and an addition of

Ca2+ ions will raise the Ca/P ratio to 1.67 [36].

• ⎯⎯→H⎯2O • + + 8 CaHPO 4 2H 2O Ca 8H 2 (PO4 )6 5H 2O 2 H 3PO4 11 H 2O (6)

• ⎯⎯→H⎯2O + + 5 Ca 8H 2 (PO4 )6 5H 2O 4Ca10 (PO4 )6 (OH)2 6 H 3PO4 17 H 2O (7)

7

1.2.3 Miscellaneous methods

Some other routes have been reported to successfully synthesize HA, such as:

. o 1. Heating a stoichiometric mixture of MCPM (Ca(H2PO4)2 H2O) and CaCO3 at 1200 C in water steam and N2 [37]. It is important to know that if the system is at equilibrium and the overall Ca/P ratio is not 1.67, a small amount of second phase will be formed.

2. Boiling CaSO4 with strongly alkaline Na3PO4 solution [38].

3. Heating Chlorapatite [25] or Bromapatite [39] in steam at about 800 to 1000oC.

There are still a few more synthesis methods. This part shows that many different ways have already been explored to synthesize stoichiometric HA powder, and the HA powders synthesized by all these different ways have different crystal parameters, crystal size, density, impurity ion contents, and probably subtle Ca/P ratio changes. These differences as well as different sintering conditions will affect material thermo- decomposition, dissolution and mechanical properties.

1.3. HA dissolution behavior and mechanism

1.3.1 non-biological HA dissolving behavior in water and acid

HA is slightly resolvable in water or acidic solution [30,31]. Several mechanisms are proposed to explain the dissolution behavior of HA [43-50]. Lower pH (<4) value 8 will dramatically accelerate this resolving process, and CaHPO4 (monetite) or

. CaHPO4 2H2O (DCPD, brushite) is more stable than HA at pH<4 (Figure 1.2). In a thermodynamical analysis, the equilibrium diagram of the predominant calcium and phosphorous species in solution is shown in Figure 1.3, while Table 1.1 gives all the corresponding chemical reactions which could happen in Figure 1.3. If HA is placed in an aqueous solution in the absence of any soluble salt of calcium or phosphorous, the system will equilibrate along the dotted line (Z-Z’) shown in Figure 1.3. For example, stoichiometric HA is put into certain water solution, after dissolving the pH of solution is

5, if there is no Ca or P containing ions in solution at beginning, then the equilibrium

2+ - solution should drop into the Ca and H2PO4 predominant area. The equilibrium

-2.6 2+ -2.8 - 2- solution should contain 10 M Ca and 10 M H2PO4 ions. The other ions i.e. HPO4 ,

+ - 2+ - CaH2PO4 , CaPO4 are all very limited comparing with Ca and H2PO4 . If the pH value

2- of the equilibrium solution rises to above 7, the predominant anion becomes HPO4 .

The general dissolution reaction of HA in water can be written as an equilibrium

2+ 3- 2- - - of HA, Ca , PO4 , HPO4 , H2PO4 , H3PO4 and OH ions:

(n -1)H O + Ca (PO ) (OH) ⇔ 5Ca 2+ + 3x PO 3- + 3x HPO 2- + 3x H PO - c 2 5 4 3 1 4 2 4 3 2 4 (8) + + - 3x 4H 3PO4 ncOH

(nc: the mole number of protons consumed when 1mol HA is dissolved. xi: (i = 1~4) defines the mole fractions of the different phosphate forms present in the solution. )

Besides the general dissolution reaction, nc and the xi are related by the three phosphate ion equilibria:

- + H3PO4 + H2O <=> H2PO4 + H3O pK1 = 2.19 (9a) 9 - 2- + H2PO4 + H2O <=> HPO4 + H3O pK2 = 7.18 (9b)

2- 3- + HPO4 + H2O <=> PO4 + H3O pK3 = 12.3 (9c)

By combining equation 8~9, one can obtain the relation of the mole of H+ ions consumed and the mole of the Ca2+ ions released from HA:

1 + 3(x + 2x + 3x ) R = 2 3 4 (10) c 5

+ 2+ (Rc: number of H ions consumed when one mole Ca is dissolved, which depends on the pH of the solution.)

The equation 8 fits the experiment results well [40-42], even when the dissolution experiments were realized in the presence of non-stoichiometric calcium and phosphate concentrations initially.

Solubility product Ksp of HA exists. Ksp is the value of KIP at saturation of HA:

2+ 5 3- 3 - Ksp = [Ca ] [PO4 ] [OH ] (11)

o At 25 C in stoichiometric or nonstoichiometric conditions, pKsp of HA is about 57.5 [40,

41, 44, 45]. In the presence of 8x10-2M KCl, the HA solubility decreases with temperature increases.

The HA dissolution rate Rd can be described as the following form:

n Rd = K(CS-C) (12)

CS is solubility, C is the salt concentration in solution, and K is a constant depending on the surface area. Exponent n defines the reaction order. This dissolution rate is related to

HA solubility, which can be drived from equation 11. The dissolution rate can be described as:

= 1/9 1/9 n R d K(K SP - K IP ) (13)

10 Many dissolution models give different reaction orders based on the experimental conditions used. According to classic Nernst’s theory [46,47], a static boundary layer of fluid between the solid surfaces exists immediately adjacent to which the solution is saturated, and the bulk solution. Then n is equal to 1. Here K = DA/ δ, in which D is the diffusion coefficient in the solid adjacent Nernst layer of thickness δ and A the surface area. If initial concentration = 0, solution volume = v and assuming K is constant, the solution concentration will be [51]:

1/9 -kt/v C(t) =KSP (1- e ) (14)

A surface reaction controlled process is the reason for higher n values for which the concentration near the crystal surface is close to the concentration in the bulk solution.

The HA dissolution behavior is a very complicated process. Many parameters affect HA dissolution rate:

1). If the rate is essentially diffusion controlled, a direct proportionality between flux and stirring speed should be observed (usually for HA implants, static body fluid condition should be assumed) [48], indicating that, a significant contribution from the surface reaction to the overall kinetics cannot be excluded from certain experimental conditions

[42].

2). The rate of dissolution of HA at a constant pH is fast initially, and then it declines continuously and falls to a really low level, which is slow than the rate based on a simple first-order law supposing a steady state. All the mathematical expressions of first-, second-, third-order, or simultaneous two parallel first-order equations are able to fit the experimental data over varying periods of time, but none would fit all the data in every case and under all conditions [49]. The order of the reaction n changes with dissolution 11 time and saturation degree from 5 or more to 1 [45]. It reminds us that this decrease in the rate of dissolution with the extent of reaction can be related to the gradual changes of the dissolution surface properties. There should be a rapid adjustment of surface composition occurs initially when crystals are put into the solution [52,53]. This behavior, which is reinforced in the presence of impurities or adsorbed molecules acting as dissolution inhibitors, is important for the understanding of the dissolution mechanism [54].

3). HA dissolution is strongly affected by the H+ concentration, but H+ diffusion from the bulk to the solid interface is not the controlling step. The rate is 100 to 3000 times lower than that it should be, but this shows that there is no interfacial H+ gradient. An (H+)m dependence was observed. The m value is in the range 0.57-0.81 and depending on the experimental conditions and degree of conversion [55-57].

Based on the experimental data, several HA dissolution kinetic models have been proposed. The most important models are two-site model [59-61], polynucleation model

[62-64], and semipermeable interfacial model [65-67]. All the models have shown limitations and drawbacks. These models were shown to deal with different aspects of apatite dissolution and none of them was able to describe the dissolution process in general [68].

Semipermeable interface model: Due to the page limit of this paper, only the semipermeable interface model is discussed in this paper. Since H+ concentration and the

HA solubility gradients were unable fully to describe the dissolution rates [55, 57], researchers assume that the controlling step is the calcium diffusion at the interface.

Many experimental observations show that the amount of calcium ions is changing and is important during all the dissolution process. Calcium ions will be continuously adsorbed 12 onto the HA surface after introduced into the dissolving media, but the amount becomes smaller with the progress of the reaction [45, 55]. It is shown that this accumulation depends on various factors such as the pH, the Ca/P ratio in solution, the sample conditioning, and the undersaturation degree in solution. The amounts of calcium accumulated are 5~50 times higher than the value estimated for a monolayer adsorption.

The result suggests that either multiplayer adsorption or an important increase of the number of adsorption sites during the dissolution. HA interface is saturated and even slightly supersaturated for calcium with respect to the solubility product of HA [65]. This probably is important for HA biocompatibility. Figure 1.4. is a calcium concentration profile, in which the calcium saturated interface restricts the diffusion towards the bulk so that the interfacial diffusion becomes much slower than the solid dissolution reaction.

Since the HA dissolution is always congruent, the dissolution rate can be described by the calcium species. At constant pH, the transport through the interfacial layer is given by the following equation:

RCa = PCa(aS-a0) (15) and

-RH = RCRCa = (5/3)RCRP (16)

(aS: the calcium ion activity at the solid saturated interface, a0: calcium ion activity at the interface between the surface layer and the Nernst layer,

PCa: the mass transfer coefficient characterizing the low permeability of the layer [65],

RH, RCa, and RP: respectively the coupled fluxes of consumed protons, the released calcium and phosphate ions,

Rc: the number of protons consumed when 1 calcium is released (see equation 10). ) 13 The transport rate for calcium through the Nernst layer is simply described by

o RCa = P Ca(a0-a) (17)

o a is the calcium activity in the bulk and P Ca is the classical mass transfer coefficient in the absence of chemical reaction. Since at steady state, both fluxes must be equal, the rate of consumed protons is described by

o P Ca R = - R (a - a) (18) H C 1+ k S

o -1 Here k = P Ca/PCa. k characterizes the reduced permeability of the surface layer with respect to the permeability in the Nernst layer. 1/(1 + k) is the rate reduction factor

(RRF). This relation well describes the experimental results in stoichiometric and nonstoichiometric conditions [45, 65]. k value is about 2000 from pH 3.7 to pH 5, while at pH ≈ 7, k value declines to 300, and the RRF value increases, which illustrates a relative acceleration of the dissolution process. The RRF value depends on sample conditioning and the amount of adsorbed calcium. The RRF shows an exponential type decay up to a quasi-plateau during the dissolution up to equilibrium. At this pesudo- steady state of dissolution, the RRF is found to be only pH dependent but not on the composition of the initial solution [45].

Some researchers propose that the adsorbed calcium ions cover a part of the surface area and limit the accessibility to the surface. The available area for hydrogen attack is much less than the total material surface area, which leads to a much lower dissolution rate in this calcium rich layer [70].

Another important phenomenon in HA dissolution kinetics is the etch pit formation. Some HA surface areas are dissolved faster than the rest area, and then etch

14 pits form. The etch pits formation probably is related to the structural defects, such as dislocations and inclusions [71-73]. Presence of dislocations accelerates dissolution; because they give rise to the strain energy--the cause in crystals favor etch pit formation.

The pits appear at the dislocation outlets, they are 0.1 – 10 µm in size and the size increases when the dissolution progresses. Different acids have no specific influence on etch pit formation and growth [71-73]. The recent investigations on apatite dissolution using atomic force microscopy provided new information on this point. For example, the growth process of a single pit on apatite was followed since its dimension was 50 nm [79,

80].

The HA dissolution is a very complicated process which is still under study. Now many papers are found to focus on many different issues which could greatly enhance or inhibit HA dissolution in water and acid, such as:

1). Physical properties include form (particulate or bulk), porosity (micro- or macroporosity, non- or interconnecting porosity), HA/Polymer form, surface area, and crystallinity (reflecting crystal size, crystal perfection, and grain size as affected by sintering conditions) [43].

2). Ca/P ratio of HA not equal to 1.67, phase impurities resulting from sintering, i.e. tetracalcium phosphate (TetCP), TCP and CaO will increase HA dissolution rate.

2- 3). CO3 ions in HA forms carbonated HA, which exhibits higher dissolution rates than pure HA [81, 54].

- 4). F ions in HA forms fluorapatite Ca10(PO4)6F2 (FAP), with a pKSP of 59.75 [82], is less soluble than HA and dissolves appreciably slower than HA [83]. F- ions not only can be involved during HA synthesis, but also can exchange with OH- in HA in F- containing 15 solution, which is effective for the prevention of dental caries and increase both resistance against demineralization and remineralization rates for the in vitro and in vivo conditions [84-86]. The mechanism is still not clear.

5). HA dissolution rate decreases due to surface modification by surfactant adsorption

[87,88].

6). Some divalent (Sr2+, Sn2+, Zn2+ and Fe2+) and trivalent (Al3+, Cr3+, Fe3+ and La3+) cations are strongly adsorbed onto HA with often important ion exchanges for Ca2+, which can inhibit HA dissolution. While the nature of ion exchange is still not clear [89,

90].

Among these effects, point 1) and 2) will be discussed late in this research.

1.3.2 Biology-related HA dissolving

After implantation, HA ceramic undergo numerous dissolution/precipitation processes induced by biological fluids. This process begins to degrade the ceramic and influence later cellular degradation (i.e. by adsorption of numerous extracellular matrix proteins). Proteins are probably the first species which attach implant right after implantation and implants are continuously interacting with salivary or blood proteins

[31]. However the interaction of HA with biomacromolecules is very complicated and not clear yet. Neutral, acidic, and basic proteins interact significantly with the HA surface. Some proteins are adsorbed onto HA surface and show high adsorption affinity for HA surfaces. These proteins such as salivary proteins are active inhibitors of 16 mineralization [91-93]. More importantly, biochemical and crystalline structural properties of implant material affect the capacity of cells such as monocytes and macrophages to produce proteins and hormones, which will regulate cell activities and thus affect the material dissolving rate [95, 96]. This will affect HA degradation rate.

The cell-related degradation generally includes phagocytotic mechanism

(fibroblasts, osteoblasts, monocytes, and macrophages) and acidic mechanism to reduce pH of microenvironment and resorb the synthetic substrates (osteoclasts) [99]. Since the cells are involved in HA degradation in biological environment, the HA degradation behavior have been greatly changed, and the degradation rate is greatly enhanced.

Mesenchymal cells including fibroblasts, endothelial cells, osteoblasts, and bone- marrow stromal cells are present at the implantation ceramics. These cells participate actively especially in the fibrous encapsulation of implanted ceramics with micromovements. This fibrous encapsulation provents the formation of bone ceramic direct bonding and the degradation processes [98]. The encapsulation is especially important to those bioinert implant since very few ceramic bone direct bonding will form at their interface. Also different mesenchymal cells are present at different implantation sites. Subcutaneous implants are surrounded by more fibroblasts due to high collagen content in these implantation sites. Bone apposition implants are surrounded by more osteoblasts and bone-marrow stromal cells, since the implants are remodeled by these cells. Mesenchymal cells can induce the solubilization of HA material [100-103].

Crystal-cell contact is required to induce HA dissolution [102]. The dissolution process doesn’t appear to be modulated by bone-resorption agent such as parathyroid hormone.

17 Mesenchymal cells degrade HA through a phagocytosis process. Many studies have shown the capability of osteoblastic cells to phagocytose HA crystals. Some phagosomes containing HA particles ingested by human bone cells [104]. This phagocytic activity follows the cellular activation which is marked by protein synthesis and stimulation of RNA transcription. HA crystals undergo dissolution into the phagosome. Similar results have been noted with fetal mouse calvaria [105] and rat osteosarcoma [106]. Fibroblasts are also found to possess the same phagocytic ability

[107-111].

After implantation, biomaterials induce an inflammatory reaction caused by the wound inflicted during the surgical act. The material structural or chemical parameters

[112, 113] as well as the implantation site are known to influence the inflammation process [114, 115]. Macrophages derive mainly from bone marrow precursor cells that divide, producing monocytes that circulate in the blood. In a second step, these cells migrate into the connective tissue, where they mature and are called macrophages. These cells are distributed throughout the body. They are characterized by an irregular surface with pleats, protrusions, and indentations—a morphologic expression of their active pinocytotic and phagocytic activities. Monocytes/Macrophages are among the first cells to colonize the biomaterial surface after implantation, could play a crucial role during biomaterial degradation [116]. The major functions of macrophages are the ingestion of particles, digestion of these particles by lysosomes, and secretion of an impressive array of substances that participate in defensive and reparative functions. Detailed process was studied by Benahmed et al [240]. In their in vitro study, two types of phagocytosis were found when cells came into contact with HA ceramics. The first mechanism concerned 18 the internalization of HA crystals together with a small amount of culture medium by monocytes/macrophages closely attached to the ceramic surface. In this process, the envelope membrane of the phagosomes disappears and releases HA crystals into the cytoplasm and thereby facilitating interaction with the organelles. The phagocytosed particles then undergo dissolution. In the second mechanism, HA crystals are detached from the ceramic surface and internalized together with a large amount of culture medium. Endophagosomes then form heterophagosomes after fusion with primary lysosomes. The dissolution of the crystals occurs in the phagosome. The monocytes/macrophages may undergo differentiation with an accumulation of residual bodies and large fat droplets accelerated by intense phagocytosis in both these two mechanisms. Subsequent to these mechanisms, differentiated macrophages appear incapable of evacuating the debris formed and finally die in culture. In in vivo experiments many researchers reported the presence of monocytes/macrophages at the implantation site of calcium phosphate ceramics [117-121]. Many particles observed in the cytoplasm by electron microscopy are indicative of the phagocytosis of ceramics by monocytes/macrophages. Recently, the phagocytic activity of monocytes/macrophages has been demonstrated also in in vitro models [122-124].

Most researchers agree that the HA implanted in bone could induce the recruitment of two totally different multinucleated populations which are able to degrade the ceramic [125]. Macrophage-polykaryons or foreign body giant cells are the first type, which are formed by fusion of mature monocytes/macrophages. They are observed in chronic inflammatory tissue reaction (granuloma), which are associated with the inflammatory reaction and intervene at the implantation site at early time but disappear 19 late. Macrophage-polykaryons have a limited capacity to resorb calcified matrix. [126].

Macrophage-polykaryons contain a large accumulation of mineral crystals in vacuoles have been described in close association with the implanted ceramic [115, 127].

The second type of multinucleated cells call osteoclasts. These cells are very large, branched motile cells containing 5 – 50 nuclei, which are derived from the fusion of bone marrow-derived cells. In areas of bone undergoing resorption, osteoclasts lie within enzymmatically etched depressions in the matrix know as Howship’s lacunae. The osteoclast secretes collagenase and other enzymes which are packaged in the Golgi complex and secret H+ ions into a subcellular pocket, a ruffled border, between the cell

+ - - and the bone (Protons are from the reaction: CO2 + H2O ! H + HCO3 , HCO3 and the products of bone resorption are taken up by the osteoclast’s cytoplasm, probably digested further, and transferred to blood capillaries). This localized pH drop and enzymes accumulation promotes the localized digestion of collagen and dissolving calcium salt crystals. Osteoclasts are involved in calcified matrix resorption, recruited progressively after implantation [128, 129]. In vitro studies have also confirmed that osteoclasts are capable of resorbing ceramic [130, 131]. Recent researches find that osteoclasts are also capable of forming resorption lacunae on the CaP ceramic surface [132]. The morphology of ceramic crystals inside lacunae will be intensively modified (series of spikes aligned in a single direction), showing similarities with ceramic crystals treated by an acidic solution. This mechanism is similar to the mechanism used by osteoclasts to resorb natural bone [133,134]. When osteoclasts resorb calcified tissues by extracellular acidification, they phagocytose various biomaterial particles (latex, titanium, pholymethylmethacrylate) [135-137]. Osteoclastic phagocytosis of porous HA has also 20 been reported after implantation in sheep mandible [138]. More recently, osteoclasts cultured in vitro on CaP ceramic can develop typical ultrastructural features of bone osteoclasts and are able to degrade ceramic by simultaneous resorption and phagocytosis.

The phagocytotic mechanism is similar to that observed in the presence of monocyters/macrophages.

1.4 The unresolved issues and importance of study on HA degradation

Although HA dissolving is very important for bone remodeling and HA implant bioactivity, it should be designed and controlled according to the specific clinical use

[146]. For example, one should consider the implant site, the cells presence in the implant location, patient sex, age, blood supply in the tissue, load bearing or non-load bearing site, etc., to decide if the biodegradation is preferred or non-preferred in this specific case, and then the implant form and composition can be defined. The implant design and manufacture should satisfy the requirement. However, it is difficult to control the HA dissolving rate due to our limited knowledge of HA degradation and HA synthesis. Some reports show that some direct HA material failure in clinical use are related to HA dissolving [162]. Also technical reason is important for many implant failure, such as HA plasma spray coating always contains some amorphous phase and other HA decomposed products, which cause a much higher degradation rate and clinical failure [148,149].

21 Degradation rate can be accelerated by many reasons, which have been review in the chapter 1.3. The HA biodegradation usually is not preferred in most cases, since a long term mechanical strength of the implant is required in these cases. It is important to know that fast degradation rate could cause following problems in clinical use:

1). A mechanical strength drop, which is especially important in HA coatings.

Microcracks [155] and pits [156] can be found in plasma spray HA coating after in vitro and in vivo degradation tests. These degradation-related defects could cause implants mechanical failure after long term use [157, 158].

2). HA debris or HA particle generation. Although HA or TCP is only composed by Ca,

P, O, and H elements, which are not toxic to the body, but these debris will trigger the immune reaction, which will affect macrophages to produce certain proteins to activate osteoclast and monocyte [159, 160, 126].

3). Fast dissolving affects ceramic-induced osteogenesis, and decreases bone-ceramic direct bonding area [162].

From the view of material science, the inconsistent osseintegration of HA caused by HA dissolving is accompanied not only by the biological experimental conditions, a diversity of the HA material fabrication conditions and materials use conditions also play an important role here. Synthetic HA processing people using now includes powder processing, pressing, sintering, polishing, etc. For HA composite processing, it is more complicated according to the shape or form and performance which will be accomplished. In this study, we will focus on how these HA processings decide the final material purity, microstructure, etc. And more importantly, how they affect the HA dissolving rate and the HA in vivo or in vitro properties, such as interface bonding 22 formation, cell attachment, HA debris generation. This research will give more understanding of the HA biocompatibility and improve the clinical performance of HA.

1.5 Practical issues which influence HA degradation

1.5.1 HA impurity

The impurity of HA is popular in commercial HA products and laboratory synthesis HA [1]. The impurities in the HA research can be divided into two categories:

(1). Other elements such as Mg, C, F, etc., which are not the component elements of

HA—Ca, P, O and H. (2). Calcium phosphate materials or CaO other than HA existing in non-stoichiometric HA, which are composed of Ca, P, O and H but have a Ca/P rate other than 1.67 [29]. The first category of impurities is mainly from raw materials or containers. Carbon may be from air. To control these impurities, one should control the quality of the raw materials and the atmosphere during HA synthesis. These impurities are easy to control. Most of the researches now focus on the second category of HA impurity. Tricalcium phosphate (TCP) and CaO are most popular impurities can be found in HA products. When an as-synthesized CaP material with a Ca/P ratio between 1.5 and

o 1.667 is heated at ~900 C (prolonged heating may be required to eliminate the CO2 if present), the only products formed are essentially stoichiometric HA and β-TCP

(Equation 19) [163], provided there is steam present to prevent the formation of

Ca10(PO4)6O (O-HA). 23 ⎯⎯→~900⎯o C + β + + CaP xCa10 (PO4 )6 (OH)2 y - Ca 3 (PO4 )2 water Carbon Dioxide (19)

Irrespective of the details of the structure of the CaP, if x mol of HA and y mol of β-TCP are formed, the Ca/P molar ratio is (10x/y + 3)/(6x/y +2). x/y can be dertermined from

XRD measurements [163]. This experiment has also been used to test the Ca/P ration of our synthesized HA. This will be discussed in Chapter 4. TCP usually comes from improper HA powder synthesis and improper HA sintering. The HA synthesis methods have been reviewed in chapter 1.2. The HA powder from those synthesis methods has a relative smaller particle size and usually is composed of amorphous phase and tiny crystals observed by XRD [164]. The raw materials ratio other than 1.67 and improper synthesis control will cause the product Ca/P ratio lower than 1.67. Many papers indicate that the aging time and temperature of the HA precursor solution affect the product purity. At least 3 days aging time is required for HA synthesis. Short aging time will cause low Ca/P ratio. High precursor temperature can shorten the aging time [29]. Our result is consistent with them. (Aging the precursor solution for 2 days at 90Co then ball- milling the HA gel with solution totally eliminated α-TCP peaks under XRD, while only ball-milling the HA gel without aging, α-TCP peaks still can be detected under XRD).

Also, sintering process could cause HA decomposition, which is the reason for the presence of some decomposition products after sintering. The high temperature CaO-

P2O5 phase diagram with water present is shown in Figure 1.5. Usually, HA sintering temperature is about 1200Co. Sintering temperature above 1250Co could cause HA decomposition. Many decomposition reactions and decomposition temperatures have been reported, which show that the HA decomposition is not as simple as what the

24 equilibrium phase diagram shows. The following reactions appears more frequently in the literatures:

1). Ca10(PO4)6(OH)2!Ca10(PO4)6O + H2O!3Ca3(PO4)2 + CaO +H2O [165] (20)

2Ca10(PO4)6O! 2Ca3(PO4)2 + Ca2P2O7 + 3Ca4(PO4)2O or 3Ca3(PO4)2 + CaO +H2O

(nonequilibrium state:spray coating-high temperature and fast cooling rate) (21)

2). The natural decomposition of bone:

(1) Ca10(PO4)6(OH)2 ! 2Ca2P2O7 + Ca3(PO4)2 + 3CaO + H2O [166] (22)

(2) Ca10(PO4)6(OH)2 ! 3Ca3(PO4)2 + CaO + H2O (23)

(3) Ca2P2O7 + CaO ! Ca4P2O9 (24)

- 2- 3). 2OH (solid) <=> O (solid) + V(solid) + H2O (gas) (V: vacancy) (25)

o The reaction product is: Ca10(PO4)6(OH)2-2xOxVx. At 1230 C, HA could lose up to 75% of its constitutional water without loss of the apatite structure [167].

4). Ca10(PO4)6(OH)2 ! Ca4(PO4)2O + 2α-Ca3(PO4)2 + H2O [168] (26)

It is reported that the borderline at which the reaction occurred is at 1598±10, 1750±10, and 1838±10 K for partial water pressures of 0.613, 9.81, and 101.3 kPa respectively.

The diversity of the results from different researchers suggests that HA decomposition might be related to raw material properties, HA powder synthesis, presence of impurity, with or without steam protection, equilibrium or nonequilibrium decomposition and details of sintering procedure. Also the product material test methods might play an important role here. Until now no HA decomposition kinetics and decomposition products crystal growth related paper is available.

Jarcho pointed out “it is lack of effective understanding and casual preparative techniques or the use of common commercially available “Reagent Grade” starting 25 materials that has led to a proliferation of calcium phosphate biomaterials with variable

(and frequently nonreproducible) crystal and chemical compositions [16].” The commercial HA having Ca/P ratios ranging from 1.67 to 1.72 are common, and the necessity of maintaining a precise ratio in vivo has not been clearly mandated. Figure 1.6 is an example of impure commercial HA products. This HA powder is from GFS chemicals. These three XRD results are for HA from GFS, β-TCP and pure HA. The HA from GFS has a low sintering density and fracture strength. Also it dissolves very fast.

The XRD shows this material has a Ca/P ratio 1.62, which means it has 27wt% B-TCP.

Tissue response to pure HA phases can differ dramatically from that of HA + specific impurities. Commercially supplied “phase pure” HAs with different Ca/P ratios from 1.62~1.76 evoked different in vitro biological responses. The HA with Ca/P ratio

1.67 was found to undergo the least dissolving compared to ratios of 1.62 and 1.70. The latter two compositions showed the expected presence of small amounts of TCP and

CaO, respectively (Usually HA synthesis raw materials include Ca(OH)2. Redundant

Ca(OH)2 or improper synthesis and sintering could leave some Ca(OH)2 unreacted, which will decompose to CaO at high temperature [169,170]). The presence of intergranular

TCP or CaO should enhance degradation rate, as both are less stable in aqueous solution than HA, which can be found on Figure 1.2. TCP has a higher solubility product than

HA, which should lead to a higher dissolving rate from equation (12). The TCP ceramic is reported to have a dissolving rate 12.3 and 22.3 times faster than the pure HA ceramic, respectively, in the acid and basic media [1]. Also some researchers found that HA with

TCP impurity has a low density (about 80%), especially with α-TCP, which is the high temperature phase of TCP [30]. Biologically, the proliferation amounts of osteoblastic 26 cells are different for different Ca/P ratio HA in vitro [170]. Although CaP materials with

Ca/P ratio 1.62-1.70 are all reported to support cell preliferation and differentiation of cells, enhanced proliferation and rate of differentiation are achieved only in cells incubated on the phase-pure HA. Also Ca/P ratio affects osteoclasts reaction. Yamada et al. studied the influence of CaP ceramic solubility on osteoclastic resorption [139].

Solubility was regulated by varying the ratio of less soluble HA and more soluble β-TCP.

After two days of culture, the pure HA samples were not resorbed by osteoclasts, whereas osteoclasts resorbed the higher resorbable TCP-HA biphasic ceramic samples. Moreover, when the resorption area was measured, it was found that osteoclasts resorbed the biphasic material more extensively than did pure β-TCP. This suggests that the resorption capability of osteoclasts depends on ceramic composition and is thus directly related to ceramic solubility, whereas the monocyte/macrophage lineage is responsible for degradation of the ceramic.

Besides TCP, CaO is another common impurity can be found in HA, especially in plasma-sprayed HA coatings. CaO can be dissolved in water and form Ca(OH)2, which raises the water pH value. This high pH value makes the implant less biocompatible. In vivo the presence of CaO appears to inhibit the formation of bone in close proximity to the implant and impedes the direct apposition of bone on the implant surface. This may result from the hydration of CaO in the physiological environment, leading to a detrimental local elevation in pH via the formation and subsequent dissolution of

Ca(OH)2 from its surfaces [169]. In the case of HA as a coating or composite for structural use which is designed to sustain cells and body fluid dissolving for long time after implantation, any more resolvable impurity phase in HA will accelerate HA 27 degradation rate. In many cases, HA particles can be found within the tissue surrounding the failed implants. These HA particles might be peeled off after HA dissolving in the body or simply from the broken interface. The impurity in HA might play an important role in these implant failure cases [171].

While in some situations, biodegradation is preferred, such as temporary scaffold, temporary barrier and drug delivery device. Most of the bioresorbable materials are polymers. Some researchers want to take advantage of the high dissolving rate of TCP to manufacture so-called bio-resolvable composite, which will be used as a skeleton to induce bone ingrowth in tissue engineering and finally this osteo-progenitor cell seeded ceramic structure will gradually be dissolved by the body after the new bone formed. The only cases for this use are the CaP ceramic/ biodegradable polymer composite as a temporary scaffold and a porous CaP ceramic scaffold [150-152]. In the late case, the scaffold guides cell attachment, growth, and tissue regeneration in 3 dimensions [153,

154]. The degradable implant provides temporary, mechanical support until the natural tissue healed and regained its strength. In order for a temporary scaffold to work properly, a gradual stress transfer should occur: as the natural tissue heals, the degradable implant should gradually weaken. The need to adjust the degradation rate of the temporary scaffold to the healing of the surrounding tissue represents one of the major challenges in the design of a temporary scaffold [31]. While some other researchers found that the mixed HA+TCP form has negative effects both in vitro and in vivo, the biological response depends on the amount of the soluble phase [162]. High TCP content could not induce bone. After a six-month period of implantation, high TCP content implants show obvious degradation of the implants changing their shape and size macro 28 and microscopically. Microscopically, they show aggregates of fine particles and appearance of multinucleated cells.

Although some work has been done on HA dissolution work and HA impurity influence, the detailed dissolving behavior of non-stoichiometric HA and stoichiometric

HA in vitro and in vivo is still not clear. Especially HA as a ceramic material, the microstructure or phase difference of non-stoichiometric HA or stoichiometric HA will certainly cause the dissolving behavior difference. CaP material with different Ca/P ratios dissolving with different rate is clear now, so the microstructure detail of the dissolving behavior and dissolved surface morphological differences of these different CaP materials should be studied with different time period. A dissolving rate difference of in vitro dissolving behavior with and without cells, and in vivo dissolving behavior in different implantation sites should be compared on the microscopic level. Based on these ideas, we tried to compare three different Ca/P ratios HA dissolving behavior in different pH value phosphate buffered solutions (PBS), subcutaneous position and femoral bone by

SEM [172, 173]. From a clinical view, since the Ca/P ratio of commercial HA is always changing due to the synthesis, some commercially impure HA has already caused some clinical failures before but no tolerance limit is available for these impurities till now. It is important for the researchers to study the cells, tissue and body response to these impurities. A comprehensive knowledge of the different purities influence on the implants should be available in the future, which will be an instruction for HA products manufacture. Our work will be helpful for people to evaluate the impurities affecting HA degradation and biocompatibility from material science view-- HA impurity, pH value,

29 grain size, surface microstructure and grain boundary. This research will also help to build the rule for clinical use of HA.

1.5.2. HA form of HA composites

More and more people are interested in the form of HA particles or porous skeleton immersed in a polymer matrix to produce new biomaterials [174-180]. Examples include HA + PLGA [181-183], HA + PMMA for bone cement [184, 185], HAPEX (HA

+ PE) for orbital floor implant or middle ear prostheses [186-189]. The main reason for the HA/polymer composite development is to improve the HA implant mechanical properties. For polymer, HA improves the polymer material biocompatibility. Also some researchers are interested in imitating the real bone microstructure. HA is brittle and has a higher stiffness than bone, which might cause “stress protection atrophy (see chapter 1)”.

Combining HA with polymer decreases the material overall stiffness and brittleness, although polymer also decreases the strength of the implant. Polymer is considered as a non-bioactive material, and HA significantly improve the composite bioactivity. Some in vivo results showed that increased amounts of HA (40 vs. 20v%) in HA-PE composite increased human osteoblast proliferation rate and differentiation [193]. Similar studies showed that increasing HA incorporation into PMMA bone cement improved human osteoblast-like cell adhesion and response [194]. The incorporation of HA in PMMA enhances the biological response of fetal rat calvarial osteoblasts compared to PMMA alone [195]. 30 Generally, there are three forms of HA in all the HA composites:

1). HA implants have been applied by performing as a coating on alloy. The alloy core sustains the loading and HA coating change the implant surface properties and improves the biocompatibility. The most important samples are the artificial teeth and femoral stems. The Co-Cr-Mo stem is plasma coated with a HA coating. The stem is fixed in the bone marrow. HA coating will help the bone growth into the pores of the coating, which is an alternative to regular cement fixation [146]. An investigation by Dental Implant

Clinical Research Group between 1991 and 1995 on 2,641 implants shows that HA- coated implants, regardless of mobility at placement, integrated more frequently and exhibited greater stability than non HA-coated implants [147].

2). HA/polymer composites are widely investigated nowadays. HAPEX middle ear implants was launched commercially in the American Academy of Otolaryngology in

September 1995. The main advantages are: 1. The HA stiffens the polyethylene, and the polyethylene toughens the composite. 2. HAPEXTM shafts may be trimmed to shape using a standard scalpel. 3. The biological response to HAPEX is satisfactory. HAPEX orbital floor implant bonds firmly to the supporting bone and none extruded from the eye, unlike the silicone implant which was only retained using a soft fibrous capsule. In cell culture studies with human osteoblasts, the cells grow and spread over the composite, attaching themselves to hydroxyapatite particles on the surface. In vivo testing has shown that a strong and stable interface is developed between the material and the bone into which it is implanted [190]. HA/PMMA cement is another HA/polymer composite.

Since the PMMA cement has been described as inert materials, with fibroblastic cells observed at the cement/bone interface [200, 201], this interface is often considered to be 31 the weak link in cement-held prostheses, providing a barrier to direct fracture healing.

The addition of small amounts of HA particulate materials, offers the possibility of strengthening the bone/cement interface [202] with a direct bonding of cement to bone.

3). Porous or granule form HA materials have been developed since 1970’s. These materials are usually successfully as bone filler. The potential advantage offered by a porous ceramic implant is the mechanical stability of the highly convoluted interface developed when bone grows into the macropores of the ceramics [142-144]. Macropores must be interconnecting with the hole size in range of from 100 to 400µm in order to meet the requirements for bone ingrowth, since this hole size is required for capillaries ingrowth [145]. A successful use of porous HA without sustaining hard mechanical load is ocular implant, which has produced a significant advance in the treatment of the anophthalmic patient. This new ocular implant appears to decrease implant extrusion and migration and may produce excellent motility of the artificial eye by coupling the ocular implant to the artificial eye by means of a motility peg. Porous HA is chosen here because of two reasons: low density and good biocompatibility. In one case, after 5 years of implantation, the implant shows complete fibrovascularization with nearly complete osteogenesis with hematopoiesis [192]. The eye muscles can be attached directly to ocular implant, allowing it to move within the orbit-just like the natural eye [191]. While in porous HA bone implant, shortly after cellular degradation, osteoconduction begins inside the structure of the ceramic which is then progressively replaced by true bone characterized by a mineralized extracellular matrix, osteoblasts, osteocytes, osteoclasts and neovascularization. Haversian remodeling subsequently takes place, and physiological turnover with resorption-apposition steps can begin [97,98]. 32 There are two issues should be concerned in the HA/polymer or porous, particle

HA. First, since HA in HA/polymer composite is usually porous or particle form, HA areas exposed to body fluid and cells are many separated tiny areas, which like many

“small islands” embedded in polymer substrate. This morphology is different as bulk HA material surface. It is our contention that after the composite material is implanted in body, the non-bioactive polymer part could induce much stronger inflammatory and foreign body reaction than the HA material could. Among the three polymer substrate materials, PMMA and PLGA should be carefully studied. PMMA as a bone cement material has been used for more than twenty years. A certain amount of papers report that non-polymerized monomer (methylmethacrylate) in PMMA cement during the polymerization can be released in water [196, 197] and tissues adjacent to bone [198], which cause cytotoxic effect in media to human fibroblasts [199]. Also, shortly after the

PMMA particles exposed to the macrophages in vitro, the macrophages rapidly release inflammatory mediators and the cells are lethally damaged [203]. In vivo reports show that in prostheses that have undergone revision, the recovered interface has often been shown to contain sheets of macrophages and giant cells closely associated with PMMA particles [204-207]. In addition, several investigators have noted resorption of bone at the interface accompanying such histological changes and have suggested that this resorption may be an essential component of the aseptic loosening process [205-207]. Since the macrophages and giant cells are present on pure PMMA cement, it is important to study whether these cells are present on the HA/PMMA cement. Researchers found that a higher osteoconductivity of HA/PMMA cement was due to a higher bioactivity of the

CaP materials at the cement surface and a lower solubility of the PMMA powder to 33 MMA monomer [208], while inert polymer portion does induce foreign body giant cells

[209]. Also since macrophages and giant cells are both capable to phagocytose HA implants (see chapter 1.3.2), the macrophages covering the PMMA part should influence those surrounded tiny HA islands if the HA particle size is small to certain level, which could causes very severe HA degradation and HA particle generation. This will raise an important issue that HA will be dissolved much faster than in bulk form, while till now few people pay attention to this issue. PLA or PLGA is a bioabsorbable polymer, which is used as bioabsorbable facture-fixation devices for two reasons. First, there is no need to remove the devices after the fracture heals, as is the case with metal fixation devices.

Second, using bioabsorbable implants prevents the stress protection atrophy and weakening the fixed bone that is usually caused by rigid metallic fixation [211, 212]. A slow degradation rate is believed to result in a lower rate of inflammatory tissue reaction

[213,214]. Up to present, the overall results have been favorable [215-217], while some complications have been reported, including late aseptic swelling [218,219], osteolytic change at the implant site [220, 221], incomplete restitution of normal bony architecture in the medullary canal after implant resorption [222, 223], and early micromovement of the implant [224,225]. HA added into PLGA matrix is to improve the mechanical properties and bioactivities in the repair of bone fractures [226, 227]. After 52 weeks test in the medullary cavities of rabbits, some bony tissue contacted the composites without fibrous tissue layers, but the HA particles disappeared increasingly from the rod surfaces over time and that the spaces left by these HA particles formed many pores in the composites surfaces [228]. This increasingly degradation rate may be due to the degradation product-lactic acid accumulation with time, which lower the localized pH 34 around the HA particles and cause faster and faster HA dissolving. It is certainly a problem late if the loosing HA particles can’t be removed and accumulate in body. While due to the short test period by the researchers, this problem was not mentioned in their paper. We believe the HA/PLGA composite degradation should be carefully studied, especially the relation of the PLGA degradation products and the fast HA degradation rate as well as the in vivo problem it might cause after long time period.

Second, for porous or particular HA implants, due to the much higher material exposed surface area to weight ratio, the dissolving rate caused by cells and body fluid in these implants will be much faster than in bulk HA. With this fast dissolving rate, high localized Ca, P ion concentration and peeled HA particles and fragments could influence macrophage activity [162].

The above two issues will lead to several important hypotheses:

1). Impurity influence will be enlarged in these implants. For Ca/P ratio lower than 1.67,

TCP impurity in HA raise the dissolving rate even higher due to the composite or particle

HA form. Researchers already found that high TCP content in biphasic porous CaP implants caused aggregates of fine particles and appearance of multinucleated cells, which was not found in pure HA porous implant [162]. For Ca/P ratio greater than 1.67, since active macrophages was found around the 90-500µm CaO containing HA particles

[210], one should expect that CaO impurity in porous HA will significantly raise the pH value fast short time after implant and cause even severe immune reaction to the implants, which could cause early implant failure-- the most common failure among all the implant failure, while no research paper can be found in this area.

35 2). Also HA ceramic surface dissolving may cause mechanical failure. Since the dissolving causes some parts in HA composite thin, or may cause surface crevices

(especially in the presence of CaO, which causes surface crevices) or micro holes, which increases the local stress and makes brittle material easy to break.

3). The fast dissolving rate in both HA/polymer composites and porous or particular HA might cause some parts of HA material to be peeled off from the implant. These debris might be the source of “particle generation” in vivo, which is believed to be harmful in clinical cases.

All these three possibilities should be understood for safely using HA composites.

1.5.3 HA material particle generation

HA particle generation happens during very aggressive HA degradation, which has been observed frequently [236-239]. Most time it is a localized phenomenon, it should start at some fast degradation areas on HA surface. According to Jongebloed et al

[71-73], HA dissolving rate is not evenly distributed on HA surface under microscope.

Micro holes are formed on top of HA surface after early time even dissolving in pure water, due to the high strain energy around dislocation and inclusion (see chapter 1.3.1).

The dissolving pits appear on the dislocation outlets with the size 0.1-10 µm. While de

Groot suggested that it is actually the micropores remaining in the structure after sintering which determine the rate of bioresorbability [74]. In ceramic structures, micropores (usually less than 5µm in size) are sometimes found more or less uniformly 36 distributed within the gross polycrystalline structure and are frequently localized at the grain boundaries which connect the individual crystallites. It is quite possible that this microporosity could aid in bioresorption by causing microscopic break-up secondary to solution-mediated resorption. Thus, partically dissolved materials could slough off individual crystals or fragments that are sufficiently small to allow for aggressive cell- mediated removal. While particle generation happens not only on micropores from sintering, which can also be found on dense smooth HA implant surface [229]. Similar results were found by us [173]. Since the HA materials used in these papers are from different sources and processed with different procedures, it is hard to review all the steps which happen during HA particle generation. Usually, particle generation on dense smooth HA material surface should start on exposed grain surface and boundary. On stage one, a surface grain morphology modification starts first after a short time period.

After HA was implanted transfemorally in young adult rats for 2-4 weeks, degradation of the hydroxyapatite surface resulted in the presence of unidirectionally aligned crystallites, both unaltered and degraded HA grains present on the HA surface [230]. After HA was implanted in femora of 23 Wistar rats for periods of 1–4 weeks, dissolution of individual grains of the implant surface created a roughened surface topography [231], then the grain boundary area was apparently etched faster than the other grain surface area.

Widened gaps can be found at grain boundary. Also the orientation of the dissolution seemed to be governed by the orientation of each grain with respect to the implant surface. These results show a selective dissolution happen at early degradation period— some grain surfaces are dissolved faster than others and grain boundary area is dissolved much faster than the rest area. On stage two, a marked remodelling process was found on 37 the HA implant surface and resorption lacunae in the periphery of the cavity edges were found. Many defaults were visible in the bulk HA material. These defaults were constituted by zones with increased porosity, also some submicron particles can be found in surrounding tissue area about several microns far away from the material surface

[229]. At stage three, more pores were formed at the HA surface and more particles were released into the surrounding tissue area. Bone matrix was formed inside the HA surface pores, even when the pores were a few microns in size. Osteocytes in close proximity to the ceramics emitted cell expansions toward the material [229]. Comparing stage one and two, one can find that grain boundary play an important role in particle generation. It seems that the undissolved leftover HA grains are the origin of HA particle generation.

The generated HA particles have significant biological influence. Although HA degraded products are Ca and P ions, the particle state can still stimulate cells reaction

(details see chapter 1.3.2). The in vitro results of macrophage phagocytosis indicate that phagocytosis increases with the size and concentration of particles up to 2 µm [140].

With a larger particle size (up to 4.5 µm), phagocytotic activity reaches a plateau, suggesting that saturation depends on the overall particle volume ingested. In these conditions, the mortality of macrophages was similarly increased with size (up to 2 µm) and concentration. Smaller particles (0.6 µm) induce cell mortality only at high concentrations. These results show that macrophage response is dependent on size and concentration but independent of ceramic composition [141]. Some researchers also observed multinuclear-type cells in close proximity to HA implant after particle generation [75, 76]. In some cases, ceramic fragments are indeed found to be present within vesicles of these cells [77, 78]. Moreover, some generated HA particles have 38 clinical persistence and migrate over distances of at least several centimeters to the interface between the femoral head and the UHMWPE articulating surface [232].

Furthermore, HA particles can survive at least a single macrophage encounter [233]; the initial particle size is a factor in determining in vivo persistence. Such HA particles generated from plasma spray HA coating are believed to stimulate inflammatory reactions resulting in localized bone loss [234]. Additionally, it has been suggested that

HA particles play a role in amplifying the pathological processes involved in the development of specific cancers [235]; at worst, particle generation may increase the incidence of implant-associated tumorigenesis.

HA particle generation area is still not very clear now, and the source of HA particle generation remains unclear; also no clear microstructural changes during all the stages for these particles have been published. Since particle generation may affect long- term bone in-growth and mechanical properties, the particle generation source, process, and mechanism should be carefully investigated. The hypothesis that grain boundary dissolving causes particle generation still needs to be proved by more clear evidence.

Also many other factors may affect particle generation, such as purity and form. HA embedded in PMMA will cause much severe particle generation is found in our results

[173]. Besides HA purity and form, some other factors may affect HA particle generation if HA particle generation is mainly decided by HA dissolving (mechanical stress is possibly involved in particle generation), such as HA pressing and sintering procedure, which influence HA porosity, density and grain boundary bonding strength. The ratio of grain boundary area/grain volume decreases when the grain size increases. Large grain size should cause relative low dissolving rate and therefore a low particle generation rate. 39 Long-time sintering is a way to gain large grain according to Kingery [241]. How the grain size will influence HA dissolving and particle generation need to be investigated.

Moreover, histology tests are required since the HA bioconductivity might be affected due to large grain size.

Although HA as a bioactive ceramic has been widely used as bone filler, coating, porous and composite material, many factors of HA degradation are still not clear. Since

HA degradation could cause clinical problem after implantation, for the safety reason it is necessary to carefully study HA degradation. HA purity and form are two important issues in practical use, not much study has been focused on how these two factors influence the HA degradation rate, especially the influence of HA form. HA particle generation after implantation will cause diseases in clinical use, and the relationship of

HA particle generation and HA degradation need to be investigated. Also HA surface degradation might cause HA easy to break under stress. In vitro and in vivo models need to be designed to study HA degradation rate change and biological response according to

Ca/P ratio change and HA form change. This study will give a valuable reference to all

HA degradation related research in the future.

40

Figure 1.1 Hierarchical levels of a human long femur structural organization [31].

41

o Figure 1.2 Solubility isotherms for various phases in the system CaO-P2O5-H2O at 25 C [50].

42

Figure 1.3 Equilibrium diagram for hydroxyapatite. Solid lines are the pH-invariant lines on the solubility surface. Broken lines separate the stability domain of various solution species [58].

43

Table 1.1 List of chemical reactions in Ca-P-H2O solution system. Each of the reaction is plotted in Figure 1.3 and is marked by the respective letter [58].

44

Figure 1.4 Schematic picture of the calcium ions concentration profile through a boundary layer including the calcium adsorbed layer, Nernst layer, and bulk solution [69].

45

Figure 1.5 Phase equilibrium diagram of calcium phosphates in a water atmosphere. In shaded area HA-containing material will be obtained after sintering [161].

46 HA purity

β-TCP peaks

GFS-HA

β-TCP

our HA

25 27 29 31 33 35 37 2 theta

Figure 1.6 An example of impure commercial HA products. HA powder is from GFS chemicals. These three XRD results are for HA from GFS, β-TCP and pure HA.

47

CHAPTER 2

Ca/P RATIO EFFECTS ON THE DEGRADATIO OF HYDROXYAPATITE IN VITRO

2.1 Introduction

Hydroxyapatite (HA) and HA-containing polymer-ceramic composites are important components of current and future biomedical materials. In spite of the well- recognized importance of single phase HA, however, precisely stoichiometric HA (a

Ca/P ratio of exactly 1.67) is not always present. Best et al [242] examined two commercially manufactured “phase pure” HA’s and determined that their different Ca/P ratios (1.65 and 1.76) evoked different in vitro biological responses. Hing et al [170] found that a Ca/P ratio of 1.67 underwent the least “substrate disintegration” compared to ratios of 1.62 and 1.70; no microstructural data regarding the observed disintegration was provided.

There are several reasons for the lack of true phase purity in HA. Many of the processing routes lead to bulk or localized (especially surface) Ca/P ratios that are not the oft-quoted stoichiometric value. CaO can be present originating from either thermal decomposition [243-247] or as deliberate additions that improve thermal stability [170,

48 248]. Tricalcium phosphate (TCP) is another common product of thermal decomposition

[243, 244, 249, 250, 165]. In fact, the effects of difficult-to-detect thermal decomposition may be responsible for conflicting perceptions regarding the stability of hydroxyapatite- based ceramics. A survey of the literature suggests that the moisture content of the furnace atmosphere may be at fault: when sintered in air, higher relative humidities appear to prevent HA decomposition [251,252]. As was recently alluded to [244], a primary factor contributing to variable phase purity could be varying and investigator/season-specific atmospheric humidity. An important consequence is that any atmosphere-driven decomposition process begins in a thin layer at the outer surface, the surface that usually first contacts and conditions the biological environment.

Such subtle and/or commercially tolerated variations in Ca/P ratio could produce inconsistent biological responses in vitro, in vivo and ex vivo. In fact, variable chemical composition is already regarded as a probable source of in vivo failures [253]. The use of

HA in biomaterials applications requires that these variations be investigated; we have initiated a program in which Ca/P ratio variations and their ‘downstream’ effects are studied under physiologically relevant conditions. Both CaO and TCP are known to be more soluble in aqueous solution than HA [254-258]. However, this paper marks the first side-by-side comparison of the effects of both of these impurities under physiologically relevant values of pH. In the absence of complicating biological factors [259] our results show that both the stability of the parent material and the behavior of the dissolved species vary markedly following these exposures. In spite of the considerably greater volumetric percentages of TCP, the relatively small amounts of CaO render the bulk

49 ceramic far more susceptible to degradation. In addition, the presence of these seemingly trivial amounts of CaO is a potent trigger for media-based precipitation in vitro.

2.2 Experiment design

A mechanochemical-hydrothermal method [260] was adapted to prepare HA powders having the desired impurities. As specified by the method, (NH4)2HPO4 powder

(Fisher Scientific, Fairlawn, NJ) was slowly added to a suspension of Ca(OH)2 (Fisher

Scientific, Fairlawn, NJ) in deionized water. However, we found that it was necessary to protect the solution by a nitrogen blanket to exclude atmospheric carbon dioxide and the unintended formation of CaCO3. The Ca/P ratio was 1.67. The resulting slurry was stirred for 10 min, then attritor-milled (Szegvari Attritor) for 10 h at room temperature [260] and dried in a vacuum oven at 80°C (again excluding atmospheric CO2) for 24 h. XRD analysis (Scintag PAD-V) following synthesis showed that only HA peaks were present.

This powder was typically heat treated in air at 900 °C for 1 h prior to compaction. As a key test of method purity [262], extended exposure to 1200°C for 10 hours under moisture protection failed to express any impurities (Figure 2.1).

The Ca/P ratio was adjusted by adding either CaO (Fisher Scientific, Fairlawn,

NJ) or Ca3(PO4)2 (Acros Organics, NJ) powder to achieve the desired Ca/P ratio. Figure

2.2 shows the volume percentage of TCP and CaO in HA versus Ca/P ratio. The CaO ultimate particle size was 0.67 µm; the Ca3(PO4)2, 0.71 µm (SA-CP4 Centrifugal Particle

Size Analyzer, Shimadzu Corporation). For a Ca/P ratio of 1.62, 36.0 g Ca3(PO4)2 was 50 added to 100.0 g HA. For a Ca/P ratio of 1.72, 1.79 g CaO was added to 100.0 g HA. All three Ca/P ratios (1.62, 1.67 and 1.72) were then attritor-milled for 24 hrs in isopropyl alcohol followed by gentle drying at room temperature.

Ten HA disks of each Ca/P ratio were prepared using by pressing at 17 MPa in a

13.0-mm hardened steel die (Fred S. Carver Inc., Menomonee Falls, WI). As produced, each specimen was approximately 13.1 mm in diameter and 3.2 mm in length. All 30 specimens were sintered at 1200°C for 90 min in steam [263]; in our case this consisted of bubbling the input air through water. Measurements of density for all ten samples of each composition shows that the resulting densities were nearly identical (HA:

91.8±0.4%; HA+CaO: 92.8±0.6%; HA+TCP: 91.6±0.4%). X-ray diffraction of the non- stoichiometric ratios revealed the expected presence of CaO and TCP. After sintering, the disk faces were polished using 400, 600, 800 and 1200 grit polishing paper.

One disk of each Ca/P ratio was reserved for morphological comparison with the exposed disks. The other nine disks at each Ca/P ratio were weighed using an electronic balance (Model CA-26, ATI Cahn) having a readability of 10-5 g. They were then divided into three groups of nine disks; each group contained three disks of each ratio. These groups were then placed into 100 ml of either pH 6.8, 7.0 or 7.4 phosphate buffered saline (PBS) solutions having the composition listed in Table 2.1. These specific values of pH were chosen because they were either identical or close to physiological pH; many in vitro papers utilize relatively low pH that may simulate intra-cellular mechanisms of dissolution in vivo but are unrealistic in the context of ‘normal’ extracellular body fluids.

These solutions were then sealed and placed in a 37°C water bath (9101

Refrigerated Circulators, PolyScience) for either 3 or 15 days. At the end of these periods 51 all samples were washed with deionized water for 10 min then placed in a vacuum oven

(Model 282 Isotemp Vacuum Oven, Fisher Scientific) at 120°C for 12 hrs after which net weight changes were recorded. One specimen of these exposed Ca/P ratios was set aside for surface observation by SEM. One disk of each Ca/P ratio from each group was broken by three-point bending to allow for observation of the fracture surface.

Component Amount added to 1 L DI water KCl 0.20 g

KH2PO4 0.20 g

MgCl2•6H2O 0.10 g NaCl 8.00 g

Na2HPO4 2.16 g Glucose 1.00 g Sodium Pyruvate 0.036 g

CaCl2 0.20 g* *Added slowly as 200 mls of a 0.1 g/L solution to avoid precipitation.

The 6.8 and 7.0 pH solutions were produced via titration with 1 M HCl.

Table 2.1 Phosphate buffered saline solution composition; 0.03wt% sodium azide was also present to prevent bacterial growth.

2.3 Results

As Figure 2.3 shows, in no case was statistically significant weight change observed following only 3 days of exposure. After 15 days, small amounts of weight loss

52 were observed from the Ca/P ratio 1.62 and 1.67 specimens at pH levels of 6.8 and 7.0; as expected, the amount of weight loss was larger at the lower values of pH. Somewhat unexpected was the greater weight loss of the stochiometric compound (versus the Ca/P ratio of 1.62) at a pH of 6.8. For the Ca/P ratio 1.72 samples, 15 days of exposure caused significant weight gains at all three values of pH.

SEM analysis (Figure 2.4) showed that all three Ca/P ratio samples initially have a relatively low density of surface defects. Although there is no statistically significant weight loss after 3 days exposure, micron-level pitting still can be observed on all three surfaces. Three days exposure to a pH of 6.8 produced a relatively higher density of pitting (Figures 2.5 and 2.6). The surface damage on the Ca/P ratio 1.62 and 1.67 samples is both relatively uniform and less severe than that observed on the 1.72 samples.

After 15 days exposure to the pH 6.8 solution, the surface damage is considerably more generalized in the non-stoichiometric samples (Figure 2.7). The extent of lateral damage in the stochiometric surface is similar to that observed after three days exposure; however, the depth of pitting in Figure 2.7(b) is considerably greater than that in Figure

2.5(b). The CaO-containing specimen (Figure 2.7(c)) also shows evidence of pits having angular edges rather than the rounded pore morphology observed in the other specimens.

Figure 2.8 reveals that the source of the consistent weight gain of the Ca/P 1.72 samples was the precipitation of well-defined new phases on the surface. This precipitate was isolated to specific areas and did not form a complete layer. For the pH 6.8 and 7.0 cases, high magnification micrographs (Figures 2.8 (a) and (b)) show that these precipitates are composed of a highly porous, platy microstructure. The appearance of the precipitate in Figure 2.8(b) is remarkably similar to that of Yamada et al [265]. At a 53 pH of 7.4, mixed morphologies (plates and rounded particles) were observed (Figure 2.8

(c) and (d)).

We employed image analysis (Clemex Vision Image Analysis, Clemex

Technologies Inc.) to determine the HA grain sizes as revealed by the fracture surfaces

(Figure 2.9). The grain sizes of HA with the Ca/P ratios of 1.62, 1.67 and 1.72 were determined to be 1.16 µm, 0.79 µm and 0.73 µm, respectively .

The relatively small percentage of CaO in the Ca/P ratio 1.72 samples was expected to be isolated to the HA grain boundaries. The fracture surface (Figure 2.9 (c)) verifies this and shows that they have retained their original particle size. CaO-rich HA retains the second phase in a distinct form partly because the addition of CaO is believed to improve the thermodynamic stability of HA at these temperatures [263]. Here, the HA grain size is approximately 92% of the stochiometric grain size consistent with reports in which CaO acts as a grain growth inhibitor [264].

2.4 Discussion

The compositional requirements dictated by the desired +/- 5% variations away from the stoichiometric ratio disproportionately affect the amount of TCP required.

Compositions containing substantial concentrations of TCP have established clinical utility as bone graft substitutes [266]. However, the primary purpose of both non- stoichiometric compositions in this paper is to examine their degradation behavior relative to that of phase pure HA. In the context of a small (and therefore difficult to 54 detect using XRD) amount of HA decomposition at elevated temperatures on an exterior, exposed surface, both non-stoichiometric compositions seem reasonable [243].

The measured weight losses/gains are consistent with the SEM results. SEM has obviously greater sensitivity in detecting degradation or precipitation. The overall trend in weight loss versus pH – more acidic solutions lead to more rapid loss – are as expected. However, even the physiological (and slightly basic) pH of 7.4 also appears to be capable of causing at least limited degradation in all cases.

A recent report suggests that in unbuffered solution the expected greater degradation of the TCP-containing HA is observed [267]. In contrast, the slightly greater weight loss of the stochiometric HA at pH 6.8 (versus that of the HA+TCP) was puzzling; coupled with this is the consistent observation of deeper pits in that surface.

Classical corrosion theory states that pitting corrosion occurs as the material at the bottom of the pit is removed to protect the surrounding surface. Here, in the absence of any galvanic couple between the bottom of a growing pit and the surrounding surface, we consider the theories of Hyakuna et al which suggest that a thin HA layer forms on all exposed surfaces [268]. If replenishment of this layer occurs via transport from specific sites of initial instability/pores like those seen in (Figure 2.5(b)) in which material is continuously removed, preferential deepening of the pits has some plausibility. This may be partly driven by higher localized values of pH within those pores. The region surrounding the initial pits in the non-stoichiometric compositions is inherently less stable and may not be able to benefit from/maintain such a protective layer. Whatever the reason, we paradoxically see greater resistance to degradation and more generalized

55 degradation in the sample containing a large volume percentage (approximately 27%) of the more bioresorbable phase (Figure 2.3(a)).

Even the early stages of pitting produce surface damage spanning many grain diameters. In no case, however, were grains or grain boundaries clearly visible in the pits even at high magnifications. The angularity of much of the pitting on the Ca/P 1.72 surfaces we attribute to mechanical damage during polishing, a common method of surface preparation. Since the grains in this composition are relatively small and separated by weakly bonded CaO particles, hard carbide particles sliding across a surface could generate microcracks between the grains roughly perpendicular to the polishing direction.

Since CaO can not be dissolved in HA from the phase diagram, the CaO inclusion should be gathered at the grain boundary area. Certainly the presence of intergranular

CaO should increase the rate of intergranular corrosion by weakening the HA ceramic grain boundaries. Evidence for this is visible in Figure 2.9(c), where fracture is clearly intergranular and the CaO inclusions appear to be relatively unchanged from their room temperature condition. The greater corrosion caused by the presence of these inclusions is visible after only three days exposure to pH 6.8 solution (Figure 2.5(c)). After 15 days, however, the surface damage is more widespread and involves the apparent loss of approximately 50% of the original sample surface.

56

2.4.1 Weight gain/surface precipitation

Although we consider it a minor part of our study, the weight gain-surface precipitation results provide a compelling argument for a possible role of trace CaO impurities in triggering calcium phosphate precipitation in vitro. Formation of carbonated apatite in vitro [110, 269] or in vivo [270, 22, 23] has previously been thought to occur via the partial dissolution of HA or TCP surface grains leading to an increase in the supersaturation of the immediate microenvironment and precipitation of the new phase

[271]. However, our results show that CaO dissolution apparently provides a stronger driving force for the nucleation and growth of new phases in vitro than either HA or TCP.

Especially compelling is the fact that the volumetric amount of CaO in the Ca/P ratio 1.72 sample is about 16 times less than that of the TCP in the Ca/P ratio 1.62 sample. One reason for this pronounced effect is that rapid CaO dissolution releases Ca2+ ions necessary for calcium phosphate nucleation and causes a localized increase in pH – via the intermediate formation of Ca(OH)2 – in a phosphorous-rich microenvironment.

Higher pH’s certainly favor the formation of HA itself (the HA synthesis method we use is but one example) and this could logically drive the formation of the apparent range of calcium phosphates we observed.

Given the abundance of precipitation observed following 15 days of exposure, the absence of any crystallization following 3 days of exposure of the Ca/P ratio 1.72

57 samples was surprising. Careful examination of the initial surface, however, showed that much of the polished surface is featureless and does not expose CaO grain boundary inclusions. Instead, these and the grain boundaries appear to be covered by a layer of polishing debris that logically consists of primarily HA. A few scattered surface pores, however, reveal both the internal grain boundaries and the smaller pores remaining after the dissolution of the CaO (Figure 2.10). After 15 days exposure, enough of the protective layer of polishing debris has been removed to expose the CaO inclusions to solution thereby catalyzing precipitation.

2.4.2 Potential in vivo significance

While the obvious potential of the CaO-containing composition to form additional calcium phosphates (especially carbonated hydroxyapatite [265]) might be interpreted as a positive characteristic that encourages new bone formation in vivo, the highly localized increases in pH necessary for this process to occur in solution are believed to be responsible for the inhibition of new bone formation/direct bone apposition in a rabbit model [169]. If CaO dissolution is more rapid than any beneficial in vivo response, trace

CaO impurities will apparently have a stronger effect on the in vivo environment than an equivalent (or relatively larger) amount of TCP; both phases form during HA decomposition. Comparatively, these exposures suggest that more generalized corrosion will occur in the TCP- and CaO-containing specimens when exposed to extracellular body fluids. However, exposure to the in vivo environment immediately results in

58 protein deposition; an adherent protein coating could conceivably influence corrosion by decreasing or slowing fluid-surface interactions. Whether pores or other surface defects could disrupt this coating and leading to the preferential deepening observed in Figure

2.7(b) remains to be seen. Subsequent deposition of a highly adherent extracellular matrix further complicates this interpretation. Finally, the interaction of classic, cellular- based mechanisms that have specific effects on HA – such as osteoclasts – to generate highly localized surface structures (e.g., resorption lacunae) cannot be predicted.

However, these observations suggest that any macrophagic response involving the generation of acidic microenvironments should have greater influence on CaO-containing

HA than on the other two compositions.

2.5 Conclusions

CaO and TCP were deliberately introduced into a well controlled HA microstructure to produce a side-by-side comparison of the effects of these impurities on the degradation of the parent HA matrix. These reasonable variations in the Ca/P ratio produce substantial ‘downstream’ effects on microstructural-level degradation under conditions relevant to the environment created by exposure to extracellular body fluids.

While the trends in weight loss were largely as expected relatively small amounts of CaO render the bulk ceramic more susceptible to degradation than much larger amounts of

TCP. In addition, the presence of small amounts of CaO serve as a potent trigger for media-based precipitation (and weight gain) in vitro. Versus pure HA, the TCP sample 59 paradoxically exhibits greater resistance to degradation even though it contains a large volume percentage (approximately 27%) of what is classically considered to be a more soluble phase. The specific morphological observations derived from these samples will be compared to those resulting from in vivo studies utilizing identically-prepared specimens.

60 20000 . 18000 16000 . 14000 12000 10000 8000 . . 6000 ...... 4000 ...... 2000 0 25 30 35 40 45 50 55 2 Theta (degrees

Figure 2.1 XRD pattern of the hydroxyapatite used in this study following sintering at 1200°C for 10 h. (.) Hydroxyapatite. No trace of the 100% 2θ peaks of either TCP or CaO are present.

61

35% 2.0% Ca/P 1.62 Ca/P 1.72 30% 1.6% 25%

20% 1.2%

15% CaO 0.8% TCP 10% 0.4% 5% Ca/P 1.667

0% 0.0% 1.61 1.63 1.65 1.67 1.69 1.71 1.73 Ca/P Ratio

Figure 2.2 Volumetric percentages of TCP and CaO in HA versus the targeted Ca/P ratios. In terms of the volume of second phase added/present, considerably more TCP is required to achieve the same 5% variance away from the stochiometric atomic ratio.

62

(a)

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Figure 2.3 Weight versus time and pH for all three Ca/P ratios. (a) pH 6.8, (b) pH 7.0, (c) pH 7.4. After 15 days exposure, consistent weight gains were observed for the 1.72 samples under all conditions.

63 Figure 2.3 continued

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64 Figure 2.3 continued

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(a)

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Figure 2.4 Representative SEM micrographs of polished surfaces before testing in pH- modified PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72.

66 Figure 2.4 continued

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67 Figure 2.4 continued

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Figure 2.5 Representative SEM micrographs of polished surfaces after 3 days exposure to pH 6.8 PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72.

69 Figure 2.5 continued

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Figure 2.6 Representative SEM micrographs of polished surfaces after 3 days exposure to pH 7.4 PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72. In spite of an extensive search, no precipitates were observed on the surface of any 3 day Ca/P ratio 1.72 sample.

72 Figure 2.6 continued

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73 Figure 2.6 continued

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74

(a)

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Figure 2.7 Representative SEM micrographs of polished surfaces after 15 days exposure to pH 6.8 PBS. Ca/P ratios of (a) 1.62, (b) 1.67 and (c) 1.72. The arrows in (c) point to pits displaying pronounced angularity. The pores in (b) are considerably deeper than those seen in Figure 5(b).

75 Figure 2.7 continued

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76 Figure 2.7 continued

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77

(a)

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Figure 2.8 SEM micrographs of the precipitation on the Ca/P ratio 1.72 sample surfaces after 15 days of in vitro exposure. (a) following exposure to a pH of 6.8 PBS, a precipitate having a platy microstructure having greater connectivity forms; (b) following exposure to pH 7.0 PBS, a precipitate again showing a platy microstructure. (c),(d) following exposure to a pH of 7.4 a two-phase precipitate forms. In none of these cases did the precipitate cover the entire sample surface; the images in Figure 2.7 were obtained from the uncovered areas.

78 Figure 2.8 continued

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79 Figure 2.8 continued

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Figure 2.9 SEM micrographs of the cross-section of HA with Ca/P ratio of (a) 1.62, (b) 1.67, (c) 1.72. The stoichiometric sample displays primarily transgranular fracture; grain size estimates were based on those areas in which fracture was intergranular (see arrow). In (c) the CaO inclusions are visible as Ca-rich (via EDS) grain boundary precipitates having a distinctive surface morphology.

82 Figure 2.9 continued

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83 Figure 2.9 continued

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84

Figure 2.10 High magnification SEM of the interior of a large pore in the surface of the CaO-containing sample after only 3 days exposure to pH 7.4 solution. Although abundant polishing debris partially obscure some features, this pore interior displays abundant evidence of smaller pores along the grain boundaries indicating that, once exposed, the CaO inclusions dissolve leaving distinctive submicron pores.

85

CHAPTER 3

IN VIVO DEGRADATION RESULTS AND MICROSTRUCTURAL

DISASSEMBLY OF CALCIUM PHOSPHATES

3.1 Introduction

Although calcium phosphates are used extensively in a variety of clinical settings, much of our understanding of their behavior is based on specific cell-based in vivo processes. Osteoclastic resorption is probably the most-quoted mechanism for their biological utilization, often to the exclusion of any other potential cellular or extracellular processes. In addition, degradation is often assumed to be uniform over large scales [272-

274] even though the microstructure itself may not be.

In contrast, reports of the generation of calcium phosphate particles in vivo exist when the calcium phosphate in use is hydroxyapatite (HA) [275-281, 120, 35]. For relatively dense, well sintered HA observations of grain loss [281], grain boundary etching and interactions based on the crystallography of the microstructure [231, 282] have been infrequently reported.

86 The purpose of this study is to highlight similar microstructurally-based interactions observed both in vivo and in vitro. This suggests an extracellular method of particle generation that does not require the participation of osteoclasts. We refer to this phenomenon as microstructural disassembly, or the grain-based loosening of an established, fully crystalline microstructure. The fact that we observe this phenomenon both in contact with osteoblasts alone and following subcutaneous implantation is particularly compelling. This behavior is distinct from that observed within resorption lacunae and provides evidence that overall degradation is regulated by more than one process, particularly as the interface becomes relatively quiescent. This provides new appreciation for the ability of natural systems to interact with hierarchically-organized synthetic implant materials at the micron level.

Previous studies have suggested that calcium phosphate microstructure has clinical significance. De Bruijn et al [230] observed both unaltered and degraded HA grains at the surface of a bone-HA interface following 4 weeks of implantation and noted an “intimate association between unidirectionally aligned crystallites of degraded grains and the bone tissue,” showing that variations in the exposed microstructure could control integration. In a subsequent paper Davies and Baldan [231] noted that “the orientation of the dissolution seemed to be governed by the orientation of each grain with respect to the implant surface.” Grain orientation is a complex result of the microstructure’s thermal and chemical history and is dependent upon grain growth, furnace atmosphere and the presence of trace amounts of grain boundary impurities.

In this context, a characteristic often associated with failure/’loosening’ of HA- based implants – particle generation in vivo – has frequently been observed [278, 149, 87 236-238, 157, 283, 284]. Liberated HA particles, once thought to be engulfed and eliminated by macrophages, in fact have clinical persistence and can migrate over distances of at least several centimeters [232]. Furthermore, such HA particles can survive at least a single macrophage encounter [233] and are believed to stimulate inflammatory reactions resulting in localized bone loss [171]. It has also been suggested that free calcium phosphate particles play a role in amplifying the pathological processes involved in tumorigenesis [235].

If these particles have a microstructural origin, initial grain size must then control initial liberated particle size and in vivo persistence. However, no clear microstructural origin for these particles has yet been identified. Do calcium phosphates, in particular

HA, break down uniformly or randomly or is there a microstructural origin for particle generation? This paper provides data supporting the latter contention.

3.2 Experiment

3.2.1 Preparation of pure and biphasic HA

Stoichiometric HA powder (Ca/P ratio = 1.67) was prepared by a modification

[172] of a mechanochemical-hydrothermal method [260]. (NH4)2HPO4 powder (Fisher

Scientific, Fairlawn, NJ) was slowly added to a suspension of Ca(OH)2 (Fisher Scientific,

Fairlawn, NJ) in deionized water. We found it necessary protect the solution using a nitrogen blanket [110] to exclude atmospheric carbon dioxide and the unintended 88 formation of CaCO3. The resulting slurry was stirred for 10 min, then attritor-milled

(Szegvari Attritor) for 10 h at room temperature [260] and dried in a vacuum oven at

80°C again excluding atmospheric CO2) for 24 h. The median particle size was determined to be approximately 0.7 µm (CAPA-4, Horiba). XRD analysis (Scintag PAD-

V) following synthesis showed that only HA peaks were present. This powder was typically heat treated in air at 900°C for 1 h prior to compaction. Extended exposure to

1200°C for 10 hours under moisture protection [110], a key test of method purity, failed to express any impurities as determined by XRD analysis.

The Ca/P ratio of a separate lot of this powder was then adjusted by adding

Ca3(PO4)2 (Acros Organics, NJ) powder to achieve a clinically-relevant [262] biphasic

Ca/P ratio of 1.62. This mixture then underwent the ‘normal’ milling and heat treatment steps experienced by the pure HA (Ca/P ratio of 1.67).

Ten HA disks of each Ca/P ratio were prepared using by pressing at 17 MPa in a

13.0-mm hardened steel die (Fred S. Carver Inc., Menomonee Falls, WI). As produced, each specimen was approximately 13.1 mm in diameter and 3.2 mm in length. All specimens were sintered at 1200°C for 90 min in air with moisture protection [288].

Measurements of density for all ten samples of each composition showed that the resulting densities were nearly identical (HA: 91.8±0.4%; HA+TCP: 91.6±0.4%) and similar to HA densities produced by other investigations [285-287]. X-ray diffraction

(not shown) of the non-stoichiometric ratio revealed the expected presence of TCP. SEM of fracture cross-sections was used to theestablish that the grain sizes were 1.16 and 0.79

µm in the biphasic and pure HA’s, respectively. The disk faces were polished using 400,

600, 800 and 1200 grit polishing paper to produce uniform surface topographies (see 89 Figure 3.1). The final specimens were approximately 9.6 mm in diameter, 2.6 mm in length, and 0.55 g in weight. One disk of each Ca/P ratio was reserved for morphological comparison.

3.2.2 In vitro sample preparation

To improve resemblance of the initial in vitro solid-liquid interface to that in vivo, cell cultures were conducted on the composite surfaces shown in Figure 3.2. Our goal in designing and fabricating these specimens was to expose a narrow bar of HA approximately 20 microns in width and 10 mm in length to the effects of cell culture.

Surrounded by relatively inert PMMA, the effects of osteoblasts would be relatively concentrated compared to that experienced by a ‘standard’ 13 to14 mm diameter disk composed only of HA.

The HA itself was prepared as described in the previous section. HA bars (2.2 mm x 2.2 mm x 10 mm) were cut from these larger disks using a low speed diamond saw by carefully angling the disks within the jig. The HA bars were placed in the bottom of a simple cylindrical mold and immersed in a 30 wt% PMMA (mw: 120,000, Aldrich

Chemical Company), 70 wt% MMA (Aldrich Chemical Company) solution containing

0.25 wt% benzoyl peroxide (Aldrich Chemical Company) as a catalyst. The use of polymeric PMMA (1) increases solution viscosity prior to polymerization to make the bar position less sensitive to vibration or other mechanical disturbances and (2) decreases the amount of hydrogen gas evolved during the polymerization and thus the possibility of 90 bubble formation. The monomer solution was carefully added to 12 mm diameter glass tubes glued to a glass plate. Once filled, the tubes were placed into a sealed aluminum reaction chamber. This chamber was then placed into a water bath; the following heating schedule was used to produce fully polymerized disks free of internal bubbles: 40°C, 30 mins; 43°C, 15 mins; 45°C, 45 mins; 46°C, 12 h; 48°C, 15 mins; 50°C, 30 mins; 65°C,

30 mins; 77°C, 30 mins; 90°C, 30 mins. The resulting polymerized disks were freed from the glass tubes and heated to 80°C under vacuum for 12 h to ensure that no residual monomer remained. Subsequent polishing required great care to expose only the desired amount of HA. As the tip of the HA bar is approximately a right angle, a simple geometric calculation shows that (exposed HA bar width) = 2 x (removed HA bar height). Thus a ~20 µm bar requires the removal of 10 microns from the tip of the bar and the surrounding PMMA. The PMMA disks containing HA rods were initially polished with 1200 grit polishing paper and periodically examined using light microscopy to detect the point at which the tip of the HA bar was just barely exposed as a long strip.

This surface was then polished using 0.05 µm colloidal alumina to produce a uniformly flat, highly polished surface in which changes are readily observed. Since calcium and phosphate dissolution in these systems is influenced by HA surface area [289, 290], this method of preparation standardizes the initial surface area, surface energy, and roughness. While subsequent images will show that polymerization shrinkage has exposed the less-polished edges of the bar, the desired form – a small amount of ceramic surrounded by a much larger amount of polymer – was achieved.

91 3.2.3 Exposure to primary osteoblast culture

Cultures were conducted utilizing the following procedure: osteoblasts were isolated from 19-day-old fetal rat calvariae as described by Bellows [291] with several modifications [292]. Briefly, following an initial treatment of calvariae for 10 minutes at

37°C with 570 U/ml Type II collagenase (Worthington Biochemical Corp., Freehold,

NJ), the cells released from calvariae by two 10-minute and two 20-minute sequential collagenase digestions were pooled. Cells were grown overnight at 25,000 cells/cm2 on plastic dishes (Corning, Corning, NY). The cells were then detached, using 0.05% trypsin and 0.53 mM EDTA in Hank’s buffered salt solution (Gibco, Grand Island, NY). Cells were then plated at a density of 36,000 cells/cm2 on the experimental substrates and tissue culture plastic. The cells were grown in alpha-Minimum Essential Medium supplemented with 10% heat-inactivated fetal bovine serum, 0.5% Fungizone, 1% L- glutamine, 0.1% gentamicin, and 0.5% non-essential amino acids (Gibco).

The HA-PMMA disks and the PMMA controls were removed from culture on days 3, 9 and 15 and washed twice with 0.1 M PBS buffer at 37°C and fixed in 2.5% glutaraldehyde in 0.1 M PBS buffer (pH 7.4) overnight. Samples were then dehydrated through a graded ethanol series (30, 50, 70, 90, and 96%) to absolute ethanol. These were transferred to absolute ethanol for critical point drying with CO2 in a critical point dryer

(Denton model CPD). The dried samples were sputter coated with approximately 80 –

100 angstroms of 60% Au-40% Pd alloy and examined using a JEOL JSM 820 scanning electron microscope operating at 7-10 kV. Media-only samples were prepared and

92 observed in the same way. The edges of the PMMA-HA disks were marked so that the location of the bar beneath the cell layer could be easily determined by tilting the stage.

3.2.4 Subcutaneous implantation

Implantation in non-osseus sites has a long history of providing important information regarding biological interactions with synthetic HA [276, 293] without the placement complications of osseus sites. To achieve this, the samples previously fabricated were then divided into groups of three disks; each group contained one disk of each ratio. The disks were first individually sterilized in a gas autoclave using ethylene oxide gas at 50°C. All food was withheld from the animals for a period of 12 to 24 h before induction of anesthesia to minimize risks of gastric reflux and abdominal distention. General anesthesia was induced using xylazine (0.05 mg/kg body weight, IV), ketamine (4 mg/kg, IV), and diazepam (0.04 mg/kg, IV). An endotracheal tube in placed and surgical plane anesthesia maintained using isoflurane gas vaporized into 100 % oxygen. A 15 x 15-cm square area of the lateral abdomen was clipped and prepared for aseptic surgery. A No. 10 scalpel blade was used to create a 1-cm diameter myocutaneous pocket for placement of the implant. The implant was inserted and the skin closed with

No. 0 polypropylene suture in a simple interrupted suture pattern. Each subject received one implant from each of the study categories (HA and HA-27v%TCP).

93

3.2.5 Bone implantation

For in bone implantation, a 10-cm longitudinal incision was made centered over the cranial medial cortex of the tibia. Each implant disc was inserted until the superficial aspect of the disc was level to the surrounding cortical bone. Each subject received one implant from each of the Ca/P ratios.

3.2.6 Specimen Harvest

On day 15 post-implantation, each animal was humanely euthanized prior to sample retrieval. A surgical incision was made overlying the implants and the tissue environment surrounding the implants evaluated. Each implant was individually identified and stored at -10°C prior to analysis. The explanted specimens were immersed in 3% glutaraldehyde (diluted with DI water, then filtered) solution, placed into a refrigerator for 24-48 h, followed by additional rinsing in DI water, exposure to the graded ethanol series and critical point drying. The dried samples were sputter coated with approximately 80 – 100 angstroms of 60% Au-40% Pd alloy and examined using a scanning electron microscope (Model XL-30 FEG, Philips Electronic Instruments,

Mahwah, NJ) operating at 7-10 kV.

94 To do histology evaluation at the bone-implant interface, a fluorescent label

(oxytetracycline, 22 mg/kg body weight, IV) was administered on days 0 and 15 of the study. Histological evaluation of the bone-implant interface was performed on sections of

50 to 90 µm in thickness using an Exakt Cutting and Grinding System to allow evaluation of the bone-implant interface.

3.3 Results

3.3.1 In vivo study under skin and in bone on bulk HA disks

For all day 3 samples in bone and skin, a few soft tissue interlayer started to form on top of all disk surface. The disk surfaces were all degraded; holes and slot can be found on top of the surface (Figure. 3.3a-c). While for day 15 samples, a histology evaluation showed that new bone formation on HA surface. New osteoid deposition and woven bone formation were noted adjacent to all surfaces. The disk surface degradation is much more severe than day 3 samples. SEM examination of each of the samples post- implantation reveals that each has a distinct surface morphology: the HA+1.6v%CaO

(Ca/P=1.72) sample is heavily damaged, with 20 - 40 µm-wide pores in its surface and approximately 20 µm wide particles approaching release from the surface (Figure. 3.4a); the HA+27v%TCP sample shows a few fairly large (5-10 µm) pores surrounded by large, undamaged areas (Figure 3.4b); the pure HA sample (Figure 3.4c) shows an abundance

95 of small (1-5 µm) pores spread out across the surface. No obvious resorption lacunae were present in any of these samples. In contrast to bone apposition, the results for exposure to non-mineralized tissue shows trends that are qualitatively similar but exaggerated versions of those seen in bone apposition.

On these oxytetracycline-labeled histological sections, new bone start to form on day 15 mineralized samples. It is clear that new bone formation line (white shining line) has already been formed on pure HA surface (Figure 3.5a). New osteoid deposition and woven bone formation were found on whole sample surfaces. Only very few fibrous tissue interlayers were found on HA samples. Most interface area shows a bone-to-HA direct bonding. The HA+TCP implants also showed new bone formation (Figure 3.5b), but bone-to-HA direct bonding was present only in some areas. On the HA+CaO implants not only new bone was formed, but also a thick collagen fiber layer was found

(Figure 3.5c). Also, a few implant fragments were noted in surrounding tissues. These stars show the fragments position. This demostrates HA with only a little amount of CaO could cause HA less biocompatible.

3.3.2 In vivo bone apposition and subcutaneous tests on HA/PMMA composite:

This form - an HA bar surrounded by polymer - was subsequently implanted utilizing the previous bovine model. After 15 days, samples were harvested. The results show the relative higher degradation on these HA bar surface than bulk HA surface

(Figure 3.6). Some HA grains might be pulled out and cracks can be seen between the 96 grain boundaries on bulk HA surface after 15 days degradation in bone (Figure 3.6a).

While on HA bar surface (the HA bar is embedded in PMMA and only a narrow area is exposed, just as described previously), after same days degradation in bone (Figure 3.6b), many grains are obviously losing, and some debris float upon the HA surface. The polishing surface is gone. When exposed to the non-mineralized tissue environment

(Figure 3.6c), HA grain surface degradation are even more severe than bone apposition samples. The grain shape has changed after degradation- the angular area of most loosing grains shows a unique shape after damage. These results show that degradation in apposition to either bone or non-mineralized tissue is considerably more rapid when HA is surrounded by PMMA in this specific form.

3.3.3 Grain pullout in vitro

SEM images showed that the morphology of the cells growing on the PMMA portions of these specimens was identical to that of the PMMA controls. At day 3, the cells were highly adherent and vigorous; a strongly confluent layer forms in which the individual cells appear to be rounded yet well extended in all directions without evidence of blebs [294] on their external membranes.

We observed that at day 3 the integrity of the cell layer above the bar did not survive the rigors of the drying processes used to prepare these specimens for electron microscopy. This allowed visual access to the HA beneath the cell layer and its

97 interaction with the extra-cellular matrix (ECM) (Figure 3.7). In these images we consistently observe that individual grains have clearly been extracted, either one-by-one or in small groups, from the initially monolithic ceramic surface. In all cases, a direct microstructural interaction (especially on the exposed rougher sides of the bar, Figure

3.7b) between the ECM and the ceramic surface is visible. Individual HA grains clearly appear to be better attached to the ECM than to the bar itself. Portions of the bar not in contact with the secreted ECM do not undergo grain extraction (Figure 3.7c). Interactions with the surface were truly microstructural: grain faces and facets are unrounded, indicating that relatively little, if any, dissolution occurred prior to extraction at this early time period.

After 9 and 15 days (Figure 3.8), adhesion of the grains to the ECM is still evident but the grains themselves have undergone considerable degradation. The original HA microstructure was transformed into a loose array of heavily etched individual grains

(Figure 3.7b). The average grain/particle diameter was considerably reduced.

For comparison’s sake, we also examined the behavior of HA bars exposed to media alone without cells. After 15 days exposure and rinsing to remove salts deposited by the media the original polished surface is still visible (Figure 3.9). At no point is grain pullout evident; the actual microstructure itself remains difficult to observe suggesting that polishing debris still cover the grains themselves.

3.3.4 Grain pullout in vivo

98 In vitro exposure to cells provided clear evidence that biological systems can interact with synthetic materials at the microstructural level; as will be shown, this behavior has now also been observed in vivo.

After 15d exposure to the non-mineralized tissue environment, the ECM on

HA+27v%TCP also peeled back during sample preparation (Figure 3.10a) and extracted hundreds of individual grains (Figure 3.10c) in a manner visually identical to that observed in Figure 3.7a. The phase identity of each of these grains (HA vs. TCP) after extraction from the duplex microstructure could not be unambiguously determined via

SEM-EDS.

An approximately 1-2 µm-thick layer of ECM has peeled off the specimen surface (Figure 3.10d). Numerous HA grains were extracted from the surface and were adhered to this layer. An examination of the surface from which the grains were extracted

(left-hand side of Figure 3.10b) shows the expected cavities; grains that remained associated with the parent microstructure are clearly evident at the bottom of these cavities. Interestingly, the surface morphology of the ECM (right-hand side of Figure

3.10b) has reproduced that of the underlying biphasic HA with excellent, nanometer-level fidelity.

The majority of the extracted grains show evidence of only intergranular fracture, consistent with grain-by-grain (or groups of grains) extraction from the microstructure. In a few cases, however, transgranular fracture of the extracted grains was evident (Figure

3.11), typically at the edges of the grain/grains.

At higher magnifications SEM was able to resolve the slight grain boundary degradation below the ceramic-tissue interface (Figure 3.12). At a magnification of 99 100,000 times, the apparently smooth surfaces in Figure 3.10 reveal a series of small steps of approximately 100 nm in height that span the width of each grain.

3.3.5 Pure HA degradation in vivo

For comparison, we observed the character of the degradation of pure HA in vivo.

Classic resorption lacunae were evident (Figure 3.13) but no grain extraction was observed. The heavily etched, rounded remnants of the original microstructure could be observed at the bottom of these lacunae (Figure 3.13b).

3.4 Discussion

A hallmark of the clinical failure of HA is the generation of “loose particles” visible during examination of the surrounding tissue. One repeatedly observed difference between in vivo and in vitro experiments involving calcium phosphates is the absence of these particles in vitro. This lack of correspondence between in vitro experiments and clinical experience decreases the significance of the former.

Under the conditions we employed, the in vitro interaction of osteoblasts and the

ECM they produce in contact with these surfaces yields disassembly mechanisms similar to those observed in the much richer in vivo environment. Simple contact of identical

100 specimens with simulated body fluid at 37°C [172] did not result in microstructurally- based degradation. In the absence of osteoclasts and the complete array of osteogenic substances generated in vivo, grains were loosened and pulled out of the HA surface after culture periods as short as three days. Slightly longer exposures (9 and 15 days) showed that this relatively simple environment can result in continued etching of the HA grains.

Clearly, mechanisms involving grain boundary dissolution are common to both environments. Given: (1) the expected abundance of calcium and phosphorous in the neighborhood of an HA implant; (2) the ability of even acellular collagen to take up calcium and phosphorous to form crystalline apatites [295] and (3) the collagen-mediated epitaxial nucleation of nanocrystalline CHA onto synthetic HA surfaces [293], it seems reasonable to expect that this, along with the irreversible adsorption of various ECM proteins and glycosaminoglycans, produces a bond strong enough to overcome the ability of these surface grains to remain attached to each other. The work of Davies and Baldan

[231] leads us to suspect that a crystallographic dependence between grain presentation and attachment may play a role; grains that are pulled out by the ECM may present common crystallographic planes.

In both cases, grain extraction appears to be a consequence of forces generated during the preparation of these samples for SEM. In vivo, however, small shifts in tissue constituency [296] could also provide tensile forces sufficient to extract loose grains. In addition, the documented ability of osteoblasts to retract collagen gels in vitro [297] suggests that, in vivo, they (or perhaps fibroblasts) could retract deposited collagen anchored to a layer of loose HA grains by epitaxially precipitated HA. However, our assumption that grain boundary dissolution is necessary bears further examination. Could 101 the nature of the bond between the ECM and the exposed grain surface be stronger than that of the grain and its surrounding microstructure? Estimates of grain boundary strength in hydroxyapatites having a similar microstructure suggest values ranging from 40-90

MPa [298]. This is well above the capabilities of any biological process at this level. The logical conclusion, that grain boundary dissolution must play a key role, drove us to re- examine the exposed grain surfaces at higher magnifications (Figure 3.12), producing evidence of stepwise dissolution of the interfacial material between the grains.

The observation of such microstructural factors during HA osseointegration tend to be largely incidental. An example is given by Niki [299], who examined cross-sections of the HA-mineralized tissue interface in a rabbit model and determined that the mode of

HA fracture changed from transgranular to intergranular as the crack neared the interface.

The width of the zone of intergranular fracture increased from approximately 5-8 µm at 4 weeks to 15-20 µm at 24 weeks [299]. We detected a very similar transition (Figure 3.14) even though this is a subcutaneous environment; in both our work and that of Niki we attribute the observed behavior to in vivo dissolution/corrosion of the material along the grain boundary (Figure 3.12) (not to a fundamental change in fracture processes) that allows the surface grains to be extracted. This dissolution at depth has important consequences for the integrity of the ceramic-tissue interface.

Fundamental studies of the mechanical properties of HA in aqueous environments provide relevant background and these are believed by some to be controlled by “ slow crack growth” mechanisms [300, 301]. However, Nonami and Wakai [302] conducted laboratory experiments that clearly showed that under stress and in the presence of

102 moisture, crack progression changes from transgranular to intergranular. They hypothesized that this occurs via the dissolution of a thin (only a few nanometers in thickness) layer of TCP at the HA grain boundaries that may exist even in nominally stoichiometric (Ca/P = 1.67) compositions [303].

This body of data suggests that both the mechanical properties and the biological integration of HA are extremely sensitive to purity and, perhaps, difficult-to-detect (by

XRD, SEM or FTIR) products of decomposition at the grain boundaries. In our case, the biphasic HA certainly contained TCP although the majority of it exists as separate grains in a mixed HA-TCP microstructure. In the grain extraction observed during pure HA degradation in vitro, small amounts of TCP (or some other calcium deficient phase) could exist in XRD-undetectable forms. At the grain boundary, these precipitates would be rendered unusually soluble due to their lattice misfit with the neighboring HA grains. In addition, since their dimensions would be limited to only a few nanometers their solubility may become entirely dominated by the surrounding higher energy interfaces.

Finally, we note that the degradation of the pure HA grains in vitro is faster than that observed for biphasic HA in vivo. This suggests that concentration differences between the two environments drive faster dissolution of the embedded HA bar* in vitro.

Using this in vitro behavior as an estimate of grain degradation rate we can project that the initial grain size will likely control in vivo persistence. This assumption rests on the idea that (1) disassembled HA grains are sitting on the implant surface and (2) that they can be actively degraded before/after they leave the surface. Following the generation of

*In vivo, the HA-PMMA composite sample provoked a strong macrophage response (to the PMMA) that more rapidly dissolved the HA (data not shown). 103 separate particles via dissolution of the grain boundaries, and the mild/absent foreign body response characteristic of many calcium phosphates, even relatively modest initial grain sizes could result in lengthy in vivo persistence. Conversely, very small grain sizes could result in short persistence times and more rapid release of degradation products.

3.5 Conclusions

Direct microstructural interactions between calcium phosphates and the biological milieu have been observed. Preliminary evidence suggests that degradation of these surfaces via osteoblast culture shares specific characteristics with the in vivo interactions between these ceramics and soft tissues at early time periods. That these interactions were common to both experiments constitutes one of the few known examples of in vitro-in vivo correspondence. Interestingly, the degradation of phase pure HA in vitro was more rapid than that of biphasic HA in vivo. It appears that often ignored microstructural factors – grain size, shape, and grain boundary strength/impurities – may in fact contribute to/control degradation. In addition, this data suggests a connection to smaller- scale observations of collagen-mediated epitaxial HA nucleation and growth on pre- existing HA grains as a mechanism by which the strength of a microstructure in which the grain boundary has been degraded can be overcome.

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(a)

Continued

Figure 3.1 Control microstructures of as-polished surfaces via SEM. The Ca/P ratios were (a) 1.62 and (b) 1.67. Polishing debris cover the surface and make direct microstructural observations difficult.

105 Figure 3.1 continued

(b)

106

14.5 mm

PMMA disc ≈20 µm

≈2.0 mm HA rod

≈2.8 mm

Figure 3.2 PMMA disk containing an embedded HA bar. The upper surface was carefully polished to expose a small (20 µm wide) portion of the bar. The HA occupies only 0.12% of the total specimen surface.

107

(a)

Continued

Figure 3.3 HA with Ca/P ratio 1.67, 1.62 and 1.72 degradation for day 3 under skin. All photos were in same magnification. Large holes and slots started to show on all Ca/P 1.72 (HA + 1.6v% CaO) samples (c). Not much soft tissue can be found on top of all day 3 samples.

108 Figure 3.3 continued

(b)

Continued

109 Figure 3.3 continued

(c)

110

(a)

Continued

Figure 3.4 HA with Ca/P ratio 1.67, 1.62 and 1.72 degradation for day 15 in bone. Holes on pure HA surface are relative small and abundant (3.3a,3.4a). A few fairly large (5-10 µm) pores are surrounded by large, undamaged areas on Ca/P 1.62 (HA+27v%TCP) sample surface (b). Slots and pores are bigger than day 3 samples on Ca/P 1.72 sample surface (c). All day 15 samples were covered by soft tissues after sample harvest. These tissues were carefully removed using ultrasonic in ethanol before SEM test.

111 Figure 3.4 continued

(b)

Continued

112 Figure 3.4 continued

(c)

113

(a)

Continued

Figure 3.5 Oxytetracycline-labeled histological sections, new bone start to form on day 15 mineralized samples. New bone formation line (white shining line) has already been formed on pure HA surface (a). The HA+TCP implants also showed new bone formation (b), but bone-to-HA direct bonding was present only in some areas. On the HA+CaO implants not only new bone was formed, but also a thick collagen fiber layer was found (c). The two particles being evolved by the HA+1.6v%CaO composition are each marked with a ‘*.’

114 Figure 3.5 continued

(b)

Continued

115 Figure 3.5 continued

(c)

116

(a)

Continued

Figure 3.6 Comparison of degradation on bulk HA surface and on HA/PMMA composite surface after 15 days. Bulk HA surface in bone apposition (a). HA bar surface of HA- PMMA in bone apposition (b). HA bar surface of HA-PMMA under skin (c).

117 Figure 3.6 continued

(b)

Continued

118 Figure 3.6 continued

(c)

119

(a)

Continued

Figure 3.7 Grain pullout after 3 days in vitro. In (a), grains are extracted from the polished face of the bar. The image in (b) shows that grain extraction also occurs when the ECM comes into contact with the rougher, unpolished sides of the bar. Significantly, the left- and right-hand sides of the image in (c) show that without contact with the ECM (apparently due to bridging of the bar-PMMA gap), no grain extraction is visible.

120 Figure 3.7 continued

(b)

Continued

121 Figure 3.7 continued

(c)

122

(a)

Continued

Figure 3.8 Longer term degradation of HA in vitro – (a) 9 and (b) 15 days. In (a), extraction by the ECM is still visible but the underlying grains have undergone considerable degradation. In (b), a higher magnification shows that after 15 days the average diameter of the grains remaining on the surface continues to decrease.

123 Figure 3.8 continued

(b)

124

(a)

Continued

Figure 3.9 Surface morphology following 15 days exposure to media alone. (a) shows the bar after exposure and mechanical extraction from the PMMA. Some surface discoloration is present probably due to contamination by media salts. (b) shows a higher magnification of the surface in which polishing lines (fine vertical scratches) are still visible.

125 Figure 3.9 continued

(b)

126

(a)

Continued

Figure 3.10 Micrographs of biphasic HA with a Ca/P ratio of 1.62 after 15 days subcutaneous implantation. At low magnifications, the image in (a) shows the ‘peeling’ of the ECM layer on the left side of the micrograph. A higher magnification image (b) shows demarcation between peeled away and still-adherent ECM. The ECM (the right- hand side of the image) precisely replicates the as-polished surface morphology. Without clear evidence of ECM-based extraction from the surface, we could easily mistake this for the HA surface. A small portion of the hundreds of grains studding the peeled-off layer in (a) is shown in (c). Extraction involving groups of (as opposed to just single) grains can clearly occur. In (d) the layer of ECM is ~1-2 µm thick; HA grains can be seen adhering to the opposite surface.

127 Figure 3.10 continued

(b)

Continued

128 Figure 3.10 continued

(c)

Continued

129 Figure 3.10 continued

(d)

130

Figure 3.11 Transgranular cracking (visible in the collection of grains at the center of this image) observed infrequently.

131

(a)

Continued

Figure 3.12 High magnification of superficial degradation of biphasic HA after 15d subcutaneous testing. The step-like cracks on the grain surfaces suggest successive grain boundary degradation.

132 Figure 3.12 continued

(b)

133

(a)

Continued

Figure 3.13 Resorption lacunae (a) present in phase pure HA after 15d subcutaneous implantation. Grain extraction was not observed for this material in vivo. The bottom of the lacunae (b) shows some remnants of the original microstructure that have undergone relatively severe degradation. While grain dissolution appears to be governed by the individual characteristics of each grain, it is clear than the generation of “loose” particles at the bottom of the lacunae is far less likely than dissolution.

134 Figure 3.13 continued

(b)

135

(a)

Continued

Figure 3.14 Cross-section of the 15d biphasic HA-tissue interface. The left-hand portion of each image shows a region displaying strictly transgranular fracture. The right-hand portion (near the surface) shows intergranular fracture. In some portions of the fracture surface in (a) the depth of the intergranular zone is highly variable; the region of transgranular fracture marked with an ‘*’ is almost at the surface while intergranular fracture in the same image extends inwards to a depth of 20 microns. In (b), a higher magnification shows that the interface between the intergranular and transgranular regions is very distinct. No mixed-mode fracture is observed.

136 Figure 3.14 continued

(b)

137

CHAPTER 4

FURTHER STUDY ON HYDROXYAPATITE DEGRADATION

4.1 Introduction

From Chapter 2 and 3, it is clear that HA degradation is affected by Ca/P ratio

(Ca/P ratio 1.62, 1.67 and 1.72), HA form (bulk form vs. HA/PMMA composite form), and implantation site (bone apposition vs. subcutaneous apposition). The Ca/P ratios

(1.62, 1.67 and 1.72) were chosen according to the most commercial HA products Ca/P ranges [16]. Since most calcium phosphate materials turn into HA plus CaO or TCP depending on their Ca/P ratio after sintering [163], it is important to understanding the trends of Ca/P ratio influence more clearly. More detailed and broad Ca/P ratios should be chosen for this study. To avoid the complication caused by in vivo test and in vitro cells and medium tests, pure water should be a good starting point for this study.

Another important issue is derived from the in vitro degradation test in chapter 2.

PBS is an important solution for biological test, while a good simulation of the in bone condition is α-MEM. PBS has strong buffer strength due to its high ion concentration.

However, in body condition, 5% CO2 will keep buffering the body solution around the

138 implant. CO2 buffer is probably not as strong as phosphate buffer, but the CO2 resource is unlimited under body condition. O2 is consumed and CO2 is released from individual cells during the metablism. CO2 is dissolved into the body fluid, so the CO2 concentration in body fluid around the implant should be constant. To simulate this condition and to avoid the complication of in vivo test, a humidity incubator is used in this study. The incubator can control the temperature; humidity and CO2 concentration is air, which gives a good simulation of body condition. The pH value of the sample solutions is an indicator of degradation strength.

Several different sintering procedures will be designed to achieve different sintered HA grain sizes. Since big grain has a relative low boundary area/volume ratio, degradation and particle generation on HA with bigger grain size should be less severe than on HA with smaller grain size. Also grain size might affect macrophage activity in vivo. Since HA is a ceramic material, grain boundary site is considered as defects with relative high energy. If any dissolution happens, the grain boundary site should be dissolved much easier than bulk area. This could be another way to control HA degradation from material science view. Usually sintering temperature and time is the critical elements to control ceramic grain size. In this study, a certain amount of CaO is mixed with HA powder first. Then different temperatures and times are combined to obtain various HA grain sizes. These HA samples then sustain the same degradation test in water. The pH values of these sample solutions are measured to calculate the water penetration depth in each of these HA samples, which will be an indication of HA degradation strength.

139

4.2 Experiment

4.2.1 Ca/P 1.55~1.78 non-stoichiometric HA sintering

Pure HA powder was synthesized by mechanochemical-hydrothermal method

[260]. The pure HA powder was calcified at 800°C for 30 mins. The Ca/P ratio >1.67 non-stoichiometric HA was achieved by adding CaO (Fisher Scientific, Fairlawn, NJ) powder to pure HA power. The Ca/P ratio <1.67 non stoichiometric HA was achieved by adding TCP powder (EM Science, Gibbstown, NJ) to pure HA powder. For Ca/P ratio

1.55, 1.59, 1.61, 1.63, and 1.65, 21.61g, 7.89g, 4.77g, 2.61g, and 1.03g TCP powder was mixed with 10g HA powder respectively. For Ca/P ratio 1.667, 1.68, 1.70, 1.72, 1.74, and

1.78, 0.011g, 0.045g, 0.112g, 0.178g, 0.245g, and 0.379g CaO powder was mixed with

10g HA powder respectively. The mixtures were ball-milled for 24hrs using ethisopropyl alcohol as a solvent in case any dissolving could happen with water as a solvent, and then it was dried at 110°C for 24 hrs. Although non-stoichiometric HA powder could be synthesized directly from solution, which means more than stoichiometric amount

Ca(OH)2 or (NH4)2HPO4 is added according to the theoretically calculation, in real situation, this surplus Ca(OH)2 or (NH4)2HPO4 is very likely dissolved in solution and it is difficult to prevent them from leaching. The mixed powder was then pressed under 80

MPa. After pressing, specimens were sintered at 1200°C for 120 mins in steam. After sintering, each specimen was approximately 9.87 mm in diameter and 1.90 mm in 140 thickness. The relative density of all specimens was about 97.5% (HA theoretical density is 3.16 g/cm3). An automatic polishing procedure was followed using an automatic polishing head (Automet, Buehler). Each disk were polished with 400, 600, 800, 1200 grit polishing paper, and finally 1 µm diamond paste. Each disk surface was checked with optical microscope to make sure all the scratches were gone.

XRD analysis (Scintag XDS2000) was carried on these well-polished samples with Ca/P ratios from 1.55 to 1.78. For each Ca/P ratio, 5 samples were tested. The XRD data were put together to calculate whether the TCP or CaO peak heights are coherent with the theoretical calculation. This test can be used to check the sample real Ca/P ratio.

4.2.2 Degradation Test

Three solutions were adopted for the degradation test: 1). pH 7.4 DI water with few amount KOH to adjust pH to 7.4, 2). pH 7.4 PBS solution (composition refer to

[172]) and 3). α-MEM solution (No.32561, Invitrogen Corporation). For the water and

PBS degradation tests, testing disks and 20 ml vials were sterilized with 70% ethanol, and then one testing disk and 5 ml solution were added to each vial. Then the vials were sealed and put in 37°C water bath. For the α-MEM degradation test, vials were autoclaved and testing disks were sterilized with 70% ethanol. (For the sterilization reason, instead of 5 ml, 10 ml α-MEM was added into each vial initially. At each test time point, only 1 ml α-MEM was taken out for pH test, and the tested α-MEM was not put back.) The vials were put into a 37°C humidity incubator (Water-Jacketed Incubator, 141 Model 3326, Forma Scientific) with 5% CO2. Vials were not sealed; so sterilized air could reach the medium in vails. Samples were also divided into day 3, day 7, and day 15 test times. 11 samples were assigned to each test time and each test solutions.

Additionally, for pH test, three blank pH 7.4water, PBS, α-MEM groups (each group has

11 samples) were added as control groups, which were tested under the same condition as the other groups.

The pH value of solution in each vial was measured with pH meter (Model 710A,

Orion Research, Inc.) at time 1, 3, 8 … 360 hours (15days). For the water and PBS degradation test, each time after a vial was opened to test pH with argon protection, it was sealed again and put back in water bath again.

After degradation test, the disks were cleaned with DI water and ethanol then dried in air. The blank disks and degraded disks were all examined under optical microscope (50X). All the images were taken on sample surface center area. The disk surface images were then imported to computer for surface degraded area quantitative analysis. The etched pits were relative darker than the polished area, these areas were marked with color using “Clemex Vision ™ Professional” software and the stained area percentage was calculated.

4.2.3 HA sintering at different temperatures and times

142 Serveral different sintering temperatures and times were chosen to gain different grain sizes. At low temperature and short sintering time, small grain size will be obtained due to the sintering kinetics. In this study, totally 8 sintering times and temperatures were selected.

After sintering, samples were well polished as previous procedure to remove all scratches on sample top surface. Then the samples were gently sonicated in DI water to remove the overlying polishing debris and expose ceramic surface. All samples were thermal etched at 1050oC 30 minutes to expose the grain boundaries. SEM was used to examine grain size. On each sintering temperature and time point, 4 samples were tested and on each sample surface 4 SEM photos were taken at different areas. Totally 320 grains were measured to gain the average grain size for each data point.

Finally, the 1100oC, 90 mins samples, 1200oC, 90 mins samples, and 1250oC, 50 hrs samples with Ca/P ratio from 1.55 to 1.78 were selected to do the degradation test.

Samples were put in pH 7.4 DI water with small amount of KOH for 15 days. During this

15 days, pH values were recorded to examine the degradation rate difference among these three groups.

4.3 Results

4.3.1 XRD result for Ca/P ratio 1.55~1.78 sintered samples

143 These 12 XRD patterns (Figure 4.1) were used to prove that the HA samples which were obtained from this simple powder mixing procedure have the correct Ca/P ratio finally. The 12 XRD patterns theoretically indicate the Ca/P ratios from 1.55 to

1.78. Above 1.67, only CaO peaks were present, and below 1.67, only β-TCP peaks were present. The peak height should have a linear relationship with Ca/P ratio. The relation was shown in Figure 4.2. The Ca/P ratio and the CaO or TCP main peak height should have the relationship: Ca/P ratio ∝ I0/(I0 + I1).

From our results, a good linear relationship is found based on these XRD data. On

Figure 4.2a, at Ca/P 1.65, the very weak TCP peaks can’t be distinguished from the background noise. The errors of Ca/P 1.74 and 1.78 are relatively greater than other data errors, but the reason is unknown.

4.3.2 Solution influence on HA degradation

During degradation, some ions in HA as well as CaO might be able to be dissolved into solutions or biological fluid. Different solutions will affect this dissolution.

We tested this influence in pH7.4 water, PBS, and 5% CO2 buffered α-MEM(minimal essential medium). 5% CO2 buffer is used to simulate body conditon in α-MEM test.

After 15 days, optical microscope shows that both pure HA and HA with 1.7wt% CaO were etched in DI water, PBS, and α-MEM. But it seems 1.7wt% CaO cause more dissolution pits than in pure HA samples. Etched sample surfaces in α-MEM are shown in Figure 4.3. 144 Figure 4.4 is the pits area percentage which is gained from computer analysis. For pure HA, etching area is about 4%in water and 5% in PBS and α -MEM. But for Ca/P ratio 1.72 samples, because of the CaO presence, etched surface area increases to about

8% in water, 12% in PBS and α -MEM. And in both cases, at beginning, HA or CaO-HA dissolving rate in α -MEM is not as fast as in PBS, while they seem not slow down after 3 days, which is due to CO2 buffering.

Figure 4.5 shows all pH test results of pH 7.4 water, pH 7.4 PBS and α-MEM. In all three solutions, the pH raised by pure HA is very limited, which means the ion release from HA should not cause clinical problem. In Figure 4.5a, pH raised relative fast in water with CaO dissolving from HA-CaO disks. After 18 hours, the average pH value is over 10.5. It is above body cell tolerence. After 15 days, the highest pH value reaches

10.86. The reason for this continual pH increase is because water keeps penetrating sample surface and CaO is being dissolved into water continually. Since pH of body fluid is stable at 7.4, pH change will cause cell malfunction or death. CaO in the etched layer could be dissolved into body fluid and cause clinical problem. In pH7.4 PBS, since the strong buffer function of PBS, very little pH change is found in PBS solution. After

15 days, pH raise to about 7.56. In CO2 buffered α -MEM, pH raise to over 8.2 initially after 2hrs, then it gradually decrease to lower than 7.4 since it is buffered with 5% CO2.

The hydroxy ions release fast at beginning, then it slows down. This suggest CaO could cause clinical problem short time after implantation, while pure HA almost doesn’t affect pH.

145 Based on these pH data for CaO-HA in water in Figure 4.5a, assuming the water penetration is evenly distributed on sample surface, penetration depth has been calculated in Figure 4.6.

From Figure 4.6, it is clear that water penetration rate is fast at the first three days, and then it levels off. The penetration depth can reach 7µm after 15 days, which is 20 times larger than grain size of 1200oC 90 mins sintered HA samples. This could be the reason for the HA grain pull out in vivo, since degradation is more severe in vivo.

The pH change in 5 ml DI water is shown in Figure 4.7. When Ca/P ratio is lower than 1.67, solution pH doesn’t change much. The pH value is within 7~8. When Ca/P ratio is higher than 1.67, solution pH dramatically increases to over 11.

4.3.3 HA microstructure vs. degradation rate in water

The 1150°C for 90 mins, 1200°C for 90 mins and 1250°C for 50 hrs sample surfaces after thermal etching at1050oC were shown in Figure 4.8. It is clear that sintering temperature greatly affects the grain size. After 1200oC 90 mins sintering, which is the normal sintering temperature and time used in most cases, the grain size is 0.36um for pure HA, and HA with CaO. After 1100°C, 90 mins sintering, sintering density is relative low (0.91). To get larger grain size, 1250°C, 50 hrs and 1350°C, 24 hrs sinterings were tried. After sintering, grain grows to 0.84 um and 3.5 um respectively. However, 1350°C is too high for HA sintering even with steam protection. Usually HA decompose at over

146 1300°C. XRD shows TTCP and α -TCP formed. The reaction should be the equation

(26).

1150°C for 90 mins, 1200°C for 90 mins and 1250°C for 50 hrs sintered Ca/P ratio 1.55~1.78 samples were chosen for degradation test in water. The degradation result is shown in Figure 4.9. The pH value of water with 1150°C samples rises higher than other two. For 1250°C 50 hrs samples, degradation rate is low since the grain size is bigger.

4.4 Discussion

Although these results were generated in vitro, the ion release and surface etching likely control the in vivo performance. Truly phase pure HA has very little grain boundary impurities; slight variations away from purity lead to the existence of either

TCP or CaO. Even relatively large amounts of TCP have no negative effects on pH; relatively small amounts of CaO do have substantial effects on pH and therefore negative physiological effects. Very small amounts of impurities can result in substantial increases in localized pH that are themselves far higher than normal physiological values. While dipping in DI water for 3 days would certainly remove surface-bound CaO, this step would do nothing to mitigate in vivo exposure of CaO weeks, months or even years post- operatively as the parent HA matrix is slowly absorbed. CaO is the least biocompatible impurity phase present in plasma spray HA. This unpredictable emergence could

147 certainly add to the challenge of osseointegration and negatively influence even long- term clinical performance.

To avoid CaO dissolving from HA implants, one possible way is to raise sintering temperature and time of HA. This will lead to a relative large HA grain size, thus the ratio of grain boundary area/grain volume drops. This means the water penetration depth drops and degradation rate drops both in vitro and in vivo. Particle generation should also be less severe in this case.

4.5 Conclusion

In vitro tests show that small amount of CaO present in HA will cause local pH increase, which could be the reason for its low biocompatibility.

HA microstructure plays an important role in HA degradation rate. HA Grain size differences cause different degradation rates which is proved by in vitro pH tests.

To completely understand the relationship of HA degradation and biocompatibility, consistent and reproducible connections between the observed in vitro and in vivo effects and materials characteristics should be provided.

148 β-TCP + HA

Ca/P 1.667

Ca/P 1.65

Ca/P 1.63

Ca/P 1.61

Ca/P 1.59

β-TCP Ca/P 1.55

25 30 35 40 2-theta

(a)

Continued

Figure 4.1 XRD patterns of non-stoichiometric HA. (a) Ca/P<1.67, (b) Ca/P>1.67. 149 Figure 4.1 continued

CaO-HA

Ca/P 1.78 CaO

Ca/P 1.74

Ca/P 1.72

Ca/P 1.70

Ca/P 1.68

Ca/P 1.67

25 30 35 40 2-theta

(b)

150 β-TCP + HA

1 0.9 0.8 0.7 0.6 0.5 0.4 TCP-HA I0/(I0+I1) 0.3 0.2 0.1 0 1.5 1.52 1.54 1.56 1.58 1.6 1.62 1.64 1.66 1.68 Ca/P ratio

(a)

Cotinued

Figure 4.2 Linear relationship between the impurity peak height and Ca/P ratios. (a) Ca/P<1.67, (b) Ca/P>1.67.

151 Figure 4.2 continued

CaO + HA

1 0.98 CaO + HA 0.96

0.94

0.92 I0/(I0+I1) 0.9 0.88

0.86 1.66 1.68 1.7 1.72 1.74 1.76 1.78 1.8 Ca/P ratio

(b)

152

(a)

Continued

Figure 4.3 Surface etching after 15 days in α-MEM. Images taken by optical microscrope (50X). Black areas are etching pits. (a) Pure HA surface, (b) CaO-HA surface.

153 Figure 4.3 continued

(b)

154 Ca/P 1.67 HA surface pits area percentage 7 6 5 4 3 water 2 PBS 1 α-MEM

pits area percentage (%) area percentage pits 0 0 5 10 15 etching time (days)

(a)

Continued

Figure 4.4 Pits area percentage gained from computer analysis. 5ml DI water, PBS, and α-MEM were used here. (a) HA surface, (b) CaO-HA surface.

155 Figure 4.4 continued

Ca/P 1.72 HA surface pits area percentage 16 14 12 10 8 6 water 4 PBS

pits area percentage (%) area percentage pits 2 α-MEM 0 0 5 10 15 etching date (days)

(b)

156 pH change in water 11

10 Water blank HA CaO 9 pH value

8

7 0 5 10 15 time (days)

(a)

Continued

Figure 4.5 pH test results of (a) pH 7.4 water, (b) pH 7.4 PBS and (c) α-MEM.

157 Figure 4.5 continued

pH change in PBS 7.6

7.56

7.52

7.48 PBS blank

pH value HA 7.44 CaO

7.4

7.36 0 5time (days) 10 15

(b)

Continued

158 Figure 4.5 continued

pH change in a-MEM 9 8.6 blank a-MEM 8.2 HA CaO 7.8 pH value 7.4 7 051015 time (days)

(c)

159 water penetration depth 9 8 7 6 5 4

depth (um) 3 water penetration 2 1 0 051015 time (days)

Figure 4.6 Water penetration depth on CaO-HA disk surface calculated from Figure 4.5a. Water penetration is fast at the first three days.

160 12

11

10 pH 9

8

7 1.55 1.6 1.65 1.7 1.75 1.8 Ca/P ratio

Figure 4.7 pH change in 5 ml DI water. The vertical broken line represents Ca/P 1.667.

161

(a)

Continued

Figure 4.8 The sample surfaces after thermal etching show the grain sizes. (a) 1150oC, 90 mins, (b) 1200oC, 90 mins, (c) 1250oC, 50 hrs.

162 Figure 4.8 continued

(b)

Continued

163 Figure 4.8 continued

(c)

164 pH change vs. Ca/P ratio 11.5 11 10.5 10 9.5

pH 9 1150C 90mins 8.5 1200C 90mins 8 1250C 50hrs 7.5 7 1.55 1.6 1.65 1.7 1.75 Ca/P ratio

Figure 4.9 Degradation results in DI water after three different sintering procedures.

165

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