Engineering Bioactive Polymers for the Next Generation of Bone Repair
A Thesis
Submitted to the Faculty
of
Drexel University
by
Emily Y. Ho
in partial fulfillment of the
requirements for the degree
of
Doctor of Philosophy
April 2005
© Copyright 2005
Emily Y. Ho. All Rights Reserved. ii
Dedications
To my father,
Ferdinand Fook Hong Ho
iii
Acknowledgements
I would like to give grace to God for His gifts and blessings.
I am grateful to the kindness and guidance of my advisor, Dr. Michele Marcolongo.
Without her continuous advice and financial support, this work would not be possible.
Many people guided me throughout this study and I would especially like to mention,
Dr. Anthony Lowman and Dr. Garland Fussell, for sharing their vast knowledge with me.
Also I thank Dr. Wei-Heng Shih and Dr. Frank Ko for their suggestion in my pre-defense and defense presentations. I appreciated the contribution from Dr. Riad Gobran and Dr.
Surya Kalidindi in my thesis proposal presentation.
Further, I must extend thanks to my friends who assisted me in various experiments:
Jonathan Thomas, Abhijeet Joshi and Marco Cannella on mechanical tests; Thomas
Juliano on nano-indentation analyses; Hui Dong Li on viscosity measurements and
Jennifer Vernengo on polymer characterizations. With their expertise, many of my experiments were made feasible.
Many thanks to my friends, Dr. Grace Hsuan, Dianne Phelan, Kishore Tenneti,
Ranjan Dash, Hoa Lam, Gwenaelle Proust, Rita Truongcao, Chris Massey, members of
Marcolongo’s Lab and members of Lowman’s Lab, who helped me in countless ways.
Special thanks to my Mom and Dad, Grace and Ferdinand Ho, for their care and patience; to my sister and brother, Maria and Alvin Ho, for their encouragement and emotional support; to my aunt and uncle, Oi Chun and William Lam, for their generosity and hospitality. Their endless love and blessing made my life fruitful. Without them, I could not have completed this thesis. iv
Table of Contents
List of Tables ...... viii
List of Figures...... ix
Abstract...... xii
1. Introduction...... 1
2. Background...... 2
2.1. Human Bone...... 2
2.1.1. Structural and Material Properties of Bone...... 3
2.1.2. Mechanical Properties of Bone ...... 4
2.1.3. Bone Cells...... 6
2.1.4. Bone Regeneration...... 7
2.2. Bone Replacements...... 8
2.2.1. Motivation ...... 8
2.2.2. Requirements of Bone Replacement ...... 9
2.3. Natural Bone Replacement ...... 11
2.4. Synthetic Bone Replacement ...... 14
2.4.1. Calcium Phosphate Materials and Bioactive Glasses...... 14
2.4.2. Biodegradable and Non-biodegradable Polymers...... 16
2.4.3. Ceramics / Polymer Composite...... 19
2.5. Tissue Engineered Scaffold ...... 23
2.6. Injectable Materials...... 24
2.6.1. Polymeric Injectable Materials...... 24 v
2.6.2. Ceramic Injectable Materials...... 26
2.6.3. Injectable Composite Materials...... 27
3. Research Goals...... 33
4. An Solid Bioactive Polymeric Composites: HA/PMMA ...... 35
4.1. Introduction...... 35
4.2. Materials and Methods...... 38
4.2.1. Material Processing...... 38
4.2.2. Bending Tests...... 39
4.2.3. Compression Test...... 39
4.2.4. Physicochemical Analysis...... 40
4.2.5. Nano-indentation...... 41
4.2.6. Data Analysis...... 43
4.3. Results...... 43
4.3.1. Global Bending Behavior...... 43
4.3.2. Local Mechanical Properties of HA...... 44
4.3.3. Local Interfacial Mechanical Properties...... 44
4.3.4. Surface Bioactivity...... 46
4.3.5. In vitro Global Elastic Mechanical Behavior...... 46
4.3.6. In vitro Interfacial Mechanical Properties...... 47
4.4. Discussion...... 48
4.5. Conclusion ...... 55
5. An Alternative Approach to Bone Replacement ...... 79
5.1. Introduction...... 79 vi
5.2. Hydrogels...... 80
5.3. Poly(N-isopropylacryamide)...... 81
5.4. Preliminary Studies...... 83
5.4.1. PNIPAAm Co-polymer ...... 83
5.4.2. Mechanical Properties of PNIPAAm-based Hydrogels ...... 84
5.4.3. Bioactive Material Component ...... 84
5.4.4. Phase Transformation...... 86
5.4.5. Rigid Polymer Supplement ...... 87
5.4.6. Polymerization Kinetics...... 88
5.5. Summary...... 88
6. An injectable Bioactive Hydrogel Composite ...... 97
6.1. Introduction...... 97
6.2. Materials and Methods...... 100
6.2.1. Synthesis PNIPAAm-MPS-PEGDM ...... 100
6.2.2. Polymer Characterization...... 100
6.2.3. LCST Determination ...... 101
6.2.4. Injectability and Viscosity Studies...... 101
6.2.5. Change in Volume Estimation ...... 102
6.2.6. Bioactivity Characterization...... 102
6.2.7. In vitro Compressive Test ...... 103
6.3. Results...... 104
6.3.1. Physical Properties...... 104
6.3.2. Bioactivity Kinetics...... 106 vii
6.3.3. In vitro Mechanical Properties ...... 108
6.4. Discussion...... 108
6.5. Conclusion ...... 115
7. Future Work...... 131
8. Novel Contributions...... 133
List of References ...... 135
Vita...... 152
viii
List of Tables
Table 2.1: Mechanical properties of cortical and trabecular bone 26, 27, 30...... 29
Table 4.1: The chemical content of the four groups of composite systems...... 57
Table 4.2: The bending properties of various composite systems...... 58
Table 4.3: Local Young’s modulus and hardness at various positions in uncoupled composite after 10mN loading...... 59
Table 4.4: Comparison of the global mechanical properties of PMMA-MAA treated HA/PMMA composites and human mandible 27, 187-189...... 60
Table 5.1: Mechanical properties of PNIPAAm-based hydrogels reported by various research groups...... 90
Table 6.1: Ion composition and concentration of three biological medias: PBS, SBF and SBFx2...... 116
ix
List of Figures
Figure 2-1: Bone Physiology. Courtesy Gray's Anatomy 35th edits Longman Edinburgh 1973...... 30
Figure 2-2: An illustration of the molecular structure of Poly(glycolide-co-lactide)...... 31
Figure 2-3: Chemical formula of the monomer of PMMA...... 32
Figure 4-1: A schematic representation of a nano-indentation applies onto a hydroxyapatite particulate upon loading...... 61
Figure 4-2: Local Young’s modulus measurements at the center of hydroxyapatite as a function of applied loads; with (a): FE-SEM images of indented hydroxyapatite at 5000x after 10mN loading, and (b): FE-SEM images of cracks induced indentation in hydroxyapatite at 2500x after 25mN loading.62
Figure 4-3: Local hardness measurements at the hydroxyapatite center of 3 to 100mN loading...... 63
Figure 4-4: Typical nano-indentation curves in a ceramic-polymer composite; with FE-SEM image of a typical indentation mark (a) on a hydroxyapatite particulate at 5000x magnification, (b): at the hydroxyapatite- polymethylmethacrylate interface at 5000x magnification, and (c): on a polymethylmethacrylate matrix at 1200x magnification...... 64
Figure 4-5: Graphic representation of the effect of various coupling agents on the local Young’s modulus at the HA-PMMA interface...... 65
Figure 4-6: Graphic representation of the effect of various coupling agents on the local hardness at the HA-PMMA interface...... 66
Figure 4-7: The local elastic modulus and global bending modulus of each coupling agent added composite group was normalized with the controls...... 67
Figure 4-8: Ca/P ratios on a 100µm2 surface area of the controls and treated composites after immersion in SBF at various immersion times...... 68
Figure 4-9: The FE-SEM/EDX images of the PMMA-MAA treated HA/PMMA composite after in vitro chemical reaction from 0 to 24 weeks...... 69
Figure 4-10: Calcium concentration in SBF after immersion of controls and treated composites from 1 to 72 hours...... 70 x
Figure 4-11: Phosphate concentration in SBF after immersion of controls and treated composites from 1 to 72 hours...... 71
Figure 4-12: Global Young’s modulus of the controls and treated composites before and after 24 hours immersion in SBF...... 72
Figure 4-13: Global ultimate compressive strength of the controls and treated composites before and after 24 hours immersion in SBF...... 73
Figure 4-14: Typical load-displacement curves of at the interface of the controls unimmersed and 72 hours immersed...... 74
Figure 4-15: Typical load-displacement curves of at the interface of the PMMA-MAA coupled composites un-immersed and 72 hours immersed...... 75
Figure 4-16: Interfacial Young’s modulus of the controls and treated composites as a function of immersion times...... 76
Figure 4-17: Interfacial Hardness of the controls and treated composites as a function of immersion times...... 77
Figure 4-18: Normalized the global and interfacial Young’s moduli of the treated composite with the controls’ properties...... 78
Figure 5-1: A simplified illustration of the polymer molecular structure of PNIPAAm- PEGDM at 25°C and 37°C...... 91
Figure 5-2: An illustration of the chemical structure of the PNIPAAm-MPS-PEGDM gel with calcium chloride salts...... 92
Figure 5-3: Synthetic procedure of PNIPAAm based hydrogel with 0 to 0.01 molar ratio of MPS...... 93
Figure 5-4: The percentage of accumulative water dispelled volume during the initial 60 minutes of immersion in a 37°C water bath...... 94
Figure 5-5: A plot of compressive elastic moduli (n=5) as a function of PMMA monomer (MMA) content...... 95
Figure 5-6: Observations of the physical properties of the polymer solutions of either 6 or 48 hours polymerized PNIPAAm-MPS-PEGDM with various MPS content (from 0 to 0.01 molar ratio)...... 96
Figure 6-1: FTIR spectra of PNIPAAm-PEGDM with 0, 0.001, 0.005 and 0.01 mole MPS...... 117 xi
Figure 6-2: Phase transformation of PNIPAAm-PEGDM with and without MPS...... 118
Figure 6-3: Viscosity measurements of PNIPAAm-PEGDM with 0, 0.005 and 0.01MPS molar ratio with respect to shear rate from 10 to 100s-1...... 119
Figure 6-4: Average volume change (n=3) of PNIPAAm-PEGDM after immersed in PBS, SBF and SBFx2 from 1 to 544 hours...... 120
Figure 6-5: Average volume change (n=3) of PNIPAAm-MPS (0.005 mole)-PEGDM after immersed in PBS, SBF and SBFx2 from 1 to 544 hours...... 121
Figure 6-6: FE-SEM/EDX images of the bioactive nodules formed within the gel network (a) before immersion; (b) after 3 days; (c) after 5 days; (d) after 30 days; (e) after 45 days and (f) after 60 days immersion in SBF...... 122
Figure 6-7: FTIR spectra indicated the phosphate attachment of the hydrogels as a function of immersion times...... 123
Figure 6-8: Normalized XRD spectra of SBF and FBS-SBF immersed PNIPAAm- PEGDM with and without 0.005 molar ratio MPS...... 124
Figure 6-9: FE-SEM/EDX image of calcium phosphate crystals in the PNIPAAm- PEGDM meshwork after 75 days in SBF...... 125
Figure 6-10: FE-SEM/EDX image of calcium phosphate in the PNIPAAm-PEGDM meshwork after 75 days in FBS-SBF...... 126
Figure 6-11: EDX data of the surface calcium (Ca) to phosphorous (P) ratios of the long term SBF and FBS-SBF immersed MPS-containing (0.005 molar ratio) gels...... 127
Figure 6-12: The composition of chlorine, carbon and oxygen was measured on the cross-section of polymer matrix of 5 and 30 days SBF immersed MPS- containing gels...... 128
Figure 6-13: In vitro elastic modulus of PNIPAAm-PEGDM with various MPS molar ratio after 5 days immersion in PBS, SBF and FBS-SBF...... 129
Figure 6-14: In vitro elastic modulus of PNIPAAm-PEGDM with various MPS molar ratio after 30 days immersion in PBS and SBF...... 130
xii
Abstract Engineering Bioactive Polymers for the Next Generation of Bone Repair Emily Y. Ho Michele Marcolongo, Ph.D. P.E.
Bone disease is a serious health condition among the aged population. In some cases of bone damage it becomes necessary to replace, recontour, and assist in the healing of the bone. Many materials have been proposed as useful replacements but none have been proven to be ideal. In this thesis, two bioactive composites were investigated for bone replacements. First reported material is a hydroxyapatite (HA) particle reinforced polymethylmethacrylate (PMMA) composite treated with a co-polymer coupling agent for mandible augmentations. The influence of the coupling agent on the local mechanical properties of the system before and after simulated biological conditions was determined by applying nano-indentation at the cross-sectional HA/PMMA interface. The local interfacial results were indicative of the global quasi static compression test results.
While the coupling agent improved the interfacial and global mechanical properties before and after 24 hours in vitro immersion, it did not affect the surface bioactivity of the system. However, the addition of coupling agent did not provide long term in vitro improvement of both local and global mechanical properties of the composite. An alternative approach of combining a bioactive phase into polymer matrix was developed.
The second analyzed material is an injectable composite with osteoconductivity and ideal mechanical biocompatibility for vertebral fracture fixations which we formulated and fabricated. A bioactive component was engineered into the macromolecular structure to facilitate the formation of apatite nucleation sites on a xiii thermo-sensitive polymer, poly(N-isopropylacryamide)-co-poly(ethyleneglycol) dimethacrylate (PNIPAAm-PEGDM), through incorporation of tri- methacryloxypropyltrimethoxysilane (MPS). PNIPAAm-PEGDM is capable of liquid to solid phase transformation at 32°C. In this study, the phase transformation temperature
(LCSTs), the in vitro mechanical properties, swelling characteristics and bioactivity of the polymers were evaluated. The addition of MPS to the polymer encouraged apatite formation and increased its compressive modulus while its LCST remained unchanged.
The challenge of this material system is to balance the network-forming and bioactivity inducing MPS with the gain in elastic recovery induced by PEGDM addition to the
PNIPAAm base, all while maintaining an injectable material system. This material platform offers a family of polymers that have a range of mechanical properties for various tissue replacements. 1
1. Introduction
Today's world of orthopedic biomaterials is expanding rapidly. Great demand for orthopedics is driven by aging demographics, product advances, increasing number of sport injuries and changing patient care strategies. Current orthopedic products in the research pipeline include stem cells, bone growth factors, gene therapy, synthetic bone fillers, bioactive implant coatings, biocomposites and resorbable biopolymers among others.
A recent report on the world orthopedic implants and products industry noted that the total orthopedic drug, implant and device market is estimated to be $4.4 billion in
2004 and may rise to $7.4 billion by 2009 1. The orthopedic implants grow with an
average annual growth rate of 11% in the global marketplace since 2002. U.S. dental
bone implant industry is estimated to obtain an annual average revenue growth of 10.6% from 2002 to 2007, with the majority in bone graft material sales 2. This is mainly driven
by the rapid expansion of the aging population (about 10% increment annually) 2. Spinal implants will also see fast sales gains as advances in fusion, fixation and neurostimulating devices improve the success rate and cost-effectiveness of related surgical procedures.
Based on Merrill Lynch’s 2002 report, the U.S. spinal implant devices market is forcasted to generate revenues in excess of $2 billion in 2004 3. The demand for reconstructive
orthopedic implants and trauma fixation products continue expanding at a moderate pace
as the incidence of deteriorated joint conditions and fracture injuries rises with shifting demographic patterns and trends toward more physically active lifestyles 1. 2
2. Background
2.1. Human Bone
Bone tissue has the ability to adapt its mass and morphology to functional demands, its ability to repair itself without leaving a scar, to rapidly mobilize mineral stores on metabolic demand. Bone is considered to be a dynamic example of Wolf’s Law of “form follows function” in biological systems 4. During development of an orthopedic
device, it is always critical to fully understand the biology of bone, its formation and
resorption that orchestrate metabolism and its in situ mechanical behavior.
Bone is a natural composite material, which by weight contains about 45-60%
mineral, 20-30% matrix, and 10-20% water. The matrix is the organic component, which is primarily composed of the protein Type I collagen 5. Similar to all collagens, Type I collagen is a triple helix that is highly aligned, yielding a very anisotropic structure. The mineral, inorganic component of bone is a form of calcium phosphate known as
Hydroxyapatite (HA). Hydroxyapatite has a chemical structure of Ca10(PO4)6(OH)2 and is present in small crystallites form (approximately 2 x 20 x 40 nm). The organic matrix provides bone its flexibility and the inorganic material is predominantly responsible for the mechanical properties of bone 6, 7.
The human skeleton can be categorized into 2 types of bone: the cortical bone and
the trabecular bone (Figure 2.1). Although both bone types comprise the same basic
materials, each one contains different proportions of the organic and inorganic materials, degree of porosity and organization. In addition, the combination of cortical and trabecular bone varies according to the skeleton regions, which is dependent on the 3 applied mechanical loading. Wolff’s Law stated that bone tends to remodel and adapt itself to the surrounding biomechanical environment 4. Various bone tissue types have
different mechanical characteristics that regulate calcium phosphate (or aptite) levels to
bear various combinations of the loading modes: tensile, compressive, bending and
torsion 8. Both cortical and trabecular bones display time-dependent mechanical
behavior, as well as damage susceptibility during cyclic loading 9, 10.
2.1.1. Structural and Material Properties of Bone
Two types of architecture are applied to describe bone: woven and lamellar.
Woven bone structure exists in newborns, callus, and metaphyseal regions. It has a fast
growth rate, is very coarse tissue with no uniform collagen orientation and may slowly
grow into a lamellar structure. Lamellar bone is slowly formed and highly oriented
tissue. Each layer of lamellar bone is embedded with differently oriented collagen. Close
to 100% of adult bone is lamellar bone 11-15.
Cortical bone is a much denser bone type with a porosity ranging between 5% and
10%. It is primarily found in the shaft of long bones and forms the outer shell around
trabecular bone at the end of joints and the vertebrae. Cortical bone is roughly 4 times the
mass of trabecular bone. The main structures in cortical bone are osteon, haversian canal and interstitial lamella (Figure 2.1). The haversian canal contains nerves and blood vessels that transport blood and nutrients to cells, and is surrounded by the trabecular bone. Each osteon is about 20µm in diameter and located less than 100µm away from the central blood supply. Thus, osteons always lie along the haversian canal, they may branch out and intertwine with each other. The interstitial lamella consists of collagen fibers surrounded by a mineralized matrix (mainly calcium and phosphate). The lamella is 4 continuous with the osteons but is not adjacent to any haversian canal 12, 16, 17. Trabecular
bone is the spongy type of bone found at the ends of all long bones and found within flat and irregular bones, such as the sternum, pelvis, and spine. The porosity of trabecular bone ranges from 30 to 90% with pores containing bone marrow. Trabecular bone is
mainly composed of the interstitial lamellae with small cavities of bone cells called
lacunae at the boundaries, and is always surrounded by cortical bone 18-21.
2.1.2. Mechanical Properties of Bone
Although bone is reported to be viscoelastic, it is more commonly characterized
by its elastic mechanical properties. Most reported studies of such properties have been
performed using mechanical testing instruments at quasi-static strain rates 16, 22. The
elastic mechanical properties of bone at the macrostructural level can vary from one bone to another as well as within different regions of the same bone, as shown in Table 2.1.
For a number of human bones (including the femur, tibia, humerus, mandible, lumbar
vertebrae, and patella), the site-specific orthotropic elastic moduli, shear moduli,
Poisson's ratios, and densities have been studied as a function of position using ultrasonic
techniques 23-25. The microstructure of bone is typically oriented in the direction of the
highest biomechanical loading in situ (in consistence with the concept of Wolff’s Law).
This microstructural directionality gives bone anisotropy of mechanical properties. Thus,
from a materials engineering perspective, bone is classified as an anisotropic composite
with open micro-pores 13.
The mechanical properties of human cortical bone from the tibia, femur, and
humerus have been found to vary between subjects, although the density remains the
same. The cortical bone has less regional heterogeneity observed comparing to the 5 trabecular bone. Therefore, the bulk mechanical properties of cortical bones are quite consistent even in various skeleton regions 26, 27. On the other hand, the bulk mechanical
properties of trabecular bone have much broader range and indeed may vary by a factor
of 2–5 from bone to bone. Since the proportion of cortical and trabecular bone of each
bone piece varies with their shape and mechanical function, the properties of bone in
each region cannot be adequately expressed in single values. It is commonly reported in a
range of values reflecting both the experimental difficulties and the observed regional
heterogeneity 28-30. The elastic moduli in the longitudinal direction were not very
different between the various types of cortical bone. There was greater modulus
variability along the length of each piece of bone than around its circumference (less than
10%). The elastic moduli in the radial direction tend to be much lower than the
longitudinal direction (as much as 1000 times less) 16, 24, 25.
In human trabecular bone, mechanical properties vary significantly around the periphery and along the length. They also show significant inter-subject differences 31, 32.
For example, within the human vertebral, the elastic strength and modulus of the trabecular bone in the superior-inferior direction are higher than those in the transverse direction by factors of 2.8 and 3.4, respectively. For a single specimen of bovine femoral bone, ratios of maximum modulus to that in an orthogonal direction can be as high as 7.4.
In human trabecular bone, by contrast, there is no difference in the mechanical properties of the humerus, the proximal tibia, and the lumbar spine. The stiffness and strength of trabecular bone from these bones were found to be lower than those of the patella, and the distal and proximal femur (patellar trabecular bone had the highest values) 28, 33. 6
The mechanical properties of a single bone lamellae were reported by Rho et al. using nano-indentation 34. Nano-indentation method emerged from hardness testing
techniques as high-resolution load and depth-penetrating instruments were developed.
Taking into account that the micro-structural technique offers means by which the
intrinsic mechanical properties of the individual microstructural components of bone may
be measured in a manner which avoids influences of the inherent defects and
heterogeneity in the bone microstructure. Roy et al. also remarked that the nano-
indentation determined local mechanical properties of vertebral bone generally are not
significantly different as a function of bone type (cortical vs. trabecular), but it is
orientation dependent, such as the local elastic modulus and hardness of the longitudinal
direction of the trabecular bone are much higher than the one of transverse direction 35.
2.1.3. Bone Cells
Bone cells are commonly grouped into three types: osteoblast, osteoclast and osteocyte. Osteoblasts are responsible for bone matrix synthesis. They secrete a collagen rich ground substance essential for later mineralization of hydroxyapatite. The collagen actually strands to form osteoids: spiral fibers of bone matrix. Osteoblasts cause calcium salts and phosphorus to precipitate from the blood and bond with the newly formed osteoid to mineralize the bone tissue. Alkaline phosphatase is contained in osteoblasts and is secreted during osteoblastic activity. The osteoblasts have estrogen receptors, which can promote the number of osteoblasts, therefore increasing collagen production36.
About 90% of all cell of human skeleton are made up of osteocytes, which are the key in controlling the extracellular concentration of calcium and phosphate, and are directly stimulated by calcitonin and inhibited by Parathyroid hormone37. These cells 7 derive from bone marrow mononuclear cells. Although the specific function of osteocytes has yet to be found, it had been widely recognized that active bone resorption takes place at the ruffled edges of osteocytes. The main feature of osteoclasts is their ability to resorb fully mineralized bone at sites. The osteoclasts secrete bone-reabsorbing enzymes, which digest bone matrix in order to remodel old and damaged bone. The differentiation, recruitment and inhibition of osteoclasts are controlled by numerous hormonal and growth factors 15, 37-39.
2.1.4. Bone Regeneration
In bone remodeling, osteoclastic resorption with concomitant osteoblastic bone
formation leads to reduction of the metaphysic 36. Active osteoclasts often have a finely striated brush or ruffled border where they are in contact with the bone. In the process of bone resorption, the ground substance of bone appears to become modified prior to resorption by osteoclasts. During osteoclastic resorption, the bone marrow cavity is enlarged that leads to an overall increase of the bone tissue. Then osteoclasts would attached to the tissue surface and proliferate 36, 37.
The renewal of bone is responsible for bone strength throughout our life. Old
bone is removed (resorption) and new bone is created (formation). During childhood and
the beginning of adulthood, bone becomes larger, heavier and denser, bone formation is
then more important than bone resorption. But bone remodeling can be limited by
diseases, aging and injuries. The bone turnover rate is significantly affected with aging
which is about 5% per decade at age 30 and is about 25% per decade at age 60 40. A high bone turnover rate can lead to a weakening of bone strength and increased fracture risk by increasing the amount of bone undergoing the formation and resorption. Thus, 8 temporary or permanent bone replacements may be required to restore the mechanical function of the bone.
2.2. Bone Replacements
2.2.1. Motivation
Bones of the skeletal system provide the supporting structure for the body. Bone can remodel and adapt itself to the applied mechanical environment. Remodeling of bone can be affected by a variety of metabolic and non-metabolic factors including weight bearing, genetics, applied stress, extracellular matrix concentration, hormones, estrogens and nutrients 36. These factors affect the formation and activity of bone cells. In addition, metabolic bone diseases and aging may prohibit bone regeneration, ultimately resulting in progressive loss of bone mineralization 7. Severe bone loss at any skeletal region can
produce fracture, pain and permanent damage, subsequently permanent bone implants
would be required to treat the problem. Osteoporosis is a common skeletal disease that
may cause severe bone loss and may result in fractures of the vertebra, hip and mandible
41. Approximately 200,000 to 500,000 vertebral fractures occur each year in US 42 that are most common among the elderly, individuals with osteoporosis, and postmenopausal women, occurring in 20% of people over the age of 70 years and in 16% postmenopausal women 43. Over 80% of vertebral fractures patients report associated back pain 43. Bone replacements would be needed to treat severe bone fractures to reduce pain and/or to promote bone remodeling. 9
2.2.2. Requirements of Bone Replacement
Williams first defined biomaterials as nonviable materials used in medical device, intended to interact with the local biological environment 32. Two major concerns in the
development of biomaterials are: (1) the material must not impinge on its host (the
surrounding biological systems) and (2) in return, the material must not be affected
adversely by its host. It is therefore essential to have a good understanding of biological systems prior to developing new biomaterials.
Katti has summarized the three general criteria for materials selection for bone replacements 44, as listed:
• The implanted material must not cause an inflammatory or toxic response that would
exceed the acceptable level.
• It must obtain mechanical properties that match with the natural bone tissue at the
implantation site.
• Its manufacturing and processing methods must be economically viable.
A biocompatible material for an orthopedic device must be surface and mechanically compatible. There are several factors that contribute to the surface compatibility of an implant, such as chemical reactions at the implant surface, the toxicity of the applied material, and the implant surface morphology 15, 45, 46. The in vivo
interaction between the implant surface and bone tissue is dynamic. During the initial
stage (within the first few seconds) after implantation, there will be only water, dissolved
ions, and free biomolecules in the closest proximity of the surface 15. As inflammatory
and healing processes proceed, the local biological condition changes that cause an
adsorbed layer of biomolecules on the implant surface until quasi-equilibrium sets in. 10
Eventually, cells and tissues will approach the surface and, depending on the nature of the adsorbed layer, they will respond in specific ways that may further modify the adsorbed biomolecules. Both the chemical composition and topography (structure, morphology) of a surface are known to be important in bone contacting implants, since they regulate the type and degree of the interactions that take place at the interface: adsorption of ions and biomolecules such as proteins; formation of calcium phosphate layers; or interaction with different types of cells (macrophages, bone marrow cells, osteoblasts) 47, 48.
Surface roughness is particularly important for the integration and stability of load-bearing orthopedic devices 47, 49. The nature of the initial interface established between an artificial material and the attached tissue determines the ultimate success or failure of the implant under cyclic loads. Surface porosity is another important factor in bone replacement 50, 51. Klawitter et al. reported that mineralized bone growth into porous calcium aluminate skeletal implants required a minimum interconnective pore size of
100µm, the ingrowth of osteoid tissue required a minimum pore size between 40 and
100µm, and the ingrowth of fibrous tissue required a minimum pore size between 5 and
11µm 52-54. Furthermore, a surface compatible biomaterial for bone replacement must consist of some level of surface wear resistance mechanisms. For instance, clinical results
show that excessive wear and wear debris is the primary cause of failure of ultra high
molecular weight polyethylene (UHMWPE) or metal hip replacements 55. Therefore,
many new hip replacements are composed of ceramic to enhance the in situ surface wear
resistance of the implant.
Matching the mechanical properties of implant with the ones of host tissue is the
key mechanical compatibility factor in the biomaterial selection as bone replacements. 11
Mismatching of stiffness will lead to a major load bearing on the implant device, and the surrounding bone tissue will experience less stress even after the fracture has been repaired. This phenomenon is called stress shielding or stress protection, which affects the bone remodeling and healing process. The underloaded bone adapts to the low stress environment and becomes less dense and consequently weak. Thus, failure of the implantation system is often found to be at the weak interface of the implant and the newly grown high porosity bone tissue 44, 47. Polymethylmethacrylate (PMMA) cements
were the first to be used in orthopaedic surgery and by far the most widely used
cementing material in joint replacements; because of the non-bone bonding characteristic
of PMMA, aseptic loosening is reported to be the most common cause of failure in
cemented arthroplasties using PMMA cement 56. Therefore, bioactive ceramics that
provide direct bone fixation have served as alterative materials for bone cement
applications, but they offer reduced in situ mechanical properties 57. Finally, the
biomaterial should withstand sterilization procedures and have an acceptable shelf life.
2.3. Natural Bone Replacement
Natural bone tissue may be crushed into powder and placed around a fracture or a fusion site. One of the advantages of using natural materials as bone replacements is that
they do not offer the problem of toxicity and may carry specific protein binding sites to assist bone healing or integration. An ideal bone graft material should possess four
elements: (1) osteogenic cells that allow the growth of the various stages of bone
regeneration, (2) an osteoconductive matrix that provides a scaffold for bone ingrowth,
(3) osteoinductive factors that can stimulate osteoprogenitor cells to differentiate into 12 osteoblasts, and (4) structural integrity that can sustain the mechanical loading in situ.
There are three commonly used bone graft materials: autograft, allograft and xenograft.
Bone taken from the patients’s own body is called autograft; it can be classified as either vascularized or nonvascularized 58. It is currently most successful graft material
available with advantages include the capability of immediate bone formation, ease of
incorporation and lack of risk of disease transmission. However, as many as 25% patients
experience difficulties with their bone grafting, such as chronic donor site pain, long
operative time, increased blood loss and risk of transfusion 59. Also autograft is not viable
for patients with severe bone damage that required a large amount of bone graft
materials.
The vascularized grafts retain their existing network of nutrient blood vessels
which allow immediate adaptation of the bone graft with the implantation by providing
an instant and intact blood supply. Therefore these types of bone grafts are particularly
well suited in poorly vascularized recipient beds, such as those exposed to radiation
therapy. Possible donor sites for osseous cranio-maxillofacial reconstruction include
radial forearm, scalcium Phosphateula, anterior iliac crest, fibula and metatarsal. A major
drawback to this form of transplant is that the surgical harvesting of vascularized
autograft is very time consuming, extremely invasive and can create significant morbidity
at the donor site.
Allografts are tissues transplanted from one person to another and are the most
frequently used substitutes for autograft bone in spinal fusion surgery 58, 60, 61. The use of
allograft bone as a graft material has been expanded in recent years as a result of
improved methods of procurement, preparation and storage, technical advances in 13 surgical methods, and the desire to avoid donor site complications associated with using autograft. Although bone allografts are versatile and widely used, especially in spinal surgeries, concerns exist regarding their ability to consistently achieve a successful fusion and the possibility of infectious disease transmission, for example, HIV, Hep B, Hep C.
Allograft has a high infection rate up to 10-12% and more than 80% of infected allografts are associated with clinical failure 60. Thus, knowledge regarding methods of graft
procurement, testing, processing, and storage is necessary for its efficient and safe usage.
Xenograft is a bone graft between two different species, such that bovine, porcine
or coralline bone can be implanted into human graft site. The selection of xenograft depends on each patient’s need. Several studies demonstrated that coralline xenografts have similar performances as autografts when used as filler in defects secondary to trauma, tumors and cysts 62, 63. Similar to allografts, xenografts can be subject to
problems of immunogenicity and have the tendency to denature or decompose in room
temperature or temperature below their melting point. Several research groups recently
developed xenograft combined with osteoinconductive factors, such as bone morphogenetic proteins (BMPs) to enhance the in situ bone generation 60, 64. BMP can
stimulate the migration, differentiation, and activity of potential bone-forming
mesenchymal cells. It exists in the extracellular bone matrix and remains inactive until
demineralized by acid extraction. Once exposed, BMP induces the formation of cartilage
and bone in vivo and stimulates the bone repair process 65. 14
2.4. Synthetic Bone Replacement
2.4.1. Calcium Phosphate Materials and Bioactive Glasses
During the last two decades, ceramic materials have been widely used in many medical applications, hip prosthesis, cardiac valves and dental implants. The bioactive property of ceramic materials has received most attention in the field of hard tissue replacement. Ceramics provide fixation by biological ingrowth of the local tissue into the implant pores, or onto the implant surface a process known as biological fixation 66, 67 through the formation of a biologically active hydroxycarbonate apatite layer on their surfaces in vivo.
Calcium phosphate-based ceramics have been used in medicine and dentistry for nearly 20 years. Synthetic hydroxyapatite (HA) is one of the most popular calcium phosphate materials for bone replacements. Depending on the temperature, the presence of water and pH value, the stable phase of calcium phosphate can be varied. At 37°C and pH > 4.2, only HA is present, but if pH < 4.2, the dicalcium phosphate or brushite will be present.
Synthetic hydroxyapatite has a very similar chemical composition,
Ca10 (PO4 )6 (OH ) , as the natural mineral component of hard tissue. It is characterized by
a calcium to phosphorus ratio of 1.67. Highly crystallized HA tends to be stable in vivo
and is resorbed at a very slow rate of 5% per year via osteoclast activity 68 . But more
soluble forms of HA can be produced by reducing crystal size and adding impurities 69
that would reduce its mechanical properties. HA is generally observed to be non-
bioresorbable and therefore suitable for long-term surgical procedure. High porosity HA 15 provides better chemical bonding at the implant-tissue interface and allows biological bonding through ingrowth of the host tissues into the implant pores 69, 70. But high
porosity can give rise to poor mechanical properties. Also HA is generally very brittle
and hard, leading to difficulties in reshaping and handling during surgical operations.
Therefore, HA alone is not an ideal solution for bone repairing materials. Another
attractive member of the calcium phosphate family for medical applications is tricalcium
phosphate (TCP, Ca3(PO4)2), which plays an important role as a bioresorbable bioceramic
71. TCP would form a apatite layer in body environment and has been used for bone
repair in the form of ceramic blocks, granules or calcium phosphate cements. It is a weak
bioceramic and thus cannot be used on its own as major load-bearing implants in the
human body.
The mechanism of bone bioactivity of calcium phosphate systems can be
summarized into a variety of events: (1) the ionic dissolution from the bioactive surface;
(2) the calcium and phosphate from precipitation from the biological environment onto
the ceramic; (3) continuous ions exchange and structural arrangement at the ceramic-
tissue interface; (4) an amorphous calcium-phosphate layer would be formed at the
ceramic-tissue interface; (5) immediately the amorphous calcium phosphate would be
crystallized and became apatite, the stable phase of calcium phosphate in body
environment. Cellular activity events, such as osteoblast cell attachment, proliferation
and differentiation, also simultaneously occurred at the ceramic-tissue interface with the
bioactivity sequence. An observation of bone tissue formation within the calcium
phosphate implant matrix showed that the interdiffusion of the bioactivity events from the 16 surface boundary layer into the center of a calcium phosphate scaffolds depends on the a several factors: such as the porous size and distribution of the scaffolds.
There are several common bioactive glasses commercially available such as
Bioglass®; bioactive glass-ceramics such as Ceravital®, A/W glass ceramic, or machineable glass-ceramics. Bioactive glass materials are capable of forming an interfacial bond with bone in vivo. However, the time dependence of bonding, the strength of bond, the mechanism of bonding and the thickness of the bonding zone differ for the various materials. Bonding to bone was first demonstrated for a certain compositional range of bioactive glasses consisting SiO2, Na2O, CaO, and P2O5. The
three key elements in these glasses that make the material reactive in an aqueous media
are: 1) less than 60 mol% of SiO2, 2) high content of Na2O and CaO, and 3) high CaO to
48, 71, 72 P2O5 ratio . Many bioactive silica glasses contain about 45 wt% of SiO2 and 5 to 1 molar ratio of Ca to P. Also, glasses with substantially less Ca to P molar ratio are known to be non-bioactive. However, addition of as little as 3 wt% Al2O3 to the system would
prevent direct bone bonding.
2.4.2. Biodegradable and Non-biodegradable Polymers
Polymers used in medical applications can be degradable or not upon exposure to
the biological environment. Biodegradable polymers are materials that will eventually be
resorbed and metabolized after being introduced into a living organism. No additional
surgery is required to remove of the implanted material. Biodegradable polymers are used
as temporary scaffolds for biomechanical and biochemical support. As new tissue begins
to grow within the implantation site, these materials are intended to degrade and leave
behind the regenerated tissue. Poly(DL-lacitde-co-glycolide), PLGA, is a biodegradable 17 polymer. This polymer consists of two monomers: polyglycolide (PGA) and poly-L- lactide (PLA) from the poly(s-hydroxyl acids), as shown in Figure 2.2. These materials have been extensively studied and are FDA approved for certain indications 73. These
polymers have been investigated as porous scaffolds for replacement and regeneration of
a variety of tissues and fixation devices for orthopedic applications 73-76. PLA is more
hydrophobic and less crystalline than PGA and degrades with a lower degradation rate.
The degradable rate of PLGA can thus be easily adjusted by altering the ratio of the two monomers in the formulation. In orthopedic applications the co-polymer can be exploited to create materials that degrade in concert with bone ingrowth. Furthermore, porous
PLGA is known to support osteoblast migration and proliferation, a necessity for bone tissue regeneration 74, 77, 78. However, such polymers on their own are too weak to be used
in load bearing situations. For example prior to degradation, they have less than 1GPa
compressive modulus and less than 1MPa compressive strength and may loss their
mechanical properties up to 50% after several days of implantation 73. Also during the
degradation of PLGA, the polymer molecules would lower the in situ pH level, which
raises difficulties in controlling its degradation rate and may induce an inflammatory
reaction.
Non-biodegradable polymers are designed to withstand high mechanical stresses
experienced with articulating joints; thereby, replace the damaged cartilage and attempt
to support the tissue. This group of polymers is important when the body is unable to
regenerate tissue at the implantation site. Polyethylene (PE) based polymers are
commonly used non-biodegradable polymer group ([CH 2 − CH 2 ]n ). PE is a homo-
polymer, consists of polymer chains made up of single repeating units. Two common 18 forms of PE used as orthopedic devices are ultra high molecular weight polyethylene
(UHMWPE) and high-density polyethylene (HDPE). UHMWPE is widely used as the load bearing surface in joint replacement systems for the past 40 years 79, 80. It was mostly
sterilized under gamma sterilization until recently. Gamma sterilization of UHMWPE in
air-permeable packaging was abandoned because it generated entrapped free radicals 80.
The free radicals may lead to oxidization of the UHMWPE implant in air and contributed
to mechanical degradation and embrittlement of the UHMWPE, before implantation.
Also the wear resistance of this polymer is not desirable and its concomitant sequelae
continue to limit the longevity of the implants 73. Now alternative sterilization methods
and processing methods that can enhance the wear resistance of the polymer are under
investigations. HDPE is a highly linear polymer with a density ranging from 0.94 to 0.97
g/cm3 is obtained with excellent toughness and good chemical resistance. HDPE was
applied as the polymer matrix in a polymer / ceramic composite as a middle-ear drum
replacement 81.
Polymethylmethacrylate (PMMA) is another commonly used non-degradable
polymer in medical devices. It is a self–curing material through the polymerization of
methylmethacrylate within a short time, which consists of repeating homologous chains
of methylmethacrylate monomer (Figure 2.3). Through the chemical manipulation of
these properties and modification of the polymer, controlled behavior variations of this
material can be achieved.
PMMA is typically used as a filler material (bone cement) for joint replacement
surgical procedures. It was originally developed for making dentures, then in 1950s Sir
John Charnley found that it was also suitable for stabilizing prostheses, and is still used in 19 over 80% of hip and over 90% of knee replacements 82. Presently, a variety of commercially manufactured PMMA cements are available. Although commercially available PMMA cements may differ in their components, each powder generally
contains an initiator of polymerization, benzoyl peroxide, and a radiopaque substance, or
opacifier, either barium sulfate or zirconium dioxide, in addition to oligomers of PMMA
itself. The composition differences in these cement provides a range of mechanical
properties of the final products in situ. The quality control of manufactured polymers
relies on the determination of the molecular weight of the polymers, as variations in
mechanical performance of a polymer can be directly correlated to variations in this
characteristic. Generally, larger molecular weight polymers are stronger and more resistant. Medical grade PMMA preparation necessitates manufacturing conditions limited to the clinical milieu, which includes lower curing temperatures and pressurization more suitable to in vivo applications. When PMMA bone cement cures in situ, volumetric shrinkage of PMMA due to density changes has been calculated and experimentally verified to be between 5 to 7% depending on the type of cement and mixing method 83. It was also noted that shrinkage of the polymerising cement in vivo,
migh result in the development of porosity, both at the bone-cement and cement-implant
in cemented hip replacements 84. Redsidual stresses results from the shrinkage of PMMA cement that may lead to reduced fatigue life of the implant system.
2.4.3. Ceramics / Polymer Composite
Ceramics have sufficient biocompatibility that they have high strength along with excellent corrosion resistance. Drawbacks of ceramics are that they are very brittle, exhibit low fracture strength, display low mechanical reliably and are difficult to 20 fabricate. On the other hand, polymers tend to be over flexible and of insufficient strength to meet the mechanical demands as hard tissue replacement. Ceramic coatings on metal and polymer/ceramic composite are a current alternative choice to overcome many shortcomings of materials mentioned above 85, 86.
Various poly(D,L-lactide-co-glycolide) / calcium phosphate composites have
been widely studied 73, 74, 87, 88. PLGA is a biodegradable polymer with an ability to
promote the osteoblast migration and proliferation. But this material does not possess the
same bioactive and osteoconductive properties as calcium phosphate material does.
Therefore a combination of the PLGA and calcium phosphate fillers becomes an attractive alternative as a synthetic bone regeneration composite. While the addition of calcium phosphate (such as HA) leads a more stable pH value than that of pure PLGA; the acidic degradation byproducts of the PLGA polyester scaffold is buffered by the HA.
Zhang and Ma made bone scaffolds from a mixture of HA power with PLGA 89. They
remarked on the non-linear degradation rates of various PLGA polymers that are not
comparable to the growth rate of the host tissue. The difference of the two rates may lead to an overall volume reduction of the implant device. They also reported the overall shrinkage of 50:50 PLGA/HA (75:25 of PLA:PGA in PLGA) is less than 10% and has about 85% porosity 90. Minimizing shrinkage enhances the maintenance of the
mechanical properties of the implant device while they degrade in the body 73. Based on
many research studies, degradation rate and mechanical property of bioactive degradable
composites are governed by optimizing the volume fraction of the two components, the
ratio of the two monomers in the PLGA, the volume fraction of HA in the composite and
the proper processing method. Some preliminary result shows that the 75:25 PLGA 21 scaffold degrades slowly in 6 to 8 weeks 76, 90. Dururan and Brown reported that the
tensile a flexural strengths of their 80:20wt% HA-(85:15)PLGA composites exhibited
25% and 10% reduction, respectively, after 24 hours immersion in 37°C simulated body
fluid 91.
One of the commercially available PLGA/HA scaffolds is the OsteoGraf™, which
is used as bone graft material. PLGA scaffold foam of 75:25 (PLA: PGA ratio) is made from solvent casting/salt leaching technique. The foam consists of a median pore size of
< 50 microns and porosity of > 90% is then rapidly places in a 50/50 polymer/HAP solution. The resultant scaffold’s porosity can be controlled by the amount of salt added, while the pore size is dependent on the size of the salt crystals. Unfortunately the suspension of HA will decrease the porosity slightly (from > 90% to 85.7%. Although this material has not yet provided a good mechanical strength, it showed that bonding layer was greatly reduced from its ability of natural bone ingrowth characteristics 92, 93.
Bonfield’s research group developed one of the first bioactive ceramic / polymer composites for bone replacements with the tradename HAPEX™. They reinforced hydroxyapatite fillers into high density polyethylene 81, 94. The volume fraction of the
hydroxyapatite is the key factor in the resultant modulus. The upper limit of HA content
is 40vol% in HDPE/HA composite in maximizing the mechanical property. On the other
hand, the least HA content is 20vol%, below that limit the composite becomes inert. A
non-adherent fibrous calcium phosphate covered the composite and no chemical bonds
were formed between the implant and the host tissue. Hence a range of 0.2 to 0.4 volume
fraction of HA was determined to be optimal. HAPEX™ could be trimmed by the
surgeon with simple surgical tools. This allowed the surgeon to obtain a better fit, and 22 the implant to have more contact with the host’s tissues. HAPEX™ was developed for total and partial replacement of the conducting middle ear bones. This application is a low load bearing device. Current study of this material has been focused on improving the load bearing property of the material by giving an orientation to the HDPE molecular chains. Also in order to improve the interfacial strength between the two component phases and hence to improve mechanical properties of the composite, various HA particle preparation and filler/matrix coupling methods are being investigated, such as silane coupling and polymer grafting with different HA particle sizes 95, 96. An HA/HAPEXTM middle ear implant is designed to replace damaged malleus and incus ossicles and to restore hearing of patients.
In all composite systems, the degree of contact and the cohesive forces at the interface are of considerable importance; for instance, stress is transmitted from one phase to another across the interface, and the extent of mechanical failure taking place depends on these factors. Therefore, knowing the interfacial mechanical properties, the bulk mechanical behavior of two phase composites can be predicted. The interfacial mechanical behavior of composites has been analyzed using nano-indentation. Nano- indentation is a depth-sensing technique that can accurately characterize the mechanical properties of almost all types of solid materials and the interfacial region of various composites at a small scale 97-102. Hodzic et al. showed that the nano-indentation test can
be employed as a useful tool in the investigation of the interfacial mechanical properties
of their glass/polymer composites with and without silane coupling agent 101, 102. They
determined the silane coupling agent and water degradation play significant roles on the
interfacial mechanisms of their composites 103. 23
2.5. Tissue Engineered Scaffold
One deficiency of commercially available bone replacement materials relates to the lack of both tissue-forming cells and the material’s mechanical compatibility with the host tissue. Biomaterials seeded with stem cells and growth factors in biocompatible scaffolds, faster and more successful regeneration of the natural tissue can be achieved.
In summary, three essential elements are needed to successfully engineer a biological tissue or organ: 1) tissue-forming cells, 2) biocompatible scaffolds conducive to normal cell functions; and 3) quantitative measures of tissue’s regenerative outcome. In tissue engineered scaffold, mesenchymal stem cells (MSCs) may be included to offspring bone cells. These stem cells can be readily extracted from the bone marrow of adult mammals
(including humans), and induced to differentiate into natural tissue. Growth factors such as basic fibroblast growth factor (FGF-2), transforming growth factor β1 (TGFβ1), bone morphogenetic (BMP-2) also would be applied in the tissue engineered scaffolds to promote bone formation. The based materials of the scaffolds can either be permanent or biodegradable, naturally occurring or synthetic, need to be non-toxic, ideally osteoinductive, osteoconductive,and mechanically compatible with native bone to fulfill their desired role in bone tissue engineering 104-108. These materials provide cell
anchorage sites, mechanical stability, and structural guidance and within an in vivo
milieu, provide the interface to respond to physiological and biological changes and to remodel the extracellular matrix in order to integrate with the surrounding native tissue.
Several research groups developed tissue engineered scaffolds using degradable
polymers or bioactive ceramic / polymer composites 108-113. Lu et al. conducted studies on
using a polymeric-based BMP delivery system to investigate the feasibility of using 24 muscle-derived cells for bone tissue engineering applications 114. A two-dimensional,
osteoinductive matrix was formed by suspending BMP in a dissolved PLGA solution.
Their studies showed the efficacy of BMP released from a polymeric substrate to induce
expression of the osteoblast phenotype and secretion of a mineralized matrix by
musclederived cells cultured on the construct. Another research group in UK, Yang et al.
generated a three-dimensional porous PLGA scaffold incorporating osteoblast simulating
factor-1 115, 116. They demonstrated that these materials have the potential of in vivo bone
formation that exploits cell-matrix interactions and, significantly, realistic delivery
protocols for growth factors in bone tissue engineering. Kaito et al load
ed BMPs in a HA filled degradable polymer composite for bone defect repairing. They
found that its tfailure load increased to 800N from 100N after 8 weeks in vitro. With
other finings, this material presents an excellent combination carrier/scaffold delivery
system for rhBMP-2 that it can strongly promote bone tissue regeneration and provide
sufficient mechanical properties in vivo 117.
2.6. Injectable Materials
2.6.1. Polymeric Injectable Materials
Developing a non-invasive method to replace and treat bone tissue is always desirable. An injectable bone replacement device that required no or a minimum of open surgery can minimize surgeries risk and maximize the treatment outcomes. Ideally, these materials should be injectable with established systems of syringes and needles, set in minutes without excessive heat generation, possess the biocompatibility or bioactivity to
favor bone regeneration, and be strong and stable enough to allow for early and active 25 loading. So far there is however no product that fulfils all necessary criteria for these types of treatments.
The polymethylmethacrylate injectable bone cement is most commonly used 118,
119. However, the commercially available formulations are altered by clinicians to
produce cements that are more amenable for use in vertebroplasty. For example, the
monomer to powder ratio is changed to decrease viscosity and increase working time but
this change affects the properties of the materials 119, 120. In additions, the in situ polymerization of PMMA cement is an exothermic reaction, the temperature was measured in range of 39 to 113°C in a study by Belkoff and Molloy 118. Such elevated
temperature is sufficiently high to pose a risk of thermal necrosis for the vertebroplasty patients. Many recent clinical reports showed a substantial number of patients with osteoporosis develop new fractures after undergoing percutaneous vertebroplasty; two- thirds of these new fractures occur in vertebrae adjacent to those previously treated 121, 122.
The two major factors may lead to the new fracture and implant failure are: mismatch of
modulus of the PMMA cement (1.8 to 3GPa) with cancellous bone (0.3GPa) in the vertebra and PMMA cement provide very limited bone ingrowth at the implantation site
123. Further, concerns about potentially serious patient care issues involved with the use
of PMMA bone cement in the spine has been expressed by The U.S. Food and Drug
Administration (FDA) 124, 125. Possible complications of PMMA bone cement include
leaching out of unreacted MMA monomer in the body, leaking of PMMA bone cement
into the local biological environment, heating the tissue during in situ PMMA
polymerization and triggering mechanical failure at the bone-cement interface. 26
The compressive biomechanics of vertebroplasty have been investigated by several groups on the effect of the surgical procedure and cement type 118, 126, 127 as well as volume of cement injected 41. Sun et al. examined the effect of volume of PMMA
cement on the load to failure and stiffness, 5 %, 10 and 20% volume of PMMA was
injected into vertebral bodies. Results showed that injecting with 20% by volume of
PMMA resulted in an increase compressive strength of 36% over control vertebrae,
although these vertebrae experienced frequent extraversion of the cement through vascular channels or the spinal canal 128. The authors remarked that the vertebral stiffness
and strength increases were directly proportional to PMMA filler volume. Since the in
situ volumetric shrinkage of PMMA was recently reported in prosthesis systems that can
initiate cracks in the implant systems 129, similar problems may also associate with the
injectable PMMA cement in vertebroplasty.
2.6.2. Ceramic Injectable Materials
Several in vitro studies with injectable calcium phosphate cements show their
feasibility and mechanical effectiveness. Animal studies confirm their biocompatibility
and osteoconductivity 122, 130-136. An example of these bioactive injectable cements is the
commercially available self-hardening calcium phosphate composites, Norian® 62, 137, consisting of equimolar amounts of tetracalcium phosphate and dicalcium phosphate anhydrous, hardens when mixed with water and forms a resorbable hydroxyapatite (HA) as the end-product. This material provides excellent mechanical properties and direct bone ingrowth ability, but it set in situ after injection. The hardened ceramic materials often contain porous and lead to poor cement-bone interfacial mechanical properties 138,
139. Also this cement usually is in a form as a paste and would be harden within a short 27 period of time, which is difficult to inject in site 132, 140. Thus, with clinical complications
of these ceramic cements currently preclude their clinical use. Belkoff et al. group
showed that pre-fractured vertebrae injected with either Cranioplastic (PMMA) or
Bonesource (HA). The PMMA cement restored the `strength` (load to failure) of the
vertebral body beyond that of the pre-fractured state, while the HA cement restored the load to failure of the vertebral body to that of the pre-fractured state. Neither PMMA nor
HA cement was successful in restoring the stiffness of the pre-fractured vertebrae.
Differences were also found in the ability to restore strength in the thoracic versus lumbar regions. In a separate study, An’s group compared PMMA cement to a calcium phosphate cement. They likewise found that PMMA injected after fracture resulted in higher failure loads than the controls while injection of the vertebral bodies with CaP cement restored the failure loads to the level of the controls. The mean stiffness of the
CaP cement vertebroplasty group was considerably smaller than the PMMA group and both were greatly reduced from the controls (25% and 50%, respectively).
2.6.3. Injectable Composite Materials
Several research groups showed that bone-like apatite could be formed on/in a three-dimensional, structured organic hydrogel matrices at normal temperatures and pressure in vitro using a novel alternate soaking process for a relatively short time. It was shown that three-dimensional hydrogel–apatite composites were easily prepared using this alternate soaking process 138, 141. For instance, the Taguchi research group focused on
apatite coating on hydrophilic polymer grafted polyethylene films 142, and other
commonly used hydrophilic polymers, such as, poly(acryl amide) (PAAm) and 28 poly(acrylic acid) (PAAc) were also employed 143. However, in general, these systems are generally developed as cement or as scaffold form of bone replacements.
Weiss et al. has been developed an injectable bone substitute combining a polymeric water solution (such as, hydroxypropylmethylcellulose and water) viscous phase with bioactive calcium phosphate ceramic granules 138, 144-147. Comparing to ionic
polymers (such as PMMA), a nonionic hydrophilic polymer matrix would lead to a
negligible in vitro volumetric contraction of the composite matrix from the spontaneous
ionic reticulation of the bioactive granules in the simulated body fluid. However, they are
just mixing the ceramic powder with the polymeric solution, which allow very limited
interaction between two phases. Therefore, the interfacial mechanical properties between
the polymer matrix and the calcium phosphate fillers would be minimal. Today they are
working on the chemical improvement of the polymer phase of the system in order to
enhance the mechanically properties of the injected material. 147, 148.
One composite cement of bioactive ceramics with bisphenol-a-glycidyl
methacrylate (bis-GMA), and triethylene glycol dimethacrylate (TEGDMA) has been
developed as an injectable materials for vertebral fixtuation, named Cortoss®. This
cement form apatite on it surface in the body environment, and bond to the bone through
this apatite layer 118. Unpolymerized monomers, however, dissolve from the surface of
these cements and a less dense layer forms under the apatite layer 149 . This less dense
layer may be unstable, and result in the bonding between these cements and the bone
being liable to degradation over a long period. In any case, the dissolution of monomers
from these cements is not desirable. 29
Table 2.1: Mechanical properties of cortical and trabecular bone 26, 27, 30.
Elastic Flexural Flexural Density Compressive Bone Type Modulus Strength Modulus [116] Strength [MPa] [GPa] [MPa] [GPa] Cortical Bone 1.8 - 2.2 170 7 - 30 180 20 Trabecular Bone 0.3 - 1.9 40 - 60 0.05 - 0.5 - 2 Adult mandible 1.8 - 2.0 90 -140 2 - 10 - 5 - 12 Adult vertebra 0.3 – 0.6 - 0.1 - 0.25 - - 30
Figure 2-1: Bone Physiology. Courtesy Gray's Anatomy 35th edits Longman Edinburgh 1973. 31
Figure 2-2: An illustration of the molecular structure of Poly(glycolide-co-lactide). 32
Figure 2-3: Chemical formula of the monomer of PMMA. 33
3. Research Goals
In 1980s, the main research stream of bone substitutes was combining bioactive ceramics with various polymers. These composites were intended to provide a wide range of mechanical properties and to mimic natural bone characteristics. This group of composites can lead to natural bone ingrowth resulting in biological fixation through osseo-conduction process. Unfortunately, these bioactive ceramic filled polymer composites tend to lose strength rapidly when exposed to the aqueous physiological environment. Failures occur mainly at the interface of the bioactive ceramic filler / polymer. In this work, we investigated two alternative methods in reducing the interfacial failure of bioactive ceramic reinforced polymer composites. Our first approach involved using coupling agent treatment to enhance the interfacial properties of a bioactive ceramic (hydroxyapatite) / polymer (polymethylmethacrylate) solid composite. We hypothesize that by enhancing the interfacial bonding between hydroxyapatite particles and polymethylmethacrylate matrix, the local interfacial properties will be improved and the in vitro mechanical properties will be better retained. Further, by introducing the local interfacial mechanical behavior between the particles and the matrix we can predict global mechanical properties of the composite. The specific aims are:
1. To determine the in vitro local interfacial elastic properties of treated and untreated
HA/PMMA composites.
2. To evaluate the in vitro bending properties of treated and untreated HA/PMMA
composites.
3. To understand the surface bioactivity of the treated and untreated HA/PMMA
composites as a function of immersion times in vitro. 34
Our second approach is to molecularly engineer bioactive ceramic fillers (calcium salts) into the backbone of a co-polymer. We believe that incorporating silane with the thermoreversible hydrogel poly (N-isopropylacrylamide - ethylene glycol) will result in an organic-inorganic hybrid that will be bioactive and have enhanced mechanical properties over silane-free treated controls. The specific aims are:
1. To characterize polymer molecular interactions, liquid to solid phase transformation
temperature and phase transition behavior as a function of cross-linking agent
compositions.
2. To investigate the effect on volume change, in vitro mechanical properties and
bioactivity for different concentrations of silane in protein-free and protein-containing
biological fluid in a range of immersion times.
35
4. An Solid Bioactive Polymeric Composites: HA/PMMA
4.1. Introduction
During the last two decades, bioactive ceramics, such as hydroxyapatite, have became widely examined for use in dental implants, such as mandible replacements 67, 68,
150-153. The bioactive properties of these ceramics have received much attention in the field of hard tissue replacement. Bioactive ceramics provide fixation to bone through biological ingrowth of the local tissue through a series of chemical reactions between the ceramic surface and the bone 67, 96, 154. One way to predict the in vivo bioactivity of a
material is to determine the in vitro performance of the specimen by immersing the
specimen in a simulated body fluid (SBF) and determing the temporal ion dissolution and
precipitation. The hydroxyapatite layer formation in vitro corresponds to bone bonding in vivo 154-159.
The local mechanical properties of many synthetic biomaterials has been analyzed
using nano-indentation 97, 160-163, which is a depth-sensing technique that can accurately
characterize the mechanical properties of almost all types of solid materials at a small
scale 97-100. If plastic deformation occurs, then there is a residual impression left in the
surface of the specimen. In some cases, the permanent deformation or residual impression
is not the result of plastic flow but may involve cracking, phase changes or structural
changes within the specimen. Several groups have conducted nano-indentation on
ceramics to determine hardness and Young’s modulus from the analysis of the load-
displacement unloading curves 100, 160, 164-167. They found that both Young’s modulus and
hardness exhibit significant peak-load-dependence. Further, indentation size effect (ISE), 36 while a concern in the analysis of metals 97, is still under debate in the testing of
ceramics. ISE may influence the calculation of the hardness and Young’s modulus, which
is determined in part from the estimation of the indentation projected area. However, none of these studies have fully analyzed the characteristics of the resulting nano-
indentation marks. The imprint of the indentations provides details on whether the
indentations fracture or behave differently with various peak-loads.
The indentation tests yield information that involve localized, non-uniform
deformation or point-contacts, such as dental occlusal contacts with surface asperities
during chewing and wear 99, 161, 168, 169. Indentation can quantify the local mechanical
properties on small surface areas, for example, at a ceramic-polymer interfaces of a
ceramic-filled polymer composite, which is difficult to define under global mechanical
tests 34, 101-103, 170. The morphological characteristics of the resulting nano-indentations
and the resulting fracture mechanism can be evaluated from the magnified images of the
indentations using scanning electron microscopy 35. Since a major failure mechanism of a
two-phase composite is the interfacial de-bonding, the interfacial mechanics of the
composite leads to better understanding of the global mechanical behavior 96, 171-173. Xu et
al. 161, 168, 169 applied nano-indentation to the investigation of the effect of different
whiskers on the reinforcement of dental resin composites. In these studies, the materials
experienced significant changes in hardness and elastic modulus values as the indenter
shifted from matrix to interface and from interface to fillers. While this study gave new
insights into interfacial mechanics, additional information on indentation size, shape and
local fracture mechanism in the system might further help to elucidate the interfacial 37 behavior. Indentation sizes are on the order of a micron, which may provide challenges in optically imaging the indentations at the interface.
In our research, we investigated a composite material of hydroxyapatite (HA) and polymethylmethacrylate (PMMA) for mandibular augmentation 174. We examined
methods to enhance the interfacial bonding between the bioactive ceramic particles and
the polymer matrix. One approach to promote the interfacial bonding of a bioactive
ceramic/polymer composite is by introducing a coupling agent that is capable of
advancing chemical bonding between the ceramic and polymer of the composite. The
modified interface is intended to reduce debonding at the filler-matrix interface leading to
a stronger, higher modulus material, especially after immersing in an aqueous solution
175-177. We hypothesize that by enhancing the interfacial bonding between the HA and
PMMA, the in vitro mechanical properties will be retained without compromise to the bioactivity of the material. Further, we put forth that the local mechanical behavior of the ceramic/polymer interface will predict the global mechanical properties of the composite.
The aim of this study was to determine the most effective coupling agent for promoting the global bending and local elastic mechanical properties of the HA/PMMA composites.
Also to examine the in vitro elastic mechanical properties and bioactivity reaction
kinetics of a hydroxyapatite (HA) particulate reinforced polymethylmethacrylate
(PMMA) composite with and without the selected coupling agent. 38
4.2. Materials and Methods
4.2.1. Material Processing
We developed hydroxyapatite (HA) / polymethylmethacrylate (PMMA) composite systems with volume ratio of 55:45 as dental implant materials. The HA scaffolds were fabricated through a patented three dimensional printing process (3DP) called Theriform™. Scaffolds were produced from HA powders of 20±5µm average particle size (Ceramed Dental LLC, Lakewood, CO). A binder was added to bind the particles for better control of the porosity. This process provides a multi-layer stacking feature that can apply to complex shape design with thickness variations. The resulting scaffolds had dimensions of 10mm by 10mm by 4mm. Subsequently, the HA scaffolds were sintered at 1250°C for 1 hour and the polymer binder was melted. Pure HA blocks were then infiltrated with molecular weight 120,000g/mol methylmethacrylate monomer
(MMA) with benztroyl as the initiator. The PMMA added blocks were pressurized in vacuum-sealed PVA bags at 60psi, and were followed with a curing process at 60°C then at 120°C in a stepwise fashion.
The sintered pure HA scaffolds were immersed with one of the three coupling agent solutions (see Table 4.1) prior to PMMA infiltration and curing process. The uncoupled composites served as controls. The specimens were polished with 1µm and
0.5µm wet alumina polishing paste, to an optical finish. The in vitro specimens were subsequently immersed in a simulated body fluid (SBF) at 37°C for immersion periods from 1 to 216 hours. The simulated body fluid (SBF) contains ions to represent the ionic
2+ 2- concentrations of plasma: 2.6mM Ca as CaCl2, 1mM HPO4 as K2HPO4·3H2O, 152mM 39
+ 2+ + 2+ Na as NaCl, 135mM Cl as CaCl2, 5mM K as KCl, 1.5mM Mg as MgCl2·6H2O,
3- 2- 157, 158 27mM HCO as NaHCO3, and 0.5mM SO4 . The chemicals were dissolved in
deionized water and buffered to pH 7.6 at 37ºC with 1.0mN Tris-HCl (Sigma, St. Louis,
MO) 155, 158. After immersion, each specimen was rinsed in ethanol and was ultrasonically
cleaned in deionized water for 30 minutes 157, 178 . Prior to nano-indentation, all specimen
surfaces were marked with three 2mm by 2mm grids using a diamond scriber to indicate
the test position of each nano-indentation array.
4.2.2. Bending Tests
Four point bending was performed on the controls and each of the three treatment
groups to determine the global bending strength and elastic modulus of the HA/PMMA
composite. This test was selected to simulate, in part, the load bearing action of a
reconstructed mandible during mastication. The four point bending test setup was
prepared in accordance with ASTM C1161-95. Testing was performed on an Instron
mechanical test machine Model 1331 (Canton, MA) using a constant load rate of
2.5mm/min until failure occurred. Load-displacement (N/mm) curves were generated on
a data-acquisition device of the testing machine. The deflection value was estimated from
the displacement value at the failure load.
4.2.3. Compression Test
Compression testing of the HA/PMMA composites to failure was performed
using an Instron mechanical test machine Model 1331(Canton, MA) operating at a
constant crosshead speed of 3mm/min. The dimension of each composite block was 4mm x 4mm cross sectional area with 2.5mm thickness. Prior to testing, specimens were dried 40 in air for 1 hour after 24 hours immersion in SBF. Load-displacement (N/mm) curves were generated on a data-acquisition device of the Instron mechanical testing machine.
The maximum load sustained by each specimen was determined as the peak force at failure during the test. The ultimate compressive strength of each specimen was calculated by dividing the maximum load by its cross-sectional area. The Young’s moduli of the composites were calculated using the slope of the initial elastic region of the stress-strain curves.
4.2.4. Physicochemical Analysis
Prior to mechanical analyses, the chemical composition and surface morphology of the unimmersed and immersed composite surfaces were analyzed using a field emission environmental scanning electron microscope (FE-SEM) equipped with
-9 Electron-Dispersive Spectroscopy (EDS, Philips XL series, Netherlands) at 1.4 x 10
mtorr vacuum level and 10keV voltage. Based on previous literature, the formation of a
bone bonding material surface is achieved through the formation of an apatite layer in
vitro, where an apatite layer with a calcium-to-phosphate ratio higher than 1.4 is known
to promote direct bone bonding 156, 158, 159. The Ca/P atomic ratios of three 200 µm2
surface areas of a composite were analyzed under EDS to estimate the average surface
Ca/P atomic ratio of a specimen. The in vitro specimens were immersed in SBF for 3
hours to 24 weeks. After immersions, solutions were analyzed for calcium and phosphate
ion concentrations. Calcium concentrations of solutions were measured by flame atomic
absorption light spectroscopy (AAS) (Perkin-Elmer, models 2380 and 5100 PC, Norwich,
CT), while phosphate concentrations in solution was measured as a phosphomolybdate 41 complex (molybdenum yellow) using UV-visible spectrophotometer (Biochrom, LKB
4053 Ultrospec K, Cambridge, England) 179.
4.2.5. Nano-indentation
Nano-indentation was performed using a Nanoindenter XP® (MTS Systems). In general, indentation was made on samples by applying a load and measuring displacements utilizing a Berkovich indenter in 7x3 arrays. As shown in Figure 1, the indent tip penetrates the surface as the applied load increases. After indentation, the specimens were coated with less than 0.01µm of gold. Using a field emission scanning electron microscope (FE-SEM) (Philips XL series, Netherlands), we reviewed the shape and morphology of the nano-indentation marks up to 5000 magnification at 1.4x10-9mtorr
vacuum level and 10keV voltage. Based on this approach, we defined accurately the
position of each impression and the details of each crack generated by the indentation.
Also, we correlated the collected measurements of Young’s modulus and hardness from the output of the indentation with the indent mark details.
Using Oliver-Pharr method 180, the three key parameters needed to determine the
hardness (H) and the elastic modulus (E) are the peak load (Pmax), contact area (A), which
is dependent on the indentation penetration depth at peak load (hc), and the initial
unloading contact stiffness (S). These parameters are measured based on the load-
displacement data. The contact stiffness (S) is measured only at the initial unloading and no restrictions are placed on the unloading data being linear during any portion of the unloading. The fundamental relations applied to determine H and E are:
P H = max A Equation (1) 42
where P is the load and A is the projected contact area at that load, and:
S π Er = 2β A Equation (2)
where Er is the reduced elastic modulus and β is a correction factor that depends on the geometry of the indenter. The projected contact area (A) is a function of the contact depth (hc) at the elastic region during unloading, see Equation (3). For the
Berkovich tip, β is 1.034 and θ is 65.3°.
A = 3 3h 2 tan 2 θ c Equation (3)
The elastic modulus of the test materials, E, is calcaulted from Er using:
1 1−ν 2 1−ν 2 = + i E E E r i Equation (4)
where υ is the Poisson’s ratio of the test material, and Ei and υi are the elastic modulus and Poisson’s ratio of the indentation tips, respectively. For a high temperature sintered hydroxyapatite, υ is about 0.25 150, and for a diamond indentation tip (as used in
97 this study), υi = 0.07 and Ei is 1140GPa .
Three studies of the local mechanical properties of the HA/PMMA composite were performed using nano-indentation. In the first study, we examined the effect of load level on H and E of the HA particulate surface of the composite because both H and E have been observed to be dependent on load level in nano-indentation 97, 165. Loads of 1,
3, 5, 8, 10, 15, 20, 25, 50, 75 and 100mN were applied on the HA particulate surfaces. 43
Load-displacement data was collected. FE-SEM images were analyzed for the resulting indentation marks. H and E were calculated according to the Oliver-Pharr Method.
The second analysis employed a load at the optimal range (10mN) to create several 3x7 arrays of the indentations on the unimmersed controls and specimens from each of three treatment groups. In the third study, the in vitro local interfacial mechanical properties of the SBF immersed controls and PMMA-MAA treated composites were determined from several 3x7 indentation arrays at the optimal load. All load- displacement data were correlated to the FE-SEM image of the indent mark. E and H were calculated for positions on the HA particles, “at the interface” (defined as 2µm from the interface) and on the PMMA surface of the composite.
4.2.6. Data Analysis
In the bending, compression and nano-indentation tests, 5 specimens of each group (controls and treated composites) were tested. One-way ANOVA analyses were performed for all mechanical properties results at 95% confidence level (p values less than 0.05).
4.3. Results
4.3.1. Global Bending Behavior
The global bending properties of the HA/PMMA composites with and without coupling agents were examined, see Table 4.2. The test results revealed that the bending strengths of both co-polymer coupled specimens (PMMA-MAA and PMMA-MA) are statistically greater than the controls. The bending strength of PMMA-MAA treated group is 27% higher than the controls with p=0.001, while the bending strength of 44
PMMA-MA treated group is 18% higher with p=0.004. However, there was no statistical significance in bending properties between the two co-polymer treated groups (p=0.06).
Neither bending strength nor modulus of silane coupled composites showed significant difference from the controls (p=0.69 and p=0.89, respectively). With the addition of
PMMA-MAA coupling agent (PMMA-MAA), the bending modulus of the composites obtained a 15% increase in comparison with the untreated controls (p=0.02). But the other co-polymer added system, PMMA-MA coupled composites, displayed no statistically significant from the controls (p=0.19).
4.3.2. Local Mechanical Properties of HA
The calculated Young’s modulus and hardness of the hydroxyapatite from the unloading curves are plotted as a function of the peak loads in Figures 4.2 and 4.3, respectively. The local mechanical properties decreased with increasing peak load up to
8mN. At the 10mN to 20mN load range, these properties became relatively constant. At
25mN and greater peak loads, a dramatic reduction was seen in both E and H. Using FE-
SEM, we were able to clearly image the indentations and linked the test outputs to the proper material phase. At 10mN, the nano-indenter left a permanent impression onto the ceramic without inducing any fracture, see Figure 4.2a. However, load level of 25mN and higher resulted in cracking of the HA particle, see Figure 4.2b. Therefore, the remainder of the testing was conducted at 10mN.
4.3.3. Local Interfacial Mechanical Properties
Table 4.3 displays the Young’s modulus and hardness values from 5 indentations at positions in the HA particles and at the HA/PMMA interface. The typical load- 45 displacement graphs of nano-indentations at the matrix (PMMA), the interface
(HA/PMMA) and the ceramic (HA) of the control composite are shown in Figure 4.4, along with the corresponding FE-SEM images (Figure 4.4a, 4.4b and 4.4c). All indents have a peak-load of 10mN. The load-displacement curves in the ceramic showed a much steeper slope upon loading than the ones at the interface and in PMMA. Upon unloading, the interface curve portrayed a rather polymer-like response. The polymer load- displacement curve was even less steep than the interface one and it obtained more plastic deformation.
Figure 4.5 showed an increase in the Young’s modulus of the co-polymer coupled
(PMMA-MA and PMMA-MAA) HA/PMMA composite from the uncoupled composite
(controls), 44% increment (p=0.01) and 50% increment (p=0.001), respectively; although they were not statistically different from each other. Also, there was no significant difference between the Young’s modulus of silane treated composites and the controls.
But, silane treated group exhibited a 38% reduction (p=0.002) in the local hardness compared to the controls. The two co-polymer coupling agents demonstrated no statistical improvement on the local hardness of the system, Figure 4.6.
In the normalized curves, Figure 4.7, both the global bending and local mechanical results indicated that the silane coupling agent does not enhance the bending modulus or local Young’s modulus of the composite compared the uncoupled controls (in fact, the properties are slightly reduced). However, both co-polymer coupling agent treatments improve these properties significantly with more pronounced increases detected in the local measurements. 46
4.3.4. Surface Bioactivity
Based on the global bending test, the PMMA-MAA coupled composites exhibited the highest bending modulus. Therefore, the in vitro mechanical behavior and bioactivity studies were only concentrated on the controls and PMMA-MAA coupled HA/PMMA composites as a function of immersion times. Figure 4.8 showed the Ca/P ratio of five different areas on the in vitro surfaces of the untreated and treated composite groups.
These Ca/P ratios exhibited no statistical difference (p= 0.44), while all values are well above 1.4. As shown in Figure 4.9, the morphology of the in vitro HA surface appeared rougher than the unimmersed state and fractured after 2 weeks of immersion, which was resulted from the in vitro surface reactions.
Figure 4.10 and 4.11 represent the calcium and phosphate concentrations in SBF as a function of immersion times. The observed induction time prior to a detectable chemistry change in solution was 1 hour and 3 hours for controls and treated composites, respectively. After 24 hours, the reaction proceeded at a slower rate by virtue of approaching the solid/solution equilibrium stage. The treated composites had a higher average ion uptake rate from 3 to 24 hours immersion than from 24 to 72 hours immersion. The treated composites have lower concentration of calcium and phosphate uptake than the controls between 1 to 24 hours immersion. However, the controls and the treated composites experienced similar calcium and phosphate concentrations in solution after 72 hours immersion with less than 5% difference.
4.3.5. In vitro Global Elastic Mechanical Behavior
Figure 4.12 and 4.13 showed the compressive test output. Before immersing in
SBF, the PMMA-MAA treated composites have statistical significant improvements in 47
Young’s modulus and in ultimate compressive strength of 24% (p=0.002) and 15%
(p=0.005) over the controls, respectively. After 24 hours immersion in SBF, the treated composites exhibited a 25% increase in Young’s modulus (p=0.003) over the controls but showed no statistical difference in ultimate compressive strength among the two specimen groups.
4.3.6. In vitro Interfacial Mechanical Properties
The typical load-displacement curves of the control and treated composite groups at the HA/PMMA interfacial position in the two composite groups before and after 72 hours immersion are displayed in Figures 4.14 and 4.15, respectively. The in vitro load- displacement curves of both immersed specimen groups (controls and treated) exhibited much shallower slopes than the un-immersed specimens, indicating a reduction in mechanical properties.
Figures 4.16 and 4.17 indicated the in vitro interfacial Young’s modulus and hardness of the HA particles and HA/PMMA interfaces as a function of immersion times.
After 3 hours immersion, the interfacial Young’s modulus of both groups showed no significant change from unimmersed samples (p=0.5). From 3 to 216 hours, the interfacial Young’s modulus of the controls decreased gradually but not significantly
(p=0.18) while their interfacial hardness revealed a statistically significant reduction of
35% after 3 hours immersion (p=0.005). However, upon longer immersion times, the local hardness of the controls remained constant (p=0.68) from 24 to 216 hours immersion.
Both interfacial Young’s modulus and hardness of the treated composites exhibited a significant reduction of 40% after 72 hours immersion with p=0.009 and 48 p=0.008, respectively. These mechanical properties remained unchanged after further immersion (p=0.24). Prior to immersion, there was a statistical difference of p=0.0008 among the interfacial Young’s modulus of the two composite groups, while there is no significant difference of their interfacial hardness. After 3 and 24 hours immersion, the interfacial Young’s modulus and hardness between the two composite groups exhibited statistical differences at p < 0.05. The local Young’s modulus of the treated group was
50% higher than the controls after 3 hours immersion and was 47% higher after 24 hours immersion. After 72 and 216 hours immersion, the interfacial mechanical properties of both composite groups demonstrated no statistical significant with p values above 0.05.
4.4. Discussion
Nano-indentation has been applied to analyze the local modulus and hardness of fiber reinforced bioactive ceramic/polymer composites. This test can accurately generate small dynamic loads (10mN) while providing information on the degree of energy absorbed, the Young’s modulus and hardness of the small contact surface on a composite.
In this study, we utilized nano-indentation to evaluate the local hardness and Young’s modulus at the interface of HA-PMMA composites as a function of coupling agents. We focused on the behavior of HA in order to determine the indentation size effect, the degree of cracking and the influence of the polymer matrix as a function of peak loads.
From Figures 4.2 and 4.3, we saw that HA had a relatively constant indentation response as determined by hardness and Young’s modulus in the 10mN to 20mN load range.
However, the local Young’s modulus and hardness of HA at less than 10mN were moderately higher and varied with the changing load levels (from 1mN to 8mN). These over-estimations of the mechanical properties may be attributed to the effect of the 49 indentation tip rounding and the presence of friction between the indenter and the specimen at small applied loads 97, 181. These effects likely provided less indentation contact depth and led to an under-estimation of the projected area of the indentation, resulting incorrectly in increased Young’s modulus and hardness at this load range.
At applied loads of 25mN and above, the dramatic reduction in the local mechanical properties were resulted from the fracture of the HA particulates, as shown in
Figure 4.2b. The PMMA matrix of the composite had a much lower Young’s modulus and hardness, which may limit the HA crack elongation. Also when cracking occurs in
HA, the unloading curve may shift to a much shallower slope with the consequence that
Young’s modulus and hardness are under-represented. The applied loading levels from
10 to 20mN produced constant Young’s modulus and hardness values. In addition, FE-
SEM images of the indentation from the HA particulates within the composites loaded to these levels resulted in no fracture of the surface and an acceptable indentation morphology 161, 182 in that there was no significant indentation size effect, see Figure
4.2a. One key limitation of determining the local mechanical properties of a composite system is the influence of the substrate material beneath the reinforcement. At a 2µm distance from the interface, on the particle surface, the estimated depth of the particle is about 2µm. In our nano-indentation, the contact depth of each “at interface” indentation at 10mN is about 0.5µm. The resulting indentation depth at the interface is, therefore about 20% of the depth of the particle, which should not result in artifacts due to the substrate material.
Depending upon the method of synthesis and sintering temperature, earlier reports revealed that the 1200°C sintered HA yielded a Young’s modulus of about 80 to 100GPa 50 and Vickers microhardness of about 6GPa 68, 150, 163, 170, 183, 184. Our nano-indentation determined that the Young’s modulus and hardness of HA at 10mN has an average value of 98.6GPa and 7.20GPa, respectively, which is in agreement with the published literature. Hence, we can apply the nano-indentation technique to analyze the effect of the coupling agent on the interfacial mechanical properties of the HA/PMMA composite with confidence.
At 10mN peak load, we found that the approximate area of the indent that is located in the center of the HA particle is about 1.3µm2. When the indentation is closer to the HA/PMMA interface, its projected area may be increased to 4µm2 with about 2µm breadth. Thus, we defined any indentation from 1.8µm to 2.2µm away from the
HA/PMMA interface of the 18µm to 22µm diameter HA particulates as the “at interface” indentation. It is especially important to correctly estimate the size of the indentations in order to properly space each indentation in the array with limited influence of the plastic zone effect from the adjacent indentations. Based on our FE-SEM images, we placed each indentation 15µm apart.
The estimation of the Young’s modulus and hardness based on the Oliver-Pharr method was valid because our observation of the load-displacement curves and indentation morphology are reproducible and without indentation size effects.
Unfortunately, Young’s modulus and hardness of the PMMA in this composite system were unacceptable because the nano-indentation was influenced by the embedded ceramic particles (Figure 4.4c). Thus, the resulting unloading curve was not indicative of that for pure PMMA. However, we determined the Young’s modulus and hardness values of the ceramic- polymer interface via Oliver-Pharr methods was adequate after our 51 observation of the load-displacement curves and indentation morphology. As shown in
Figure 4b, the FE-SEM photos of 5000 magnification of the indentation at the interface, the 10mN load nano-indentation led to a permanent impression while producing no cracking in the brittle material. The projected area of this “at interface” indentation was approximately 4µm2. This area can be compared to the calculated area for the same indentation via Equation (3), based on penetration depth, hc, and Berkovich tip geometry, which was 5µm2. Based on Equation (1), the calculated hardness would be 2.5GPa for the
4µm2 area and 2GPa for the 5µm2 area. The difference in the determination of the projected area was influenced by the PMMA matrix on the indentation tip during its penetration. The displacement of the tip penetration was clearly recorded in the load- displacement curve.
The nano-indentation studies showed the slopes of the elastic-plastic curves of coupled materials were altered by the different chemistries of the coupling agents. It appeared that the silane coupled composite had the lowest local Young’s modulus and hardness among the treated groups. Silane coupling agents are commonly used as adhesive promoters for bioactive ceramic reinforced composites 96, 172. Results in previous studies showed improved mechanical strength with no change in the Young’s modulus of silanized HA reinforced polymeric composites. For this HA/PMMA composite, the silane coupling agent did not provide a mechanical benefit over the un- coupled control specimens. The –OH- molecules in the HA particles were expected to bond well with the –OH- molecules in the silane coupling agent. But the presence of water and oxygen in the atmosphere likely led to the formation of addition water molecules with the –OH- groups in the HA and left the –OH- groups in the silane 52 unattached. Our data did indicate that the co-polymer coupled composites (PMMA-MA and PMMA-MAA) significantly increased the local Young’s modulus of the HA/PMMA composite at the interface. However, the local hardness of the co-polymer coupled composites did not show statistical difference from the uncoupled composite. The co- polymer coupled composites provided enhanced chemical interactions between the HA particulates and the PMMA matrix. While the global bending modulus increased compared to that of the control group, the enhancements were more pronounced on the local level, as indicated by nano-indentation.
The starting hydroxyapatite powder has a Ca/P molar ratio range from 1.50 to
1.67. Based on EDS analyses, the HA/PMMA specimens of both groups had fairly similar Ca/P ratio of above 1.4 for all tested in vitro immersion periods. After 9 days immersion, the average Ca/P ratios of both controls and treated composites approached to a constant value of 1.67, as shown in Figure 4.8. There was no statistical difference among the Ca/P ratios of the controls and PMMA-MAA co-polymer coupled HA/PMMA composites, while the calcium and phosphate uptake was virtually the same concentration for each condition after 72 hours immersion. The in vitro data indicated HA/PMMA composite with and without the co-polymer coupling agent can achieve natural bone bonding and have an average surface Ca/P ratio that is close to the 1.67 Ca/P ratio of natural hydroxyapatite.
Before and after 24 hours immersion, the advantage of the PMMA-MAA coupling agent on the global Young’s modulus in comparison with the untreated controls was significant. On the other hand, the effect of 24 hours immersion in SBF on the
Young’s modulus of controls and treated composites was insignificant. However, after 24 53 hours, surface bioactivity of the HA particles led to a significant alteration on the ultimate compressive strength of both composite groups (p=0.001 in both cases). The major failure mechanism of the immersed and unimmersed specimens was HA-PMMA interfacial debonding, which was determined from FE-SEM imaging. This failure mechanism is similar to that of many other filler reinforced composites 96, 172, 185.
Using nano-indentation, the in vitro interfacial mechanical properties of the controls and treated HA/PMMA composites were determined. The shape of the load- displacement curve is often found to be a rich source of information, not only for providing a conformation to calculate Young’s modulus and hardness of the biomaterial, but also for the identification of non-linear events. The surface bioactivity of the HA particles resulted in lower interfacial Young’s modulus and hardness of the immersed specimens in the region of the interface.
The effect of surface roughness increases with increasing radius of indenter and increases with decreasing indenter load. Thus, for small loads with spherical indenters, surface roughness can have a significant effect on the test output 97, 186. For sharper indenters (such as Berkovich indenter with a tip radius of 100nm), the effect of surface roughness is less dramatic 180. This sharp indenter led to an optimization of the effect of surface roughness under small applied load. During our observation of the HA surface bioactivity, some HA particulates began to chip after 240 hours immersion and its surface would be increasingly coarse with the immersion time. Therefore, we were not able to determine the local interfacial mechanical properties after 240 hours immersion using nano-indentation. Also, drying the composite after immersion may have led to 54 dehydration of the ceramic and polymer matrix, which likely shrunk and may have disrupted the interface.
Although the 24 hours immersed composites exhibited less reduction of local mechanical properties in comparison to the global compressive properties, there were statistical differences between the in vitro interfacial Young’s modulus and hardness of the controls and treated composites (p=0.01 and p=0.001, respectively). After 72 hours immersion, the interfacial mechanical properties of the treated composites behaved similarly to the controls. We put forth that once the HA/PMMA composites are in contact with the physiological environment, the coupling agent delays the ion exchange process at the HA/PMMA interface, as shown in the analysis of calcium and phosphate uptake.
The treated composite had an induction time of 3 hours immersion as compared to the controls, which had an induction time of only 1 hour immersion.
In comparison to the controls, a normalized graph of global and interfacial
Young’s modulus of the PMMA-MAA treated composite before and after 24 hours immersion in SBF was plotted in Figure 4.8. Although both the global and interfacial
Young’s modulus showed a similar mechanical behavior after immersion in SBF for 24 hours, a more pronounced difference was detected at the HA/PMMA interface of treated and untreated composites using nano-indentation. Thus, the nano-indentation determined local interfacial mechanical properties are indicative of the influence of coupling agent treatments on the global mechanical behaviors of the HA/PMMA composite. The limitation of the technique, however, is that local mechanical properties determination does not take in account the homogeneous nature of the material, such as the effects of porosity and particle size distribution on the mechanical properties. Nonetheless, as a 55 tool to evaluate the mechanical response to alterations in reinforcing agent/matrix interfaces, nanoindentation has proved to be a valuable tool.
Further, the unimmersed global elastic and bending moduli of the PMMA-MAA treated HA/PMMA treated composites are comparable with the ones of the human mandible tissue (Table 4.4). When the composite is subjected to the body fluid the interfacial bonding become weak and catastrophic failures still may occur at the HA-
PMMA interfacial. Figure 4.9 showed calcified bridge would gradually form on non- bioactive PMMA matrix. However, it would not mechanically sufficient for mandibular augmentation while some HA particles began to fracture and mechanically separated from the PMMA matrix (as early as after 216 hours immersion). Therefore, it is crucial to improve the bonding mechanics between the bioactive phase and the polymer matrix.
One supposed method is to develop a macro-level single phase composite as hard tissue replacements.
4.5. Conclusion
The application of nano-indentation to the interface between the PMMA polymer and HA particulate reinforcement enabled us to obtain some details of the local mechanical behavior of the composite. Based on the test data of the unimmersed specimens, the PMMA-MAA and PMMA-MA coupling agents enhanced the interfacial stiffness while maintaining the interfacial hardness of the composite systems. Silane coupling agent showed a reduced local Young’s modulus and hardness compared to the uncoupled controls.
Applying nano-indentation at the interface enabled the elucidation of the local mechanical behavior of the material, which showed that the coupling agent enhanced the 56 in vitro surface interactions at the interface. The coupling agent enhanced the interfacial bonding for a short duration, however, after 72 hours, there was no difference between the mechanical behavior of the treated and untreated materials. Nano-indentation was useful in indicating the global mechanical consequence of interfacial treatment with a coupling agent even after reactivity in simulated body fluids. 57
Table 4.1: The chemical content of the four groups of composite systems.
Volume % in Total Group Chemistry Content Polymer Content
Controls No coupling agent -
Silane MPS in acetone 5 %
95vol% Polymethylmethacrylate: PMMA-MA co-polymer 2.5 % 5vol% maleic acid
95vol% Polymethylmethacrylate: PMMA-MAA co-polymer 2.5% 5vol% methacrylic acid
58
Table 4.2: The bending properties of various composite systems.
Bending Strength [MPa] Bend Modulus [GPa] Group ± SD (n = 5) ± SD (n = 5)
Controls 53.6 ± 4.8 15.5 ± 0.96
Silane 55.2 ± 7.2 15.3 ± + 2.75
PMMA-MA co-polymer 63.4 ± 2.9 16.4 ± 0.70
PMMA-MAA co-polymer 67.9 ± 3.7 17.9 ± 1.6
59
Table 4.3: Local Young’s modulus and hardness at various positions in uncoupled composite after 10mN loading.
Young’s modulus [GPa], Hardness [GPa], Nano-indentation position ± SD, (n=5) ± SD, (n=5)
HA particulate 97 ± 4.4 7.2 ± 0.41
HA/PMMA interface 31 ± 4.8 1.9 ± 0.27
PMMA matrix 7.4 ± 0.92 0.31 ± 0.04
60
Table 4.4: Comparison of the global mechanical properties of PMMA-MAA treated HA/PMMA composites and human mandible 27, 187-189.
Unimmersed (dry) After 24 Immersion (in vitro)
Elastic Bending Elastic Bending Modulus Modulus Modulus Modulus [GPa] [GPa] [GPa] [GPa]
PMMA-MAA treated 2.0 – 2.2 16.3 – 19.5 1.7 – 1.9 0.5 – 1.4 HA/PMMA composite
Human Mandible 2 -10 5 - 12 4.8 – 6.0 -
61
1µm
Indentation tip
Hydroxyapatite
20µm
Figure 4-1: A schematic representation of a nano-indentation applies onto a hydroxyapatite particulate upon loading. 62
(a) Indentation at 10µm from HA/PMMA interface.
160 (b) 140 Cracking 120
100
80 60
40 (GPa) Modulus Young's 20
0 0 20406080100120 Peak Load (mN)
Figure 4-2: Local Young’s modulus measurements at the center of hydroxyapatite as a function of applied loads; with (a): FE-SEM images of indented hydroxyapatite at 5000x after 10mN loading, and (b): FE-SEM images of cracks induced indentation in hydroxyapatite at 2500x after 25mN loading. 63
12
10
8
6
4
Local (GPa) Harndess 2
0 020406080100120 Peak Load (mN)
Figure 4-3: Local hardness measurements at the hydroxyapatite center of 3 to 100mN loading.
64
(a) (b) (c)
Figure 4-4: Typical nano-indentation curves in a ceramic-polymer composite; with FE-SEM image of a typical indentation mark (a) on a hydroxyapatite particulate at 5000x magnification, (b): at the hydroxyapatite- polymethylmethacrylate interface at 5000x magnification, and (c): on a polymethylmethacrylate matrix at 1200x magnification.
65
60
50 40
30
20
Young's Modulus (GPa) Modulus Young's 10
0 ControlCSMASilane PMMA:MAA PMMA:MA Coupling agents
Figure 4-5: Graphic representation of the effect of various coupling agents on the local Young’s modulus at the HA-PMMA interface.
66
3.0
2.5 2.0
1.5 1.0
Hardness (GPa) Hardness 0.5 0.0 Control Silane CSMAPMMA:MAA PMMA:MA Coupling agents
Figure 4-6: Graphic representation of the effect of various coupling agents on the local hardness at the HA-PMMA interface.
67
Figure 4-7: The local elastic modulus and global bending modulus of each coupling agent added composite group was normalized with the controls. 68
2.00 Controls 2.5% PMMA:MAA 1.90
1.80
1.70
Ca/P Ratio 1.60
1.50
1.40 0 50 100 150 200 Days
Figure 4-8: Ca/P ratios on a 100µm2 surface area of the controls and treated composites after immersion in SBF at various immersion times. 69
Figure 4-9: The FE-SEM/EDX images of the PMMA-MAA treated HA/PMMA composite after in vitro chemical reaction from 0 to 24 weeks. 70
3 Control Treated composite 2.5
2
1.5 1
0.5 0 calcium concentration [mmol] concentration calcium 0 1020304050607080 immersion time [hours]
Figure 4-10: Calcium concentration in SBF after immersion of controls and treated composites from 1 to 72 hours. 71
1.2 Control
Treated composite 1.0 0.8
0.6
0.4 0.2
0.0
phorphous concentration [mmol] 0 20406080
immersion time [hours]
Figure 4-11: Phosphate concentration in SBF after immersion of controls and treated composites from 1 to 72 hours. 72
2.5 Control Treated composite 2.0
1.5
1.0
0.5 Young's modulus [GPa]
0.0 Before immersion After 24 hr immersion
Figure 4-12: Global Young’s modulus of the controls and treated composites before and after 24 hours immersion in SBF. 73
Control 200 Treated composite
150
100 [M Pa]
50
0 Strength Compressive Ultimate Before immersion After 24 hr immersion
Figure 4-13: Global ultimate compressive strength of the controls and treated composites before and after 24 hours immersion in SBF.
74
Controls 12 Un-immersed 72 hours 10 immersed
8 6 4 Load [mN]
2 0 0 100 200 300 400 500 600 700 Displacement [nm]
Figure 4-14: Typical load-displacement curves of at the interface of the controls unimmersed and 72 hours immersed. 75
Treated composites 12 Un-immersed 72 hours immersed 10
8
6
Load [mN] 4
2
0 0 100 200 300 400 500 600 displacement [nm]
Figure 4-15: Typical load-displacement curves of at the interface of the PMMA- MAA coupled composites un-immersed and 72 hours immersed. 76
Figure 4-16: Interfacial Young’s modulus of the controls and treated composites as a function of immersion times. 77
Figure 4-17: Interfacial Hardness of the controls and treated composites as a function of immersion times.
78
Global Young's Modulus
1.60 (Compression test) Interfacial Young's Modulus (Nano-indentation) 1.50
1.40
1.30 1.20
Ratio to Controls Ratio 1.10
1.00 unimmersed after 24 hr immersed
Figure 4-18: Normalized the global and interfacial Young’s moduli of the treated composite with the controls’ properties. 79
5. An Alternative Approach to Bone Replacement
5.1. Introduction
The global and interfacial mechanical properties of the HA/PMMA composite with or without coupling agents were determined using compression, bending and nano- indentation tests. Improvements in the global and interfacial Young’s modulus were obtained for the PMMA-MAA treated systems. The prior reported data showed that the mechanical properties of these materials are comparable with the one of the natural mandible tissue. Once the treated composites are soaked in SBF, their advancements in mechanical properties would slowly diminish within 3 days and would behave similar to the untreated composites.
The in vitro mechanical and chemical bonding between the two phases of the
HA/PMMA composites (with or without treatments) were weak and would break down over time. Hence, we explored an alternative method to combine two distinct materials.
Other limitations of these composites are they require open surgical procedure and only provide a limited fit to the implantation site. Thus, we proposed injectable biomaterials to treat irregularly defects and can reduce the invasiveness of the implantation procedures.
We anticipated creating a “single phase” material using molecular design to engineer bioactive ceramic into the polymer backbone as an ideal bone replacement material that is injectable at room temperature and would become rigid in body. Also, without the presence of microscopic interfaces, the superior theoretical mechanical properties of the composite may be achieved. Based on previous lab work on PNIPAAm hydrogel as an injectable thermo-sensitive nucleus implant in the spinal intervertebral disc 1, we employed PNIPAAm as the primary materials as a novel injectable bone substitute for 80 vertebroplasty. In the preliminary experiments, several factors were screened in synthesizing the bioactive hydrogel. These factors include adding polyethyleneglycol, applying tri-methacryloxypropyltrimethoxysilane, using calcium chloride, selecting a proper polymerization time, and including a rigid polymeric component. Also the compressive mechanical properties and phase transformation behavior of these hydrogel systems were evaluated.
5.2. Hydrogels
Hydrogels are three-dimensional hydrophilic polymer networks imbibing large amounts of water or biological fluids. They are polymers that are glassy in the dry state, and can be either a homogeneous polymer (by itself) or a co-polymer (with two or more different polymers). They exhibit water-swelling ability, which means that they swell when they are in contact with water. Because of its swelling ability, hydrogel holds much promise for spine replacement surgeries as well as potential usage in tissue engineering applications.
Many hydrogels has recently developed as biomaterials for applications in the medical and pharmaceutical fields. Research studies have showed that the swelling behavior of hydrogels depend on the external environments 190, 191. They can exhibit abrupt changes in their swelling behavior of the network structure, permeability, or mechanical strength in response to changes in pH, ionic strength, temperature, and electromagnetic radiation. The most commonly studied hydrogels having environmental sensitivity are sensitive to either pH or temperature 143, 192. 81
5.3. Poly(N-isopropylacryamide)
Poly(N-isopropylacryamide) (PNIPAAm) is a thermosensitive hydrogel that has received high attention for biomedical use because of its lower critical solution temperature (LCST) behavior at around 32°C in an aqueous solution 193-196. PNIPAAm chains hydrate to form expanded structures in water when the solution temperature is below its LCST, but becomes a compact gel structure by dehydration when heated to a temperature above its LCST. Below its LCST, PNIPAAm is extremely soluble in water and appears transparent. However, as its temperature is increased above its LCST, it becomes hydrophobic from the increased interactions between the isopropyl groups and
PNIPAAm precipitates out from the aqueous solution and becomes opaque. PNIPAAm hydrogels possess a three-dimensional network structure, which is insoluble but has characteristics of reversible swelling 191. The polymer chains undergo a coil (soluble)- globule (insoluble) transition when the external temperature cycles across its LCST at about 33°C 190, 197. Thus, at a temperature below the LCST, PNIPAAm hydrogels absorb water and exist in a swollen state, but shrink and display an abrupt volume decrease when the environmental temperature is higher than the LCST. With fast response of swelling ability in the body temperature, PNIPAAM would be an excellent candidate for injectable tissue replacements.
Stile et al. 198 created loosely crosslinked polymers of PNIPAAm and PNIPAAm– co-acrylic acid as an injectable hydrogel for cartilage tissue applications. With the addition of acrylic acid (AAc), the LCST of polymer increased about 2°C to 34°C. The
PNIPAAm-co-AAc did not experience the same dramatic water loss as the PNIPAAm control. In fact, the initial water content at physiological temperature actually increased 82 slightly but decreased to about 74% after 6 days in PBS. These values become important in the context of in situ gelation since a polymeric plug needs to adequately fill the cartilage defect void. If too much water separates from the hydrogel, the polymer might not yield a tight or secure fit and consequently, would decrease the potential for tissue regeneration. In vitro experiments with bovine chondrocytes demonstrated that both polymers supported cell viability and allowed cells to produce extracellular matrix molecules. These initial results showed that hydrogels derived from PNIPAAm can provide a biocompatible matrix for applications such as cartilage repair.
Silane contains silanol side groups and has been applied as an adhesive agent for binding two materials. Through the incorporation a silane with PNIPAAm, a bioactive hydrogel can be engineered. Some silanol (Si-OH) in the silane provided attachment sites for calcium and initiated mineral dissolution from the simulated biological environments, while some may condense into the siloxane serving as a cross-linker to the polymer.
Traditional PNIPAAm-based hydrogel is usually formed through the reactions of
NIPAAm monomers and commercial cross-linking agents. With the addition of cross- linking agent, the hydrogel was mechanical strengthen. However, the thermosensitive behavior of the hydrogel can be altered from the chemical cross-linking that may prohibit its application as injectable implant materials. Liu and Sheardown 199 investigated
PNIPAAm-siloxane systems as ophthalmic biomaterials. They reported that the LCST phenomenon remained in these materials (but less abrupt). These siloxane added gels have greater mechanical strength than the pure PNIPAAm polymer because the siloxane serves as cross-linkers to the polymer backbone. 83
5.4. Preliminary Studies
5.4.1. PNIPAAm Co-polymer
Surface-immobilized and stimuli-response hydrogels gain increasing attention because of their potential as in vitro carriers for regenerative medicine. These hydrogels provides mechanical properties and elasticity in its swollen state, which is important for cell anchorage to the substrate. Besides these properties, the polymer matrix permits diffusion and delivery of nutrients and growth factors. With proper molecular improvement of the PNIPAAm hydrogel, a multifunctional polymer can be produced for hard or soft tissue repair applications.
At temperature above the LCST of PNIPAAm (about 32°C), the hydrogel expel its containing water and can loose its gel properties (such as elasticity). Kono et al. 200 demonstrated that with the addition of polyethyleneglycol (PEG), a hydrophilic component, enhances the elasticity and swellability of the gel network in body fluids.
However, PEG is known for its inert behavior toward biosystems in general and to protein adsorption in particular. Optimization of the hydrophobic and hydrophilic components is the key factor of this gel system. Above the phase transition temperature, the hydrophobic surface increases, allowing for cell attachment, spreading, and proliferation, and the hydrophilic component (mostly from PEG) may initiate cell detachment 201-203. PEG can be modified with dimethacrylates (PEGDM) and these materials have shown to be biocompatible with the unreacted dimethacrylates having a relatively low cytotoxicity 204. Theoretically each PEGDM chain has two functional side groups that would link with the two PNIPAAM chain and form a branched co-polymer 84 system (Figure 5.1). Branched polymers tend to provide high flexibility that can improve the elasticity of the material.
5.4.2. Mechanical Properties of PNIPAAm-based Hydrogels
Several groups studied various mechanical properties of PNIPAAm-based injectable hydrogels, as listed in Table 5.1. In summary, these reported mechanical properties of the gels depend on the cross-linker concentration and testing temperature.
During the initial experiments, we performed compression testing to determine the elastic modulus of 6 hours reacted PNIPAAm-based hydrogel with and without 0.05mol%
PEGDM after immersion in 37°C PBS. The in vitro compressive modulus of MPS-free and MPS-containing gels are 25kPa and 38kPa, respectively. Although only 0.05mol%
PEGDM was added into PNIPAAm, about 50% enhancement of the compressive modulus after immersion was resulted. With further addition of PEGDM or a cross- linker, a higher compressive modulus may be achieved.
5.4.3. Bioactive Material Component
We are interested in developing this material for use in injectable fracture fixation, specifically for vertebral compression fractures. The goal is to create an injectable, bioactive, elastic material that has a compressive modulus similar to vertebral trabecular bone. We modified PNIPAAm with poly(ethylene glycol) dimethacrylate
(PEGDM), a hydrophilic component, with the intent to provide elasticity to the system 1.
Ohtsuki et al. 205 showed that apatite-forming ability can be provided by the addition of alkoxysilane and calcium ion to the organic polymer. The authors modified
PMMA bone cement with alkoxysilanes and calcium salts and showed the cement can 85 provide with apatite-forming in vitro. Based on these findings, we designed a bioactive organic–inorganic hybrid through copolymerization of organic polymer and alkoxysilane compounds with organo-functional groups.
It has been reported that glass in the binary system CaO-SiO2 showed the apatite formation in simulated body fluid (SBF). The reaction of bioactive glasses and glass– ceramics with calcium and phosphate containing biological solutions revealed that the apatite nucleation is triggered by a catalytic effect of silanol (Si–OH) group formed on the surface of the material and accelerated by the release of calcium ion (Ca2+) from the material into the solution. This finding brings an idea that new kinds of bioactive materials can be designed in a range of composition including Ca2+ and Si–OH group as an essential constituent to show bioactivity.
Calcium chloride (CaCl2) salts were selected as the first bioactive organic materials for initiating bioactivity in the silane cross-linking agent added PNIPAAm-
PEGDM co-polymer. CaCl2 retains solubility in biological fluids, but it can be replaced by another organic bioactive material that can be incorporated into the hydrogel network in the same fashion. The goal of adding these bioactive promoters is to provide sufficient potential bioactive sites within the polymer matrix in vitro that would lead direct bone bonding with exposure to the biologoical environment in vivo.
Three forms of alkoxysilanes were explored during the primary screening of the synthesis of the bioactive hydrogel systems, which are methacryloxypropyltrimeth- oxysilane (MPS), 3-aminopropyltriethoxysilane (APS) and 3-aminopropyltrimeth- oxysilane (AMPS). Based on the primary screening, the optimized molar ratio of APS and AMPS is about 0.004 and the optimized molar ratio of MPS was 0.01, which allows 86 maximum bioactivity and retains the phase transition behavior of the gels from minimium cross-linking. Tsuru et al. 206 recently developed organically modified silicates synthesized from polydimethylsiloxane, tetraethoxysilane and calcium nitrate, and reported there was apatite formation on the hybrid after soaking in biological fluids.
Furthermore, apatite formation was also observed on organic–inorganic hybrids containing calcium salt (such as calcium chloride) by two research groups, who synthesized the system through sol–gel processing starting from the MPS silane agent 205,
207. With the addition of MPS and calcium chloride salt in the PNIPAAm-PEGDM, the resulted chemical structure of the hydrogel would be bioactive as shown the Figure 5.2, but would retain the injectable property. The synthetic procedure of bioactive PNIPAAm-
PEGDM based hydrogels is illustrated in Figure 5.3.
5.4.4. Phase Transformation
In situ residual stress at the interfaces of bone-cement and cement-implant due to bulk volumetric shrinkage of the PMMA cement in prosthesis systems was reported 83,
129. Based on the computer modeling analyses and experiments of various research groups, the in situ mechanical properties and fatigue life of the prosthesis systems were directly dependence of the residual stress that may initiate internal cracks in the prosthesis system 83, 208. Although no similar study has yet been report on the effect of
PMMA volumetric shrinkage on the mechanical properties for vertebral fixations, some degree of residual stress is expected at the bone-cement interface that may lead to internal cracks in the surrounding bone. Therefore, an improved bone replacement material for vertebral repair should complete setting within the initial 15 minutes of implantation and has minimal volumetric shrinkage after the polymer setting. 87
We analyzed the phase transformation behavior of PNIPAAm-PEGDM hydrogels with and without 0.05mol% MPS during the initial 60 minutes. All gels were injected in a warm 10mm diameter glass container (at about 30°C) and were observed to complete phase transformation within 3 minutes in a 37°C water bath. The percentage of accumulative volume of the water dispelled from phase transformed gels as a function of immersion time was plotted in Figure 5.4. The accumulative volume of water dispelled becomes steady and no significant additional volume of water was dispelled from the gels was detected after 15 minutes of immersion. The MPS-containing gel showed a higher water dispelled accumulative volume than the MPS-free gel after 15 minutes of immersion. The presence of MPS in the PNIPAAm-PEGDM gel enhances cross-linking in the system that may affect the phase transformation behavior and the amount of water dispelled in the gel at 37°C. Nevertheless, both specimens demonstrated a rapid polymer setting and reached a stable volume within a relatively short period time after immersion
(about 15 minutes).
5.4.5. Rigid Polymer Supplement
One proposed method to improve the mechanical properties of the bioactive hydrogel is adding a rigid polymer to its PNIPAAm backbone. The first attempted rigid polymer was PMMA because it is a common bone replacement material, which has an elastic modulus of about 3GPa. With the addition of PMMA in the PNIPAAm-MPS-
PEGDM was expected to enhance the elastic modulus of the gels. Figure 5.5 is a graph of compressive elastic moduli of the gels, with 0.005 or 0.001 MPS molar ratio after 5 days immersion in two different fluids, as a function of PMMA monomer (MMA) content.
Each data represents an average of 5 measurements. The compressive test data showed 88 that PMMA reduced the elastic modulus of the gels. PMMA is a hydrophobic polymer in
SBF and it may influence the phase transition behavior of the gel because of the highly heterogeneous distribution of the hydrophilic and hydrophobic components in the polymer that leads to poor mechanical properties. A small content of an alternative rigid polymer in SBF at 37°C, such as polyethylene oxide, can be applied to enhance the mechanical properties and retain the phase transition of the PNIPAAm based systems.
5.4.6. Polymerization Kinetics
A validation of the polymerization kinetics was performed by comparing the physical properties polymer solutions of MPS added PNIPAAm-PEGDM polymer
(25wt%) with deionized water. The MPS added polymer was expected to be lightly cross- linked and resulted in a highly viscous polymer solution. However, less viscous polymer solutions were obtained from 6 hours polymerized MPS added gels appeared, and more viscous polymer solutions resulted from 48 hours of polymerization (Figure 5.6). The dissimilarity in performances of the polymer solutions is related to the degree of cross- linking of the polymer. With less polymerization time (6 hours), an incomplete cross- linked gel was properly obtained.
5.5. Summary
Based on the preliminary studies of compressive moduli and phase transformation behavior of PNIPAAm based hydrogels, with the incorporation of PEGDM, MPS and calcium chloride, the 48 hours polymerized gels showed potential to as fit candidates for vertebrae augmentations. However, the reason of reduced elastic modulus from the addition of PMMA to the MPS-containing hydrogels was not clear that is worth 89 exploring in future. As a single phase material system, PNIPAAm-PEGDM may eliminate the problems that initiated from the interfaces between the two material phases of fillers reinforced composites. Further exploration will examine the usefulness of the injectable bioactive hydrogels in stabilizing compression fractures of the vertebrae without causing undo stress risers on adjacent vertebrae because of its more closely matched modulus to the vertebral bone. 90
Table 5.1: Mechanical properties of PNIPAAm-based hydrogels reported by various research groups.
Reported Mechanical Testing Author Material Values Properties Condition (kPa)
PNIPAAm-co- PBS and Ultra Complex shear Stile and Healy 198 (4mol%) Acrylic purified water; 0.07 - 0.10 modulus at 10Hz acid at 37°C
PNIPAAm-co- PBS and Ultra Complex shear Stile and Healy 198 (4mol%) Acrylic purified water; 0.02 - 0.04 modulus at 10Hz acid at 22°C
PNIPAAm-co- N,N- Compressive Distilled water; Zhang et al. 209 5 - 28 Dimethylacrylam modulus at 22°C ide
Peptide- PNIPAAm-co- Complex shear Ultra purified Kim and Healy 210 0.10 - 0.12 (4wt%) Acrylic modulus at 10Hz water; at 37°C acid
Peptide- PNIPAAm-co- Complex shear Ultra purified Kim and Healy 210 0.02 - 0.03 (4wt%) Acrylic modulus at 10Hz water; at 22°C acid
PNIPAAm-co- Surface elastic Distilled water; Ohya et al. 211, 212 (20 wt/v%) modulus 158 - 286 at 37°C gelatin (AFM indentation)
PNIPAAm-co- Surface elastic Cell culture Ohya et al. 211, 212 (20 wt/v%) modulus medium; at 155 - 333 gelatin (AFM indentation) 37°C 91
25°C 37°C
PNIPAAm PEGDM
Figure 5-1: A simplified illustration of the polymer molecular structure of PNIPAAm-PEGDM at 25°C and 37°C. 92
Figure 5-2: An illustration of the chemical structure of the PNIPAAm-MPS- PEGDM gel with calcium chloride salts.
93
N-isopropylacrylamide (NIPAAm) Alkoxysilane (Silane) Polyethylene glycol dimethacrylate in methanol (PEGDM) 0.5% (v/v)
Mixture of NIPAAm-Silane-PEG with 0 to 0.01 molar ratio of Silane 2-2’-Azobis isobutyronitrile (AIBN)
Calcium Chloride in Methanol (CaCl2:MPS = 2:1)
Purge with Nitrogen gas
Polymerization at 65°C for 6 or 48 h
Stir the mixture overnight at 25°C
Dry in vacuum overnight at 25°C
Thermo-sensitive bioactive hydrogel
Figure 5-3: Synthetic procedure of PNIPAAm based hydrogel with 0 to 0.01 molar ratio of MPS. 94
MPS-containing (0.005 molar ratio) Gel MPS-free Gel 60%
50%
40%
30%
20%
10% Accumulative Volume of Water Dispelled [%] 0% 0 10203040506070 Time [minute]
Figure 5-4: The percentage of accumulative water dispelled volume during the initial 60 minutes of immersion in a 37°C water bath. 95
MPS 0.005 (PBS) 700 MPS 0.001 (PBS) MPS 0.005 (SBF) 600 MPS 0.001 (SBF)
500
400
300
Elastic[kPa] Modulus 200
100
0 -0.002 0 0.002 0.004 0.006 0.008 0.01 0.012
MMA molar ratio to NIPPAAm
Figure 5-5: A plot of compressive elastic moduli (n=5) as a function of PMMA monomer (MMA) content.
96
Various coupling agent compositions in PNIPAAm-PEGDM (Reaction Time: 6 hours)
0.012
0.008
0.004
MMA (molar ratio)
0.000
0.000 0.002 0.004 0.006 MPS (molar ratio)
Various coupling agent compositions in PNIPAAm-PEGDM
(Reaction time: 48 hours) 0.012
0.008
0.004
MMA (molar ratio)
0.000 0.000 0.004 0.008 0.012 MPS (molar ratio)
Thick solution and gelation at 37°C: Thin solution and gelation at 37°C:
Figure 5-6: Observations of the physical properties of the polymer solutions of either 6 or 48 hours polymerized PNIPAAm-MPS-PEGDM with various MPS content (from 0 to 0.01 molar ratio).
97
6. An injectable Bioactive Hydrogel Composite
6.1. Introduction
Current options for alleviating pain due to vertebral fracture include vertebroplasty as well as kyphoplasty 213 which utilizes two balloons inserted into the vertebral space to expand and to compress the trabecular bone tissue, creating a central cavity. Subsequent to cavity preparation, polymethylmethacrylate (PMMA) bone cement is injected into the cavity using percutaneous techniques. Typically, PMMA bone cement is used to restore the mechanical integrity of the vertebral body by stabilizing the trabecular fractures and relieving pain.
Belkoff and Molloy 118 reported that in situ temperature of the PMMA exothermic reaction can reach 39 to 113°C in the vertebral during vertebroplasty. Such elevated temperature is sufficiently high to pose a risk of thermal osteo-necrosis for the vertebroplasty patients. Another issue associated with vertebro or kyphoplasty 213 with
PMMA cement include increased rigidity of the vertebral body in comparison to the more superior and inferior vertebrae due to the high modulus of the PMMA (1-3GPa) in comparison to trabecular bone tissue (0.5GPa). This can lead to stress inconsistencies along the length of the spine and may lead to subsequent fracture. Subsequent fracture has been reported in 26% of the kyphoplasty cases by Fribourg et al. 214. In addition, pulmonary embolism leading to cardiac failure has been reported, although not frequently, for patients receiving acrylic vertebroplasty and is cause for concern when the
PMMA exits its intended locale 215. Moreover, PMMA does not offer the advantage of osteoconductivity that is appreciated with bioactive cements. 98
Numerous groups have examined bioactive cements, either calcium phosphate cements or polymeric cements containing bioactive ceramics. While the bioactivity has been accepted as an improvement over PMMA cements 135, 136, the mechanical properties of the calcium phosphate cements have been questioned for sufficient fatigue strength and high modulus mismatch to cancellous bone, which may not be adequate to prevent fracture from recurring both initially and over the life of the procedure 133. An example of these bioactive injectable cements is the commercially available self-hardening calcium phosphate composites, Norian®, for mandibular and vertebral augmentations 62,
137, 216, consisting of equimolar amounts of tetracalcium phosphate and dicalcium phosphate anhydrous, hardens when mixed with water and forms a resorbable hydroxyapatite (HA) as the end-product. This material provides excellent mechanical properties and direct bone ingrowth ability, and sets in situ after injection. However, the hardened ceramic material often contains pores and leads to poor cement-bone interfacial mechanical properties 139, 144. The handling characteristics of Norian® in the present form is non-optimal for injection, as the cement is in a paste form and therefore difficult to inject in the site 132, 140.
Many injectable bone substitutes combine polymer and bioactive ceramics to offer several advantages, such as easier tailoring mechanical properties and better interaction with the implantation site are being investigated 130, 146, 173. With sufficient viscosity, these polymer may use as injectable bone cement. Kokubo et al. extensively studied polymethacrylate / bioactive ceramic cements in the 1990s 217. They incorporated various bioactive glass beads and calcium phosphate granules to reinforce the polymer.
Among these composites, the bioactive glass reinforced composites provided the better 99 mechanical properties than the others and higher bioactivity. Unfortunately, some bioactive glass beads separated from the bone cement, which led to poor cement-bone interfacial properties 218-220. Lack of adhesion between the ceramic fillers and the polymer matrix is the major factor responsible for the filler-matrix debonding that ultimately resulted in a lower load capacility of the cement 221, 222. Currently, they are employing an alternative method of developing a bioactive injectable biomaterial using nanohybrid organic-inorganic process by adding silane coupled calcium salts. The apatite is formed locally within the polymer chain, which provides an improved ceramic-polymer adhesion with the apatite formed in situ 130, 223, 224.
We developed an injectable bioactive thermo-sensitive biomaterial as vertebrae replacements. Our goal is to create an elastic material that has a compressive modulus similar to vertebral trabecular bone. Among thermosensitive hydrogels, poly(N- isopropylacryamide) (PNIPAAm) is attractive because of its lower critical solution temperature (LCST) behavior at between 29 and 36°C in an aqueous solution 1, 191, 225.
Thus, PNIPAAm is soluble in water at room temperature, but is insoluble and indeed gels at physiological temperature. We modified PNIPAAm with poly(ethylene glycol) dimethacrylate (PEGDM) 1. The PEGDM domains provide the hydrogels with a hydrophilic component which allows the entangled gels to exhibit a wide range of swelling and mechanical properties 1. Further, we engineered a bioactive component into the macromolecular structure to facilitate the formation of calcium phosphate nucleation sites on the injectable polymer system through incorporation of tri- methacryloxypropyltrimethoxysilane (MPS). With the presence of the calcium phosphate sites, bone-like apatite can be formed on/in three-dimensionally-structured organic 100 hydrogel matrices in vitro using a short-term immersion process. 138, 141. In this study, we examined the polymer structure, the gelation temperature (LCST) and viscosity of the bioactive hydrogel as a function of MPS compositions. Also we determined the gels in vitro mechanical properties, swelling characteristics and bioactivity in different immersion media in order to determine the feasibility of these materials as a fixation medium for vertebral compression fractures.
6.2. Materials and Methods
6.2.1. Synthesis PNIPAAm-MPS-PEGDM
Nitrogen gas was bubbled through a mixture of purified NIPAAm, PEGDM in
200:1 weight ratios (about 0.05mol% of PEGDM) in methanol solvent. 2- Various molar ratios of MPS:NIPAAm (0 to 0.01) and constant molar ratio of calcium chloride:MPS of
2:1 was added to the polymer mixture with the initiator 2’-Azobis isobutyronitrite. The mixture was polymerized at 65ºC for 48 hours, and was stirred overnight at room temperature. The polymer was dried for a few days in a vacuum oven to further remove the solvent, and then the dried PNIPAAm-MPS-PEGDM was ground to fine powder.
Polymer solution was prepared with 25wt% polymer and 75wt% deionized (DI) water at room temperature.
6.2.2. Polymer Characterization
After synthesis, the polymers with various MPS contents were characterized using fourier transform infrared spectroscopy (FTIR; Magna-IR 560 , Nicolet, Madison, WI) and a field emission environmental scanning electron microscope (FE-SEM), equipped with Electron-Dispersive Spectroscopy (EDS; Philips XL 30, Netherlands), at 1.4 x 10-9 101 mtorr vacuum level and 15keV voltage. The polymer powder was dissolved in acetone for FTIR analysis. Prior to FE-SEM/EDS determination, the polymer was coated with platinum with a platinum sputter for 60 seconds.
6.2.3. LCST Determination
The LCST of the polymer solution (25wt% of polymer in DI water) was evaluated using differential scanning calorimetry (DSC; DSC 2010, TA Instruments, New Castle,
DE) as a function of MPS content. The samples were heated at 1°C/min to 50°C under nitrogen gas purge. Once a plot of heat flow vs. temperature was obtained, the LCST was determined to be the temperature at the minimum heat flow point of the curve 226, 227.
6.2.4. Injectability and Viscosity Studies
An injection process can be considered in two physical processes: (1) delivery of the injected materials in the implantation site, and (2) infiltration of the injected materials with the existing biological environment 228. The delivery of the bioactive hydrogel was examined qualitatively via the injection of polymer solutions with various MPS molar ratio through a 20 gauge needle in 37°C SBF manually with a moderate force. The in situ infiltration properties of the bioactive hydrogel were studied via the viscosity measurement of the polymer solution as a function of MPS molar ratio. The viscosity measurement of three polymer solution (25wt% of PNIAAm-PEGDM with either 0,
0.005 or 0.01 MPS molar ratio in deionized water) was performed using a dynamic stress rheometer (DSR; DSR-200 Rheometics, TA Instruments, New Castle, DE) with the cone/plate test setup at room temperature. The applied shear rate was 1 to 100s-1 for all specimens. 102
6.2.5. Change in Volume Estimation
The firm gels were prepared by molding the polymer solution at 37°C. The gels were immersed in phosphate buffered solution (PBS), simulated body fluid (SBF) or double ion concentrated SBF (SBFx2) from 1 to 544 hours. In the preparation of PBS, the
9.6g of Dulbecco’s PBS powder (Sigma-Aldrich Co., St.Louis, MO) was dissolved in DI water, the ion concentration of PBS is listed in Table 6.1. For the preparation of SBF and
SBFx2, various chemicals (as provided in Table 6.1) were dissolved in 1.0mN Tris-HCl
(Sigma-Aldrich Co., St.Louis, MO) and buffered to pH 7.6 at 37ºC. The temperature of the solutions was maintained at 37°C. The volume change of hydrogel after immersion in
PBS, SBF or SBFx2 was calculated by subtracting the gel volume at 37ºC after one hour of immersion from the gel volume after various immersion hours. The average volume change is determined based on the volume change of three specimens. Initial water dispelled volumetric change from phase transformation was estimated only in 37°C SBF and reported previously in Section 5.4.4.
6.2.6. Bioactivity Characterization
Recent studies showed protein adsorption can alter the in vitro reaction kinetics of bioactive surfaces 229, 230. The effect of protein adsorption on the bioactivity of
PNIPAAm-MPS(0.005 mole)-PEGDM specimens were analyzed using a biological media containing protein, 10vol% fetal bovine serum (FBS; Biomeda Co., Foster City,
CA)-90vol% simulated body fluid (SBF). The 37°C formed 0.005 mole MPS-containing gels (in cylinder structure with an average 5mm diameter and 7mm height) were immersed in 5ml of 37°C biological fluids, SBF and FBS-SBF, for the in vitro formation of the apatite within the bioactive polymer network. The immerse solution is renewed 103 every 30 days. After soaking for a given period, the specimens were removed from the solution, ultrasonically washed with 37°C deionized water for 30 minutes, and then vacuum dried at room temperature.
Based on previous literature, the formation of bone bonding is known through the development of an apatite layer in vitro, where an apatite layer with a calcium-to- phosphorous ratio higher than 1.4 is known to promote direct bone bonding 67. The presence of calcium and phosphate after immersion in SBF was examined using FE-
SEM/EDX (Philips XL 30, Netherlands) and FTIR (Magna-IR, Nicolet, Madison, WI), respectively. The apatite formation within the polymer after long term immersion (30, 45,
60, 75 and 90 days) in SBF was verified by powder X-ray diffractometry (XRD: Rigaku
X-ray diffractometer, Rigaku/USA Inc., Danvers, MA, USA). The calcium to phosphorus atomic ratios of 5 bioactive nodules within the polymer matrix were also calculated using
EDS after immersion in SBF and FBS-SBF from 30 to 75 days. Further, the chlorine adsorption of the 5 and 30 days SBF immersed polymer matrix was determined by quantifying the chlorine, carbon and oxygen composition across the gel thickness.
6.2.7. In vitro Compressive Test
The polymer solution was poured into a 10mm diameter polymer cylinder and the mold was heated to 37ºC solution for phase transformation. Then the firm gels were stored in the 37ºC immersion media (PBS, SBF and FBS-SBF) for 5 days to reach its equilibrium state. After the gel stabilized with the immersion solution, the effect volume change of the gels on its mechanical properties would be negligible. The specimens were tested in unconfined compression using an Instron mechanical testing system (Instron
Model 4442, Park Ridge, IL) fitted with a 50N load cell and 37ºC immersion bath with 104 either PBS or SBF. Specimens were compressed at a strain rate of 100% strain per minute. Load and displacement data was recorded at 20 points per second with the
Instron Series IX software. This data was converted to stress and strain values in
Microsoft Excel using the specimen’s initial dimensions. A compressive elastic modulus was measured as the initial linear slope of the stress-strain response at 0.05mm/mm strain. In the mechanical study, 5 specimens of each group were tested. Long term in vitro mechanical properties were also analyzed for 30 days PBS and SBF immersed specimens using the compression test. One-way ANOVA analyses were performed for compressive modulus results at 0.05 level of significance..
6.3. Results
6.3.1. Physical Properties
The PNIPAAm-PEGDM polymer and the PNIPAAm-MPS-PEGDM polymer with no calcium chloride were miscible with deionized water within 2 days and appeared to be transparent. However, with the addition of calcium chloride, the polymer solution was opaque after 2 days of dissolving and became transparent slowly within 14 days at room temperature.
Figure 6.1 showed the FTIR spectra of the gels with different chemistries.
Comparing the FTIR spectra of MPS-free gels to MPS-containing gels, peak 1030cm-1
(Si-O groups) appeared only in the materials with 0.005 and 0.01 molar ratio of MPS.
The 0.001 molar ratio MPS content in the gel was not detectable likely due to the FTIR sensitivity limit (±1%). No significant change was measured on the LCST of the
PNIPAAm-PEGDM gels regardless of MPS concentration (p=0.62). Consequently, all 105 polymer solutions (with upto 0.01 molar ratio of MPS) would transform to firm gels at about 32.6°C (the LSCTs).
Ease of injection through a 20 gauge needle was accomplished with MPS concentrations of 0.01 molar ratio or less. The gel became solid instantaneously in the
37°C SBF (Figure 6.2). Based on the DSR data, the viscosity measurements (η) of the polymer solution of 0, 0.005 and 0.01 molar ratio of MPS with respect to shear rate of 10 to 100s-1 are plotted in Figure 6.3. At 90 to 100s-1 shear rate, the viscosities of the specimens became relatively constant. For PNIPAAm-PEGDM with 0, 0.005 and 0.01 molar ratio MPS content, the viscosity at 100s-1 was determined to be 3, 118 and 83Poise, respectively.
When gels were heated to 37°C, all hydrogels became opaque and pliable in PBS,
SBF and SBFx2. But once they were immersed in SBF and SBFx2 for 24 hours, all hydrogels turned transparent and more rigid. During the rapid phase transition at 37°C, the hydrogels dramatically collapsed and released a large fraction of pore water within the first hour of heating. Then the specimens tended to swell (increase in volume) and reach equilibrium (volume change within 10%). The average volume change (n=3) of
PNIPAAm-PEGDM with 0 and 0.005 molar ratio of MPS, immersed in three biological media: PBS, SBF and SBFx2, was examined from 1 to 544 hours (Figures 6.4 and 6.5).
As shown in Table 6.1, PBS has less ion content than SBF and SBFx2 has the greatest ion content among the three immersion media. The average volume change of gel with or without MPS appeared to vary with the ion content of the media. Figure 6.4 and 6.5 showed that the PBS immersed gels showed higher average volume change than SBF or
SBFx2 immersed gels during the initial 4 days immersion. 106
In the determination of the effect of MPS addition in the PNIPAAm-PEGM, the one with 0.005 MPS molar ratio exhibited less average volume change than the one without MPS. The maximum average volume change of 14% of the MPS-containing gel was estimated after 6 and 14 hours immersion in PBS, while the maximum reading of
29% was estimated after 14 hours immersion in PBS. In brief, the less average volume changes of the bioactive hydrogel after immersion in higher ion concentrated biological fluids. After 5 days of immersion, the average volume change of all MPS containing hydrogel specimens became stable (within 10%) in all biological conditions.
6.3.2. Bioactivity Kinetics
With the addition of MPS and calcium chloride, the PNIPAAm based hydrogels becomes bioactive gels. FE-SEM images and EDX spectra of the MPS-containing gels showed bioactive sites were nucleated locally and grew three-dimensionally into calcium phosphate nodules after 30 days of immersion in SBF at 37°C. Figure 6a showed the bioactive site was embedded in the polymer matrices. In comparison of the FE-SEM images of 0, 3 and 5 days SBF immersed gel (at 10000x, 10000x and 15000x, respectively) with the images of 30, 45 and 60 days SBF immersed gel (at 1000x, 1000x and 5000x, respectively), there displayed an increase number of bioactive sites within the polymer matrices (Figure 6.6). Figure 6.7 illustrated the FTIR spectra of 0, 5 and 30 days
-1 SBF immersed MPS-containing gels. 1054 cm peak, represents the –PO4 groups, was detected in the FTIR spectra of 30 days SBF immersed gel, but it did not show in the other two spectra.
The XRD spectra were normalized with the unimmersed specimens accordingly.
Figure 6.8 demonstrated the normalized XRD spectra of the MPS-free gels after 45 days 107
SFB immersion and MPS-containing gels after 5, 30, 45, 60 and 75 days SBF immersion and 45 days FBS-SBF immersion. No indication of calcium phosphate formation was revealed in the XRD spectra of the 45 days SBF immersed MPS-free gel. The MPS- containing gels showed changes in the XRD pattern as immersion times changed from 5 to 75 days immersion in SBF. The differences in the XRD patterns between 5 days immersed bioactive specimen and the 30, 45 and 60 days SBF immersed MPS-containing specimens indicate the formation of amorphous calcium phosphate (Figure 6.8). After 75 days immersion in SBF, the XRD peaks assignable to apatite 223 were observed. The FE-
SEM/EDX results of these specimens (Figure 6.9) illustrated calcium phosphate crystals were formed on the surface of the polymer. However, the XRD spectra and FE-SEM image of the 75 days FBS-SBF immersed gels showed no apatite formation (Figure 6.8 and 6.10, respectively).
The initial detection of phosphorous in the bioactive sites of the SBF immersed gels was shown after 20 days immersion, but it was 10 days delayed for the FBS-SBF immersed specimens (Figure 6.11). Once the calcium phosphate formed in the FBS-SBF gels, its phosphorous content increased in a faster pace than the ones of SBF immersed; and after 75 days, no statistical different were detected in the calcium to phosphorous ratios in the bioactive sites between the two specimens. The FE-SEM/EDX result of the unimmersed specimens (Figure 6.6) showed the calcium chloride contained bioactive nodule was embedded in the polymer matrix. Phosphorus from the surrounding fluid was gradually reacted with the calcium rich bioactive nodules and led to calcium phosphate formations. The chlorine, carbon and oxygen composition across the gel thickness were varied between 5 and 30 days SBF immersed specimens. Figure 6.12 showed chlorine 108 was not detectable at the center of 5 days SBF immersed specimen while about 1 mol% of chlorine was detected across the 30 days SBF immersed gel thickness.
6.3.3. In vitro Mechanical Properties
The elastic modulus of 5 days SFB immersed gels was significantly superior to those of 5 days PBS immersed gel (p < 0.001), however these SBF gels did not demonstrate a significant change from those of 5 days FBS-SBF immersed gel, as shown in Figure 6.13. The control specimens (PNIPAAm-PEGDM with no MPS addition) demonstrated higher elastic modulus after immersion in SBF and FBS-SBF than those immersed in PBS (148% and 228% more respectively).With the addition of 0.005 MPS molar ratio, the elastic modulus of the PBS immersed firm gel had a 203% increment reached to 0.18MPa; while the SBF immersed one had a 293% increment at about
0.6MPa. But the elastic modulus of SBF immersed gel with 0.0075 and 0.01 MPS molar ratio had a lower value than the one with 0.005 MPS molar ratio (132% and 89% less respectively). Thus, the maximum elastic modulus of the bioactive gels was achieved with the addition of 0.005 MPS molar ratio after 5 days immersion in either SBF or FBS-
SBF. Figure 6.14 shows in vitro elastic modulus of 30 days PBS and SBF immersed
PNIPAAm-PEGDM with 0, 0.001 and 0.005 molar ratio of MPS, which showed no significant differences from the 5 days immersed samples.
6.4. Discussion
PNIPAAm based hydrogels are most widely studied lower critical solution temperature (LCST) polymer due to the drastic transition from solution (with the addition of water) at room temperature to solid (gel) above 32-34ºC. During gelation, physical 109 interactions within the polymer network occur and lead to a phase transformation. There are no initiators or cross-linking agents necessary to solidify the gel in situ which makes a clean rejectable biomaterial without the biocompatibility issues associated with residual monomer and initiator agents. Recent reports have shown the ability of PNIPPAm-based materials to support cell growth in vitro and in vivo and found that such materials exhibit little if any cytotoxicity 198, 212, 231 . Stile and Healy reported that cross-linked PNIPAAm- based polymers have the potential to serve as functional scaffolds in tissue engineering applications 198. These materials have yet to be examined for fracture fixation. The
Taguchi research group focused on apatite coating on hydrophilic polymer grafted polyethylene films 142. Other commonly used hydrophilic polymers, such as, poly(acryl amide) (PAAm) and poly(acrylic acid) (PAAc) were also employed 143. Kawai et al. have reported that carboxyl (-COOH) groups play an effective role in heterogeneous nucleation of apatite in the body environment in the presence of calcium salts 232.
Additionally, Madsen and Peppas have shown gels prepared from poly(methacrylic acid) and PEG-monomethacrylate complex with significant amounts of calcium in calcium enriched phosphate buffered solutions as drug delivery systems 185. Kabashita et al. demonstrated that with an optimized content of carboxyl groups and calcium chloride in the polymer, an apatite layer would form on a hydrogel. Once the apatite nuclei are formed on the substrate, they can grow spontaneously by consuming the calcium and phosphate from the surrounding biological environment 217, 223. While this approach of immersion formation of bioactive layers on the surface of polymers in general and hydrogels in particular have been recently described, there are to date no bioactive injectable hydrogels reported in the literature. 110
The addition of MPS to the PNIPAAm-PEGDM hydrogel system encouraged calcium phosphate formation thoughout the gel thickness after SBF and FBS-SBF immersion. This mechanism is likely controlled by the silanol groups (Si-OH) in MPS, which provided attachment sites for calcium and initiated mineral dissolution from the simulated biological environments. As shown in the FTIR spectra of the gels with various MPS content, the Si-O groups were detected for the 0.005 and 0.01 MPS molar ratio. It was observed that polymer with more 0.01 MPS molar ratio content would be immiscible in water even after 14 days at room temperature. It is possible that the high degree of cross-linking from the addition of MPS has prevent the PNIPAAm-PEGDM be soluble with water. Therefore, all reported material characterizations were performed on
PNIPAAm-PEGDM with MPS content from 0 to 0.01 molar ratio.
The solubility the gel in water was influenced with the addition of calcium salts.
The presence of calcium ions were attached to the gel matrix with MPS. The silica calcium salts may be less miscible with water; therefore an extended solubilizing time is required. In addition, solubility of hydrogels in deionized water can be influenced by many other factors. For instance, reduced temperature (as low as 4°C) can enhance the amount of hydrogen bonding interactions between water and hydrophilic amide group dominate causing the polymer to be miscible in water faster 197. Decreased polymer particle size to a sub-micron level can increase the overall surface area for polymer-water interaction that also can improve the solubility of the polymer in water.
For the compositional ranges examined, no significant influence was observed with MPS and PEGDM content on the phase transition behavior and temperature (LSCT) of the hydrogels. An instantaneous phase transition and rapid dehydration of the gels at 111
37°C in either PBS, SBF, SBFx2 or FBS-SBF were demonstrated. This sharp transition from the balancing hydrophobic / hydrophilic groups of the PNIPAAm is consistent with previously reports of injecting NIPAAm in water and PBS at 35°C 190, 197. Such rapid solidification of the polymer solution in simulated biological conditions led to the motivation the use of PNIPAAm-PEGDM with MPS as an injectable bone material. Even though rapid dehydration of the gel would lead to quick deswelling of the material in
SBF, the volume did not change more than 10% after the first hour of injection. This factor is beneficial for bone repair applications because it offers a physically stabilized material in the body first hour of surgery that would reduce the overall surgical risk.
Based on the studies of Baroud et al., for vertebroplasty ideal PMMA bone cement should maintain a viscosity of 1000 to 2000 Poise (at shear rate of 40s-1) during the initial 20 minutes of local delivery using a cannula and allow proper infiltration in bone without polymer leakage to the adjoined vertebral bodies. The room temperature measured viscosity of PNIPAAm-PEGDM with 0.005 MPS molar ratio is 169 Poise at shear rate 40s-1 (Figure 6.2). As Arai et al. reported that PNIPAAM solution presents a such as strong thermo-viscosifying effect that its viscosity increased from 0.5 to 100
Poise at its LCST 32°C due to the rapid phase transformation of the gel 231. Although the
MPS-containing gel has a much lower viscosity than the PMMA bone cement, the gel is expected to increase viscosity during injection as the local temperature may increase slightly (approaching body temperature) and set to firm gel within seconds in body.
Further exploration of the in situ viscosity of PNIPAAm-MPS-PEGDM gels in vertebral body would be necessary. 112
Swelling properties of the gels varied with immersion media, although the differences were less pronounced for the MPS-containing systems, as shown in Figures
6.4 and 6.5. Stile et al. reported their PNIPAAm based systems exhibited a reduction in volume change when the phase transition occurred in present of ions (in PBS) comparing to the ones in ultra purified water 198. An increase in ion concentration of the immersion media resulted in less solution uptake by the hydrogels, regardless of MPS presence in the system. The increased ions concentration in SBF and SBFx2 led to additional cross- linking of the gels that resulted in less swell or even de-swell of the materials in vitro.
Matsuura et al also reported the calcium ion concentration in an aqueous solution can influenced the swelling behavior of a natural hydrogel, bovine vitreous body. Less swelling of the hydrogel was measured from less calcium ion concentration. Therefore, the ions concentration of the immersion media is a key factor on the swelling behavior of the gels.
In addition, less average volume change was determined in MPS-containing gels after immersion in various biological media comparing to MPS-free gels as a function of immersion time (Figure 6.4 and 6.5). With the addition of cross-linker (MPS), the gels likely to have a higher degree of cross-linking and result in de-swell of the gels after immersion in various media. MPS contributed to increased viscosity and elastic modulus of the material, likely the result of the condensation of silanol which forms siloxane (Si-
O-Si) serving to cross-link the polymer backbone. In theory, the elastic modulus is expected to increase with the MPS content, however, the elastic moduli of the gels with
0.0075 and 0.01 molar ratio of MPS were lower than the one containing 0.0005 molar ratio. The gels with MPS more than 0.005 molar ratio may have poor distribution of 113 siloxane bonds and high content of siloxane bonds can lead to formation of a structure similar to brittle silica glass.
Figure 6.10 showed the differences in elastic moduli were more pronounced in the presence of SBF. As compared to PBS, immersion in SBF served to enhance the bioactivity reactions as well as increase the mechanical properties, even after 5 days of immersion in vitro. Previously mentioned, the SBF immersed gels is likely exhibited more cross-linking from the higher ions concentration of SBF than the PBS immersed ones. With higher degree of cross-linking, superior elastic modulus would be resulted.
Based on the FE-SEM images, EDX spectra and FTIR spectra, the nucleation and growth of calcium-phosphate were showed in the PNIPAAm-0.005 molar ratio MPS-
PEGDM matrices of 30 days SBF and FBS-SBF immersed specimens (Figures 6.6, 6.7 and 6.9, respectively). The XRD spectra revealed that these bioactive sites were amorphous even after 30 days immersed in SBF, as shown in Figure 8. These amorphous calcium phosphates have insignificant influence on the mechanical properties of the gels.
After 75 days immersion, the amorphous calcium phosphate grew into apatite crystals. As a substantial amount of apatite forms within the polymer and leads to direct bone bonding, the global mechanical properties of the gel are expected to be improved. Also the calcium to phosphorus ratio of the bioactive element is slowly reduced and tends to approach the 1.55 to 2.2 range, that signifies hydroxyapatite, the primary inorganic component of bone 16.
Moreover, 30 days SBF immersed test outputs appeared to be highly scattered and achieved superior average elastic moduli higher than the 5 days SBF immersed ones. A few complications occurred during the preparation of the cylindrical specimens for 114 compression test, such as air bubbles trapped in the specimens and tendency of the gels adhere to the surface of the mold. Several procedures were applied to limit air bubbles present in the specimens. First the polymer was ground into micron size powder and added to deaerate the deionized water in a nitrogen gas contained box. The polymer was stored in a nitrogen contained glass container for 2 days at room temperature. Then the polymer solution was poured in a sealed small cylindrical glass container under nitrogen gas and the container was placed in a 37°C vacuum oven for 2 hours until the gel solidified. The resulted specimens contained no macro-size bubbles that would affect the compression test result. Since the gel has high tendency to adhere to most polymer and glass surface. Therefore, the applied glass molds were coated with a layer of Teflon to prevent gel attaching.
The presence of chlorine ions throughout the gel thickness of 30 days SBF immersed samples appeared to have negligible influence on the mechanical properties of the system (Figure 6.14). Detection of chlorine in the center of the gel displayed that there were fluid and ions diffusion within the gel meshwork. Nucleation and growth of apatite are also controlled by specific interactions between the calcium and phosphate ions and protein 229. The adsorption of proteins on the mineral surface can alter the nucleation and crystallization rate of apatite 233. In this study, apatite formation in the
MPS-containing gel was observed after 75 days SBF immersion while the protein adsorbed layer may limit the transformation of ions to the bioactive surface that would affect the kinetics of calcium phosphate formation in the MPS-containing gels. The in vitro mechanical properties of the 5 days FBS-SBF immersed specimens are very similar to the ones of SBF (Figure 6.13). However, the long term effect of the FBS-SBF on the 115 bioactivity was highly distinct from the ones of SBF. The proliferation, crystallization and phosphate uptake of the nodules of the FBS-SBF immersed gels were much slower than the ones of SBF, as shown in FE-SEM/EDX images and calcium to phosphorous ratios.
6.5. Conclusion
With the addition PEGDM and MPS, the compressive modulus of the PNIPAAm systems reached to 0.7MPa that is the highest value ever reported. The challenge of
PNIPAAm-MPS-PEGDM systems is to balance the network-forming and bioactivity inducing MPS with the gain in elastic recovery induced by PEGDM addition to the
PNIPPAAm, all while maintaining an injectable material system. We have shown that this hydrogel family is capable of “structural” mechanical behavior, while potentially providing nucleation sites for cell attachment through its bioactive surface. The future work would be selecting the appropriate permutations for this system that would lead to higher modulus materials and closely match with the modulus of vertebral trabecular bone. 116
Table 6.1: Ion composition and concentration of three biological medias: PBS, SBF and SBFx2.
Concentration in Concentration in Concentration in Ion PBS [mmol] SBF [mmol] SBFx2 [mmol]
Na+ 6.95 152 304
Cl- 1.55 135 270
K+ 1.5 5 10
- H2PO4 4.21 1.0 2.0
Mg+2 - 1.5 3
- HCO3 - 271.5 543
SO4-2 - 0.527 1.054
Ca+2 - 2.5 5
117
1030cm-1
Figure 6-1: FTIR spectra of PNIPAAm-PEGDM with 0, 0.001, 0.005 and 0.01 mole MPS.
118
Figure 6-2: Phase transformation of PNIPAAm-PEGDM with and without MPS. 119
400 PNIPAAm-0.005MPS-PEGDM PNIPAAm-0.01MPS-PEGDM 350 PNIPAAm-PEGDM
] 300
250
200 150 Viscosity [Poise Viscosity 100
50
0 0 20406080100 Shear Rate [1/s]
Figure 6-3: Viscosity measurements of PNIPAAm-PEGDM with 0, 0.005 and 0.01MPS molar ratio with respect to shear rate from 10 to 100s-1. 120
40% immersed in PBS immersed in SBF 30% immersion in SBFx2
20%
10%
0% 0 100 200 300 400 500 600 -10% Average volume change volume Average
-20% Immersion time [days]
Figure 6-4: Average volume change (n=3) of PNIPAAm-PEGDM after immersed in PBS, SBF and SBFx2 from 1 to 544 hours. 121
immersed in PBS 40% immersed in SBF 30% immersed in SBFx2
20%
10%
0%
0 100 200 300 400 500 600 -10% change volume Average -20% immersion time (hour)
Figure 6-5: Average volume change (n=3) of PNIPAAm-MPS (0.005 mole)-PEGDM after immersed in PBS, SBF and SBFx2 from 1 to 544 hours.
122
(a)
Figure 6-6: FE-SEM/EDX images of the bioactive nodules formed within the gel network (a) before immersion; (b) after 3 days; (c) after 5 days; (d) after 30 days; (e) after 45 days and (f) after 60 days immersion in SBF. 123
3.5
3.0 1054 cm-1 2.5 PO
Absorbance 4 2.0
1.5 30 days
1.0 5 days 0.5 0 days 0.0
-0.5 1400 1300 1200 1100 1000 900 800 Wavenumbers (cm-1)
Figure 6-7: FTIR spectra indicated the phosphate attachment of the hydrogels as a function of immersion times. 124
Figure 6-8: Normalized XRD spectra of SBF and FBS-SBF immersed PNIPAAm- PEGDM with and without 0.005 molar ratio MPS. 125
Figure 6-9: FE-SEM/EDX image of calcium phosphate crystals in the PNIPAAm- PEGDM meshwork after 75 days in SBF. 126
75 days
Figure 6-10: FE-SEM/EDX image of calcium phosphate in the PNIPAAm-PEGDM meshwork after 75 days in FBS-SBF. 127
10 SBF FBS-SBF
8
6
Ca/P ratio Ca/P 4
2
0 0 1020304050607080 Immersion Time [day]
Figure 6-11: EDX data of the surface calcium (Ca) to phosphorous (P) ratios of the long term SBF and FBS-SBF immersed MPS-containing (0.005 molar ratio) gels.
128
3 2 1
specimen
5mm
5 days immersed 30 days immersed No Position Chlorine Carbon Oxygen Chlorine Carbon Oxygen 1 center 0 91.23 8.77 1.03 88.56 10.41 2middle 0.05 89.1 10.85 1.28 86.56 12.16 3surface 0.22 90.21 9.57 1.47 89.45 9.08
Figure 6-12: The composition of chlorine, carbon and oxygen was measured on the cross-section of polymer matrix of 5 and 30 days SBF immersed MPS- containing gels. 129
5d PBS 700 5d SBF 10%FBS + SBF 600
500
400
300 200
Elastic Modulus [kPa] 100
0 -0.002 0 0.002 0.004 0.006 0.008 0.01 0.012
MPS molar ratio to NIPAAm
Figure 6-13: In vitro elastic modulus of PNIPAAm-PEGDM with various MPS molar ratio after 5 days immersion in PBS, SBF and FBS-SBF. 130
900 30d SBF 30d PBS 800
700
600 500
400
300
Elastic Modulus [kPa] Modulus Elastic 200
100 0 -0.001 0 0.001 0.002 0.003 0.004 0.005 0.006 MPS molar ratio to NIPAAm
Figure 6-14: In vitro elastic modulus of PNIPAAm-PEGDM with various MPS molar ratio after 30 days immersion in PBS and SBF.
131
7. Future Work
This thesis work meant to be the first step to evaluate the feasibility of two composites as orthopedic devices. Some of the things that could enhance the findings of this study and current understanding of the composites are summarized below.
1. Bioactive solid composite: HA/PMMA
The aim of the addition of a coupling agent in the HA/PMMA interface is to promote the in vitro mechanical properties of the composite with minimal effectiveness on its surface bioactivity.
• In vitro cyclic mechanical test on the composites to better simulate the in situ bending
and compressive loading environment as mandible reconstruction up to 1 millions
load cycles.
• Since protein adsorption can greatly affect surface bioactivity of biomaterials, the in
vitro mechanical properties of the HA/PMMA composites can be highly altered in a
protein included simulated body fluid.
• Vacuum dried the composites prior to immersing in silane coupling agent can
improve the interfacial hydrogen bonding between the HA and silane. Therefore, both
global and interfacial elastic mechanical properties of the silane treated specimens
can be enhanced.
2. Injectable bioactive hydrogel: PNIPAAm-MPS-PEGDM
The challenge of this material system is to balance the network-forming and bioactivity inducing MPS with the gain in elastic recovery induced by PEGDM addition to the PNIPPAm base, all while maintaining an injectable material system. Additional 132 permutations of this system may likely lead to even higher modulus materials with increased bioactivity.
• The bioactivity can be improved with an organic bioactive material that has higher
solubility than calcium chloride. Higher solubility of calcium salt results in an
increase of supersaturation of the surrounding ion contained fluid; hence, an
accelerated apatite formation within the polymer matrix can be achieved. Also, a
combination of organic materials can be added into the systems to enhance the
bioactivity and mechanical properties. One suggestion is adding a mixture of CaCl2
salt and hydroxyapatite powder into the PNIPAAm-MPS-PEGDM. In theory, the
addition of hydroxyapatite powder can slightly enhance the in vitro mechanical
properties of the gel.
• Define the optimized combination of PNIPAAm and PEGDM as a function of the
molecular weight of each polymer. With higher molecular weight of both polymers,
the elasticity of the systems may increase while the polymer may become less rigid.
• A higher degree of cross-linking may be achieved with a different cross-linker, such
as acrylic acid, that can lead to improvements in the mechanical properties and
bioactivity of the gel without affecting its injectability.
• An alternative rigid polymer and hydrophobic in SBF at 37°C, such as polyethylene
oxide, can be applied to enhance the mechanical properties and retain the phase
transition of the PNIPAAm based systems.
133
8. Novel Contributions
This thesis work investigated two innovative ceramic / polymer composites as osteo-implant devices. Novel contributions of each material to the literature can be summarized as follows:
1. Bioactive solid composite: HA/PMMA
• Developed a method to determine the interfacial elastic properties of a bioactive
composite before and after immersion in a simulated body fluid. The interfacial
mechanical behaviors of the HA/PMMA composites were more sufficient in
analyzing the effect of coupling agent treatments and can applied to predict the global
elastic properties.
• Improved the interfacial elastic properties of HA/PMMA composite using a co-
polymer coupling agent. The co-polymer coupling agent (PMMA-MAA) showed an
improved interfacial and global Young’s modulus even after immersion in SBF for 3
days in comparison to the untreated composites. The surface treatment of the
bioactive fillers delayed the surface bioactivity of the fillers that helped to retain its
initial in vitro interfacial mechanical properties.
• Observed the long term bioactivity of a three dimensional printed HA/PMMA
composite bar. By closely arranging the bioactive HA particles, calcified bridges
bioactive material can be developed and sheltered the adjacent non-bioactive
polymer.
2. Injectable bioactive hydrogel: PNIPAAm-MPS-PEGDM
• Formulated a bioactive thermo-sensitive hydrogel through chemically bonded
calcium onto the polymer backbone, which are capable of apatite nucleation and 134
proliferation in vitro. With this technology pipeline, other polymers can also be
molecularly engineered into a bioactive polymeric system.
• Evaluated the elastic modulus of a thermo-sensitive hydrogel based material in
various biological fluids with respect to immersion times. The result showed that the
presence of ions in the media can influence the in vitro mechanical properties and
phase transformation behavior of the hydrogel materials.
• Determined the influence of a silane cross-linker on the in vitro mechanical properties
of a thermo-sensitive hydrogel. An optimized cross-linker content can maximize the
mechanical properties while retain the phase transformation behavior of the system.
Also the swellability of a cross-linker added hydrogel would attain equilibrium with
the immersion solution within hours. 135
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Vita
Emily received a Bachelor of Applied Science with Honors in Metals and
Materials Engineering from University of British Columbia in June 1998. After graduation, she worked in Hong Kong Productivity Council as a process engineer from
September 1998 to December 2000. She also was a member of the company’s Year 2000
Best Work Improvement Team.
In March 2001, Emily joined Materials Science and Engineering Department of
Drexel University. She quickly was involved in research and has focused her work on materials intended to replace bone tissue. Her work has led to a graduate student research award at Drexel University’s Research Day competition, a summer student research grant from the American Academy of Implant Dentistry, a Drexel University Dean’s fellowship, a Drexel University George Hills’ fellowship and a student professional development award from Society for Biomaterials. Emily has presented her work at national conferences and one of her work has been accepted by the journal, Dental
Materials. In addition, Emily’s talents lie in her awards for Drexel University Business
Plan Competition and SFU International Business Plan Competition. Through a series of written and oral rounds of review, Emily’s team was awarded second place for Drexel
University’s competition and also was the final 25 out of 200 teams completed in the
SFU international competition. Moreover, Emily is a recipient of Drexel University’s
Excellence in Mentorship in summer 2004 as a guidance of high school teachers’ research projects.