Related applications

The present application claims the benefit of Australian provisional patent application number 2013900475 filed on 14 February 2013, the disclosure of which is incorporated herein by reference.

Field of the Invention

The present invention relates to the synthesis and application of novel organic-inorganic composite molecularly-coupled "hybrid" materials. The hybrids of the invention have applications in wide-ranging technologies such as optics, electronics, mechanics, membranes, protective coatings, catalysis, sensors, and in particular as biomaterials. These hybrids embody the advantageous properties of both organic and inorganic frameworks such that the chemical coupling between these components yields properties that are not readily achievable in conventional composite materials such as thermoplastics and bioglasses, and mere physical admixtures thereof.

The invention has been developed primarily for use in medical applications such as biocompatible bone or tissue replacements. Although the invention will be described hereinafter with reference to this application, it will be appreciated that it is not limited to this particular field of use. For instance, the inventive materials of the invention can be fabricated from a wide range of different types of polymers including natural, synthetic, degradable and non- degradable. Different types of structures can be produced such as monoliths, fillers, coatings and injectable pastes/solutions which are hardenable in situ. In addition, the bioactivity of polymeric implants can be improved by the impregnation of these bioactive materials into such implants, or providing a surface coating thereon to improve biocompatibility and mechanical properties. The hybrid materials of the invention can also formulated as drug delivery devices capable of delivering hydrophobic or hydrophilic pharmaceutical compounds.

Background to the Invention

The following discussion of the prior art is provided to place the invention in an appropriate technical context and enable the advantages of it to be more fully understood. It should be appreciated, however, that any discussion of the prior art throughout the specification should not be considered as an express or implied admission that such prior art is widely known or forms part of common general knowledge in the field.

Bioceramics such as "Bioglass®" are favourable materials for bone grafting applications due to a relatively high biocompatibility and bonding affinity to the host tissue of the mammalian body via the formation of a biologically active hydroxyl carbonate apatite (HCA) layer on their surfaces. This HCA phase is chemically and structurally equivalent to the mineral phase in bone, thereby facilitating interfacial bonding. A thin layer of apatite forms on the glass-tissue interface, facilitating strong bond to the bone. Some

formulations can also facilitate growth of osteoblasts through the material. It will be appreciated that reference to "bioglass" throughout the present application is intended to be non-limiting, and any biocompatible glass source is within the scope of the present invention.

Bioglass® is a commercially-available family of bioactive glasses, comprising Si0 ₂ , Na ₂ 0, CaO and P ₂ 0 ₅ in specific proportions. Bioglasses differ from traditional soda-lime glasses due to the relatively low amount of silica (i.e., less than 60 mol.%), relatively high amount of sodium and calcium - and relatively high calcium : phosphorus ratio. The ratio of calcium to phosphorus promotes the formation of apatite crystals; calcium and silica ions can act as crystallisation nuclei.

Bioglasses have many different formulations. Some bind to soft tissues and bone (e.g., 45S5), some to bone only (e.g., 5S4.3 or Ceravital), some do not form a bond at all and after implantation become encapsulated with non- adhering fibrous tissue, and others are completely resorbed within few weeks.

It has been demonstrated in the past that the mol% of silica in bioglasses has a significant effect on their bioactivity. For example, bioglasses having <35 mol.% Si0 ₂ are said to be non glass-forming. Those having >50 mol.% Si0 ₂ , <10 mol.%) CaO, <35 mol.%> Na ₂ 0 are said to be bioactive, and undergo resorption within 10-30 days. Finally, bioglasses having >65 mol.%> Si0 ₂ are non-bioactive, nearly inert and become encapsulated with fibrous tissue. Despite its favourable biological profile, one of the major obstacles to the widespread adoption of bioglass in bone repair therapies has been its relatively poor mechanical properties, i.e. high Young's modulus and low fracture resistance, which make these materials brittle. Also, the bending strength of bioglass is typically within the range of 40-60 MPa, which is insufficient for load-bearing applications.

Several attempts have been made to resolve the issues associated with the brittleness and fragile structure of bioglass prior to bone grafting {see, e.g., Peltola, et al., Eur. Arch. Otorhinolaryngol. , 2012, 269, 623-8). For instance, introducing materials with a relatively low elastic modulus (e.g., polymer matrices) is a strategy used to improve certain mechanical properties of the bioglass and thereby more closely mimic the natural structure of bone (see, e.g., Hamizah, et al., J. Appl. Polym. Sci, 2012, 125, E661-E669). However, despite such measures, bioglasses generally provide an imperfect solution to the physical properties required of a bone repair/bone regeneration substrate.

One polymer which has been investigated in the past for biomedical purposes is poly(methyl methacrylate) ("PMMA"), which is a transparent thermoplastic and has excellent mechanical properties. In orthopaedic surgery, PMMA bone cement may be used to affix implants and to remodel lost bone. PMMA-based bone cements were first introduced in the 1960s (see, e.g., Charnley, J. Bone Joint Surg. Br., 1960, 42-B, 28-30) and in the interim, progress has been made toward improving the properties of such materials for the fixation of implants and filling bone voids after trauma or the removal of tumours. For example, in orthopaedic and dental operations, a mixture of methyl methacrylate ("MM A") monomer, initiators and activators is placed in the body and fixation occurs by in situ polymerisation. This approach, however, is fraught with several issues that include damage to/necrosis of the surrounding tissues as a result of temperature increase (c. 70°C) from the resultant exothermic reaction; release of toxic compounds such as residual MMA monomer, initiators and activators; inertness and lack of bioactivity, which leads to the thickening of intervening fibrous tissue layer; inhibition of cell function and growth/differentiation; increased inflammatory response; and cell death/necrosis, which can lead to the loosening of prosthesis/implants.

Another major consideration when using PMMA cement is the effect of stress shielding. Since PMMA has a Young's modulus greater than that of natural bone, the stresses are loaded into the cement and thus, the bone no longer receives the mechanical signals to continue bone remodelling. This, in turn, can result in bone resorption.

Several attempts have been made to resolve the problems of PMMA- based implants, such as minimising the amount of MMA monomer and replacing it with PMMA polymer. The release of the MMA monomer is not only an issue when implanted in the body; there are also concerns about surgeons being exposed to both MMA and benzoyl peroxide ("BPO") which is commonly used as initiator while working with bone cements.

Despite the many advantages of using either bioglass or PMMA in biomedical applications, there are significant obstacles to their widespread adoption. Bioglass and PMMA have been physically combined in the past, and several of the abovementioned problems addressed. For example, embedding PMMA into bioglass structure improves certain mechanical properties and reduces brittleness, which in turn provides a material useful in biomedical applications such as bone implants. For example, bioglass has been added in the form of powdered filler to PMMA-based biomaterials to enhance its biocompatibility and bioactivity. The amount and particle size of the bioglass powder has a significant effect on the bioactivity of the resultant composite. Decreasing the particle size of the bioglass increases the available surface area of the bioglass, which results in increased exposure of the bioactive compounds to the surroundings. However, a physical mixture of PMMA and bioglass in solid form results in a non-homogenous distribution, and a lack of adhesion between the phases which consequently enhances the risk of implant failure. Furthermore, a non-homogenous mixture of PMMA and bioglass results in inconsistent physical properties such as degradation rate and mechanical strength, which is an issue inherent in many commercially-available materials. Furthermore, addition of bioceramic particles such as bioglass to injectable formations of PMMA cements results in needle blockage, which makes the injection difficult and painful for the patient. Also, it results in inhomogeneous filling of the defected site.

It is an object of the present invention to overcome or ameliorate at least one of the disadvantages of the prior art, or to provide a useful alternative. Summary of the Invention

The present invention provides chemical bonding between organic and inorganic compounds in organic-inorganic molecularly-coupled "hybrid" materials, which exhibit synergistically-enhanced properties. Chemical coupling (e.g. covalent bonding) of organic-inorganic materials can be distinguished from the mere physical mixtures of organic and inorganic materials, as discussed in the background above. The bioactive glass-polymer hybrid materials of the invention are composites consisting of two constituents which are bonded at the nanometer or molecular level. The present invention provides nanoscale interpenetrating networks of the bioglass and polymer which have covalent coupling between them. In some embodiments, this involves careful control of the chemistry of the sol-gel process. The materials of the present invention have potential in wide-ranging technologies such as optics, electronics, mechanics, membranes, protective coatings, catalysis, sensors and biomaterials. In some preferred embodiments, the polymer is PMMA, however in alternative embodiments, the polymer may be poly(vinyl alcohol), chitosan, or other biologically compatible polymers, and combinations thereof, to produce bioactive glass hybrid scaffolds for biomedical applications.

The physico-chemical properties of the materials of the invention can be "tuned" by adjusting the relative ratios of the individual components, and thereby provide an opportunity to exploit the various properties of both organic and inorganic frameworks, as well as manipulate their functional characteristics by varying parameters such as composition, procedure, coupling and the degree of dispersion of each component within the other. Moreover, the uniform intermixing of the organic and inorganic phases to provide an interpenetrating network yields unique properties that are as yet unachievable in conventional composites, or even nanocomposites.

The organic-inorganic hybrids of the invention additionally address the issue of phase separation, which is a common problem in the prior art, by creating interaction at a molecular level. The organic and inorganic components are mixed in domain sizes of nanometers, and the molecular interactions between the phases lead to the formation of two distinct classes of hybrid.

"Class I" hybrids include a mixture in which there is a weak interaction between the organic and inorganic phases, such as van der Waals forces or hydrogen bonding, whereas "class II" is characterised by hybrids having covalent or iono- covalent bonds, which impacts upon the final material properties. The bonding between phases in the organic-inorganic hybrids of the invention can be class I and/or class II.

The present invention relates to organic-inorganic hybrids, wherein the inorganic component is bioglass, and the organic component comprises a polymer, pre-polymer, or an oligomer, or combinations thereof. A hybrid material is formed when there is chemical coupling of the organic and inorganic components, and in one preferred embodiment a coupling agent is used to covalently couple the organic and inorganic phases together and provide an interpenetrating network of each component.

Preferably the coupling agent is an organic moiety attached to the organic component which allows chemical bonding of the organic component to the inorganic network. In one preferred embodiment, an organosilane group can be provided on the organic component which acts as a "network modifier", since in the final structure the inorganic network is only modified by the organic group. A preferred organosilane coupling agents is 3-(trimethoxysilyl)propyl methacrylate ("MPMA"), which comprises both organic and inorganic moieties, thereby providing means to covalently bond the respective organic and inorganic components. In this example, it will be appreciated that the organic polymer component contains a certain percentage of MPMA. However, in an alternatively embodiment, the polymer can be grafted with trimethoxysilyl functionality. Whilst a preferred network modifier or coupling agent is MPMA, it will be appreciated that other coupling agents can be utilised, as discussed further below.

In preferred embodiments, the polymer is functionalised with moieties which provide covalent coupling to the bioglass. However, in other

embodiments it will be appreciated that the bioglass can be functionalised with moieties which enable covalent bonding to the polymer component.

One particularly useful method to prepare the hybrid organic-inorganic materials of the invention is via "sol-gel" synthesis. This method is

advantageous due to its low-temperature processing, which is a favourable method for the incorporation of an organic compound into the structure of an inorganic substance without thermal degradation. Other suitable methods of synthesis will be known to the skilled person. Furthermore, the sol-gel method is efficient in creating a relatively homogeneous distribution between the two phases as a result of dissolving both organic and inorganic compounds in a common solvent. Preferably organic molecules are mixed with metal alkoxides and/or organosilanes which undergo sequential hydrolysis and condensation reactions and trap and/or bond to the organic components in the gel structure. As an example, the hydrolysis and condensation reactions of tetraethyl orthosilicate ("TEOS"), a common precursor for the fabrication of silica networks, are shown in Equations (1) to (3) and illustrates the formation of a Si- O-Si network.

Hydrolysis:

≡Si(OR) + H ₂ 0 <→≡Si(OH) + ROH (1)

Condensation:

≡Si(OR) + (OH)Si≡ <→≡Si-0-Si≡ + ROH (2) ≡Si(OH) + (OH)Si≡ <→≡Si-0-Si≡ + H ₂ 0 (3)

In one preferred embodiment, the sol-gel method may be used to prepare polymer-bioglass hybrids. Preferably the polymer is PMMA. In one

embodiment, the presence of an organosilane coupling agent {e.g., MPMA copolymerised into the PMMA) results in the formation of Si-O-Si covalent bonds to the bioglass silica network, and results in the formation of a relatively homogenous product. The molecular structure of PMMA-co-MPMA-bioglass hybrid is shown in Figure 1 A of the accompanying drawings. A conceptual representation of the hybrid material shown in Figure 1 A is shown in Figure IB. The present inventors have utilized analytical methods such as Fourier Transform Infrared (FTIR), 1H, ¹ C and ²⁹ Si solid-state NMR spectroscopy to study the microstructure of the hybrids of the invention at the molecular level. Furthermore, the nanoscale interactions between the polymer component (e.g. PMMA) and bioglass have been investigated using two-dimensional solid state NMR analytical techniques. The results of these analyses have been used to determine the effect of molecular interaction on the physical properties of such materials, and to optimize the composition of the hybrids with a view to preventing or reducing phase separation.

According to a first aspect, the present invention provides a composite material comprising a covalently coupled polymer-bioglass.

It will be appreciated that the covalently coupled polymer-bioglass is a "hybrid" composite material, meaning that the polymer and bioglass are covalently coupled to each other on the molecular scale. In preferred embodiments the polymer-bioglass are covalently coupled. However, the skilled person will appreciate that other chemical bonds can be utilised to molecularly couple the polymer and bioglass. The constituents of the bioglass are chosen such that the bioglass is adapted to form an apatite layer on the bioglass surface. Preferably the hybrid materials of the invention mimic all the properties of bone, are bioactive, osteoconductive and biocompatible.

Bioglass component

Bioactive glasses of silicate composition, which were first developed by Hench and co-workers in 1969 (Hench LL, Splinter RJ, Allen WC, Greenlee TK. Bonding mechanisms at the interface of ceramic prosthetic materials. J. Biomed. Mater. Res. 1971;5(6): 117-141), represent a group of surface reactive materials which are able to bond to bone in physiological environment (Hench LL. Bioceramics. J. Am. Ceram. Soc. 1998;81(7): 1705-1728.). Bioactive glasses suitable for the present invention consist of a silicate network incorporating sodium, calcium and phosphorus in different relative proportions. The classical 45 S5 bioactive glass composition is universally known as

Bioglass® (composition in wt%: 45% Si0 ₂ , 24.5% Na ₂ 0, 24.5% CaO and 6% P ₂ 0 ₅ ) and is a preferred bioglass composition for the present invention. Fabrication techniques for bioactive glasses include both traditional melting methods and sol-gel techniques. The typical feature common to all bioactive glasses, being melt or sol-gel derived, is the ability to interact with living tissue forming strong bonds to bone (and in some cases soft) tissue, a property commonly termed bioreactivity or bioactivity. The bonding to bone is established by the precipitation of a calcium-deficient, carbonated apatite surface layer on the bioactive glass surface when in contact with relevant physiological fluid or during in vivo applications.

The skilled person will appreciate that the term bioglass is intended to mean those forms of glass which are biologically compatible. Glasses are defined as solid materials with large enormous structural disorder or liquid materials with large viscosity values, whereas bioglasses are considered as a class of bioceramics with extensive applications in biomedical engineering and bone replacement materials. It will be appreciated that there are a number of types of bioglass, and that any bioglass will be suitable for use in the invention. Non-limiting examples of other bioglasses are those which contain other ions such as Zn, P, Mg, Sr, F, Fe, B, K, Na, etc, and combinations thereof.

The preferred precursors to the bioglass inorganic component are tetraethyl orthosilicate and tetramethyl orthosilicate, trimethyl orthosilicate, and combinations thereof. However, it will be appreciated by the skilled person that other precursors can be used.

A typical composition for a bioglass suitable for the present invention comprises:

Si0 ₂ 60-70 weight-%,

Na ₂ 0 5-20 weight-%,

CaO 5-25 weight-%,

MgO 0-10 weight-%,

P ₂ 0 ₅ 0.5- 5 weight-%,

B ₂ 0 ₃ 0-15 weight-%,

A1 ₂ 0 ₃ 0-5 weight-%) and

Li ₂ 0 0-1 weight-% Polymer component

In one preferred embodiment, the polymer is an organic polymer, and is preferably PMMA. However, it will be appreciated that other polymers, pre- polymer or oligomers can be utilised in the composite materials of the invention.

Suitable polymers have the following polymer compositions:

• statistically polymerized monomers (e.g. monomers A and B polymerized into ABBAABAB);

• alternating polymerized monomers (e.g. monomers A and B polymerized into ABABABAB);

• gradient (tapered) polymerized monomers (e.g. monomers A and B

polymerized into AAABAABBABBB);

• block copolymers (e.g. monomers A and B polymerized into

AAAAABBBBBB) wherein the block length of each of the blocks is 2, 3, 4, 5 or even more repeat units;

• graft copolymers (graft copolymers consist of a polymeric backbone with side chains attached to the backbone); and

• mixed forms of these polymers, e.g. blocky gradient copolymers.

The polymers suitable for the invention may have different polymer architectures including linear, comb/branched, star, dendritic (including dendrimers and hyperbranched polymers). A general review on the architecture of polymers is given by ODIAN, George, Principles of Polymerization, 4th, Wiley-Interscience, 2004. p. 1-18.

Comb/branched polymers have side branches of linked monomer molecules protruding from various central branch points along the main polymer chain (at least 3 branch points).

Star polymers are branched polymers in which three or more either similar or different linear homopolymers or copolymers are linked together to a single core.

Dendritic polymers comprise the classes of dendrimers and hyperbranched polymers. In dendrimers, with well-defined mono-disperse structures, all branch points are used (multi-step synthesis), while hyperbranched polymers have a plurality of branch points and multifunctional branches that lead to further branching with polymer growth (one-step polymerization process).

The polymers may be prepared via addition or condensation type polymerizations. Polymerization methods include those described by ODIAN, George, Principles Of Polymerization, 4th edition, Wiley-Interscience, 2004, p. 39-606.

Addition polymerization methods include free radical polymerization (FRP) and controlled polymerization techniques. Suitable controlled radical polymerization methods include:

• RAFT: reversible addition- fragmentation chain transfer:

• ATRP: atom transfer radical polymerization

• MADIX: reversible addition-fragmentation chain transfer process, using a transfer active xanthate;

• Catalytic chain transfer (e.g. using cobalt complexes);

• Nitroxide (e.g. TEMPO) mediated polymerizations:

Other suitable controlled polymerization methods include:

• GTP: group transfer polymerization;

• Living cationic (ring-opening) polymerizations;

• Anionic co-ordination insertion ring-opening polymerization; and

• Living anionic (ring-opening) polymerization.

Suitable examples of monomers for synthesising the polymer component include: acrylic acid, methacrylic acid, maleic acid (or their salts), maleic anhydride, alkyl(meth)acrylates (linear, branched and cycloalkyl) such as methyl(meth)acrylate, n-butyl(meth)acrylate, tert-butyl(meth)acrylate, cyclohexyl(meth)acrylate, and 2-ethylhexyl(meth)acrylate; aryl(meth)acrylates such as benzyl(meth)acrylate, and phenyl(meth)acrylate,

hydroxyalkyl(meth)acrylates such as hydroxyethyl(meth)acrylate, and hydroxypropyl(meth)acrylate; (meth)acrylates with other types of

functionalities (e.g. oxiranes, amino, fluoro, polyethylene oxide, phosphate substituted such as glycidyl (meth) acrylate, dimethylaminoethyl (meth)acrylate, trifluoroethyl acrylate, methoxypolyethyleneglycol (meth)acrylate, and tripropyleneglycol (meth)acrylate phosphate; allyl derivatives such as allyl glycidyl ether; styrenics such as styrene, 4-methylstyrene, 4-hydroxystyrene, 4- acetostyrene, and styrene sulfonic acid: (meth)acrylonitrile; (meth)acrylamides (including N-mono and Ν,Ν-disubstituted) such as N-benzyl (meth)acrylamide; maleimides such as N-phenyl maleimide; vinyl derivatives such as vinyl alcohol, vinylcaprolactam, vinylpyrrolidone, vinylimidazole, vinylnapthalene, and vinyl halides; vinylethers such as vinylmethyl ether; vinylesters of carboxylic acids such as vinylacetate, vinylbutyrate, and vinyl benzoate. Typical condensation type polymers include polyurethanes, polyamides, polycarbonates, polyethers, polyureas, polyimines, polyimides, polyketones, polyester, polysiloxane, phenolformaldehyde, urea-formaldehyde, melamine- formaldehyde, polysulfide, polyacetal or combinations thereof.

Various initiators can be utilised to polymerize the monomer(s). The initiator can be a thermal initiator or a photo-initiator. Thermal initiator(s) suitable for use in the invention include tertamyl peroxybenzoate, 4,4-azobis(4- cyanovaleric acid), 1, 1' azobis(cyclohexanecarbonitrile), 2,2'- azobisisobutyronitrile (AIBN), benzoyl peroxide, 2.2-bis(tert- butylperoxy)butane, 1, l-bis(tertbutylperoxy)cyclohexane , 1, l-bis(tert- butylperoxy)cyclohexane, 2,5-bis(tert-butylperoxy)-2,5-dimethylhexane, 2,5- bis( tert-butylperoxy)-2,5-dimethyl-3 -hexyne, bis(l-( tert-butylperoxy)- 1- methylethyl)benzene, 1 , 1 -bis( tert-butylperoxy)-3.3 , 5-trimethylcyclohexane, tert-butyl hydroperoxide, tert-butyl peracetate, tert-butyl peroxide, tert-butyl peroxybenzoate, tert-butylperoxy isopropyl carbonate, cumene hydroperoxide, cyclohexanone peroxide, dicumyl peroxide, lauroyl peroxide, 2,4- pentanedione peroxide, peracetic acid and potassium persulfate.

Corsslinking monomers can also be utilised.

Certain resorbable polymers may be alternatively, or additionally, be used in the invention, such as: polylactides (PLA), poly-L-Iactide (PLLA), poly-DL-Iactide(PDLLA); polyglycolide (PGA); copolymers of glycolide, glycolide/trimethylene carbonate copolymers (PGA/TMC); poly (lactide ethylene oxide fumarate), other copolymers of PLA, such as

lactide/tetramethylglycolide copolymers, lactide/trimethylene carbonate copolymers, lactide/d-valerolactone copolymers, lactide/s-caprolactone copolymers, L -lactide/D L -lactide copolymers, glycolide/L-Iactide copolymers (PGA/PLLA), polylactide-co-glycolide; terpolymers of PLA, such as lactide/glycolide/trimethylene carbonate terpolymers, lactide/glycolide/ ε- caprolactone terpolymers, PLA/polyethylene oxide copolymers;

polydepsipeptides; unsymmetrically 3,6-substituted poly-1 ,4-dioxane-2,5- diones; polyhydroxyalkanoates, such as polyhydroxybutyrates (PHB); PHB/b- hydroxyvalerate copolymers (PHB/PHV); poly-b-hydroxypropionate (PHP A); poly-p-dioxanone (PDS); poly-dvalerolactone - poly-8-caprolactone, poly(E- caprolactone-DL-Iactide) copolymers; methylmethacrylate-N-vinyl pyrrolidone copolymers; polyesteramides; polyesters of oxalic acid; polydihydropyrans; polyalkyl-2-cyanoacrylates; polyurethanes (PU); polyvinylalcohol (PVA); polypeptides; poly-b-malic acid (PMLA); poly-b-alkanoic acids;

polycarbonates; polyorthoesters; polyphosphates; poly(ester anhydrides); and mixtures thereof; and natural polymers, such as sugars, starch, cellulose and cellulose derivatives, polysaccharides, collagen, chitosan, fibrin, hyalyronic acid, polypeptides and proteins.

Mixtures of any of the above-mentioned polymers and their various forms may also be used. It will be appreciated that each of the polymers listed above are functionalised with suitable silane functionality adapted to covalently link the bioglass and the polymer components, or include a suitable co- polymerised silane-containing monomer, such as MPMA.

Certain non-resorbable polymers may be used in the invention, such as: polymethyl methacrylate, poly methacrylate, poly butyl acrylate, and combinations thereof. Again, it will be appreciated that each of the polymers listed above are functionalised with suitable silane functionality adapted to covalently link the bioglass and the polymer, or include a suitable co- polymerised silane-containing monomer, such as MPMA.

Preferred natural polymers are: collagen, chitosan, fibrin, hyalyronic acid, polypeptides and proteins. Preferred resorbable polymers are: polylactides (PLA), poly-L-Iactide (PLLA), poly-DL-Iactide(PDLLA); polyglycolide (PGA); copolymers of glycolide, glycolide/trimethylene carbonate copolymers (PGA/TMC); poly (lactide ethylene oxide fumarate), copolymers of PL A and polyhydroxyalkanoates.

Preferred non-resorbable polymer are: polymethyl methacrylate, poly methacrylate, and poly butyl acrylate.

Preferred molecular weights of the polymer are between 2,000 and 1,000,000 g.mol ^"1 .

Coupling agent

It will be appreciated that a variety of coupling agents may be used to covalently bind the polymer and bioglass components of the composition together to make a hybrid material. The polymer component may be functionalised either by introducing functionality to the polymer or cross linking a suitable monomer having the desired functionality.

Non-limiting examples of coupling agents are: MPMA, (3-aminopropyl) triethoxysilane, 3- glycidoxypropyldimethoxymethylsilane, acetoxytri-tert- butoxysilane and combinations thereof. Coupling agents with the general formula are also suitable for use with the invention.

where RO is a hydrolyzable group, such as methoxy, ethoxy,

or acetoxy, and X is an organofunctional group, such as amino,

methacryloxy, epoxy, etc.

In other embodiments, the polymer and the bioglass are suitably functionalized such that the functionalisation on each component can be covalently linked to molecularly couple together the polymer and the bioglass. Uses of the composite 'hybrid' material of the invention

The hybrid materials of the invention may be formed into a variety of shapes and used for various medical purposes. For example a medical device formed from a hybrid of the invention is preferably chosen from the group consisting of: a 3D implantable scaffold, an orthopaedic implant for

reconstructive surgery, a dental implant/prostheses, a spine implant, implants for craniofacial reconstruction and alveolar ridge augmentation, for cartilage regeneration, an osteochondral defect implant, a strut, a stent or a stent-graft. However, it will be appreciated that there are many other devices which would be within the purview of the present invention.

The hybrid materials of the invention may be implanted into a body in different ways, including, but not limited to subcutaneous implantation, implantation at the surface of the skin, implantation in the oral cavity, use as sutures and other surgical implantation methods. In other embodiments, the hybrid materials of the invention may be injected subcutaneously.

In some preferred embodiments, the hybrid materials of the invention may be coated with at least one resorbable polymer material, non-limiting examples of which include polyglycolides, polydioxanones,

polyhydroxyalkanoates, polylactides, alginates, collagens, chitosans, polyalkylene oxalate, polyanhydrides, poly(glycolide-co-trimethylene carbonate), polyesteramides, or polydepsipeptides etc. Alternatively, the coating material may comprise healing promoters such as thrombosis inhibitors, fibrinolytic agents, vasodilator substances, anti-inflammatory agents, cell proliferation inhibitors, and inhibitors of matrix elaboration or expression. Examples of such substances are discussed in U.S. Patent No. 6, 162, 537. The present invention also contemplates using a polymer coating, (e.g. a resorbable polymer) in conjunction with a healing promoter to coat the implantable medical device, for example according to the reference [Wu C. Acta

Biomateilia, 2008; 4:343-353].

In one embodiment, preferably the hybrid materials of the invention are a fully synthetic bone graft substitute. Due to its interconnected pores, the material serves as an ideal osteoconductive scaffold and supports the formation of new host bone. As highlighted herein, many of the advantages of the new materials of the invention can be summarised as follows:

• Optimized porosity.

• Enhanced bone ingrowth and vascularisation.

• Avoids potential problems common for grafting methods .

• Is formable to almost any shape to suit the application.

• Easy to use.

• Combines with autologous bone marrow or blood. • Displays accelerated and enhanced osteointegration.

The uses of the present invention are many fold, including:

• For bone void fillings or augmentation in zones requiring cancellous rather than cortical bone.

• For the filling of bone defects after trauma, reconstruction, or

correction in non-load or load-bearing indications.

• For trauma and orthopaedics: Filling of voids caused by cysts or osteotomies, filling of defects arising from impacted fractures, refilling of cancellous bone harvesting sites, arthrodesis and nonunions.

• For spine surgery: Postero-lateral fusion, interbody fusion (as cage- filling material), vertebrectomies (as filling material of the vertebral implants), refilling of bone graft-harvesting sites

• For cranio-maxillofacial surgery: Reconstruction of mandibular defects and sinus lifts.

• For promoting neural regrowth as well as bone regrowth, e.g. in spinal applications.

The hybrid materials of the invention may also be formed into ribbons or fibres.

By incorporating and optimizing the coupling agent (e.g. MPMA) into the structure of the polymer-bioglass hybrid:

1. ) Gelation time was reduced from 5 days to 5 hours.

2. ) The issue of phase-separation between the polymer and bioglass was addressed at the molecular level.

3. ) Homogenous morphological, biological and physico-chemical properties were obtained for the molecularly-coupled hybrid.

4. ) The hybrid materials of the invention were shown to have promising properties for bone replacement applications.

By manipulating the process parameters such as temperature, solvent and addition of catalyst (sodium bicarbonate) for condensation of silica:

1.) The gelation time of samples was reduced to only 3 minutes, 2.) The samples were produced via a biocompatible method, i.e.

utilisation of biologically-compatible solvents, such as ethanol rather than THF, and utilisation of biologically-compatible sodium bicarbonate catalyst rather than hydrochloric acid (HF).

The combination of the features above, enables the hybrids of the invention to be utilised as injectable bone cement and interconnected porous scaffolds. The hybrid materials of the invention are mechanically strong and biologically active, which enables them to be utilised as an alternative to currently used bone cements.

Composite 'hybrid' material properties

In some embodiments, the hybrid materials of the invention can be used as a porous tissue engineering scaffold. Preferably, the scaffold has a porosity of at least 60 %, more preferably at least 80 %, and most preferably at least 90 %. The hybrid of the invention may be formulated and prepared to have an interconnected porosity of 20, 25, 30, 35, 40, 45, 50, 55, 60, 65, 70, 75, 80, 85, 90 or 95%.

In one embodiment, the pore sizes are between about 75 to about 300 μιη. However, it will be appreciated that the hybrid materials of the invention could be configured to have lower or greater pore size according to the intended or desired use, and any pore size would be within the purview of the present invention. For example, pore sizes of 5, 10, 20, 30, 40, 50, 60, 70, 80, 90, 100, 110, 120, 130, 140, 150, 160, 170, 180, 190, 200, 210, 220, 230, 240, 250, 260, 270, 280, 290, 300, 310, 320, 330, 340, 350, 360, 370, 380, 390, 400, 410, 420, 430, 440, 450, 460, 470, 480, 490, or 500 micron are possible.

The compressive strength of the hybrid materials of the invention are preferably between about 2 to 20 MPa, but can be greater than this depending on the choice of polymer and bioglass, and their relative concentrations.

Preferably the hybrid materials of the invention have a biocompatibility when placed in physiological fluid. Preferably the hybrid materials of the invention form a hydroxyapatite layer upon exposure to bodily fluids. As the skilled person will appreciate, the formation of hydroxyapatite is widely recognised as strong evidence that the body accepts the material as sui generis and is a requirement for the implant to chemically bond with living bone and tissue. The material of the invention has favourable cell-interaction properties. For example, two types of osteoblast cells (primary human osteoblast cells and mouse clonal osteoblast cells) adhere to and proliferate on the hybrid material. Preferably the material of the invention induces mineralization for cultured cells.

Preferably the hybrid material of the invention has homogenous morphological features. Preferably the hybrid material of the invention has hardness similar to the hardness of natural bone. Preferably the hybrid material of the invention shows no cytotoxicity in in vivo tests.

According to a second aspect, the present invention provides a method of production of a composite material comprising the steps of:

preparing an inorganic sol of bioglass precursors;

preparing an organic sol of polymer;

combining said inorganic sol and said organic sol under such conditions to produce a covalently coupled polymer-bioglass composite material.

Preferably the inorganic sol of bioglass precursors is selected from tetraethyl orthosilicate ("TEOS"), trimethyl orthosilicate, and combinations thereof.

Preferably the polymer is a sol of poly(methyl methacrylate)-co-3-

(trimethoxysilyl)propyl methacrylate ("PMMA-co-MPMA"). However, the polymer may by resorbable, or non-resorbable, and can be chosen from the polymers listed above.

Preferably bioglass precursors and polymer(s) are dissolved or suspended in the same solvent. Preferred solvents are selected from:

tetrahydrofuran (TFIF), ethanol, N-methyl-2-pyrrolidone ( MP), and acetic acid.

In a preferred embodiment, the inorganic sol is prepared by hydrolysing TEOS in low pH aqueous solution (pH between 1 to 6) and with the addition of calcium chloride or any other suitable calcium salt at room temperature. The aqueous solution can be acidified with hydrochloric acid. Preferably the TEOS:water molar ratio is between 1 :4 to 1 :8, for example 1 :5, 1 :6, or 1 :7. In another preferred embodiment, the predetermined volume ratio of organic sol to inorganic sol is polymenbioglass = 60:40 vol% or 80:20 vol%. Preferably the polymer is not greater than 80 vol.%.

Preferably the ratio of polymer to bioglass precursors is between 1 : 1 to 1 :6. Preferably the ratio is 1 : 1, or 1 :2, or 1 :3, or 1 :4, or 1 :5 or 1 :6.

The reaction can proceed between couple of seconds up to several minutes. However, the gelation and drying procedure may take from hours to days.

Preferably, the PMMA-co-MPMA is prepared from free radical polymerisation of PMMA and MPMA in a predetermined molar ratio, using α,α'-azoisobutyronitrile ("AIBN") as an initiator. The predetermined molar ratio of MPMA to PMMA is most preferably between about 0.1 to 0.5 molar ratio. Preferably the molar ratio of MPMA:MMA is 0.1 to avoid undesired crosslinking between the copolymer chains. A ratio lower ratio than 0.1 results in reduced, poor or a lack of formation of hybrid Si-C bonds in the resulting composite material. The polymerization reaction can take anywhere from 1 to 12 hours and be conducted at temperatures between 60 to 100°C.

In some preferred embodiments the sols are at least partially gelled. In other embodiments the hybrid material is dried and sintered.

According to a third aspect of the present invention there is provided an organic-inorganic hybrid composite when synthesised by a method according to the second aspect.

According to a fourth aspect, the present invention provides a formulation for preparing a composite material, the formulation comprising a bioglass sol, a polymer sol having functionality adapted to covalently couple the bioglass and polymer. The polymer is a copolymer of MPMA:MMA at a 0.1 molar ratio, and wherein the polymer composition is above 60 vol. %.

According to a fifth aspect, the present invention provides a kit comprising in separate containers: a bioglass sol, a copolymer sol and a catalyst.

The covalent coupling between the bioglass and copolymer sols is catalysable by the addition of sodium bicarbonate to produce a solidified hybrid composite material. Other catalysts will be known to the skilled person. In some preferred embodiments the bioglass sol and copolymer sol are in a separate containers, and in other embodiments are in the same container.

In another aspect, the invention provides a method of forming a composite material in situ or in vivo, the method comprising the steps of delivering to a subject in need thereof a polymer sol and a bioglass sol to a predetermined location in the body of the subject and in situ catalysing the formation of a hybrid bioglass-polymer composite material. The predetermined location can be subcutaneous, for example, at a site in the body in need of a scaffold for bone growth, such as a region of bone damage or weakness.

The polymer sol and bioglass sol can be administered sequentially or preferably simultaneously.

In one embodiment to prepare a biocement in situ, the method preferably comprises mixing the bioglass sol and copolymer sol at elevated temperature. Preferably the elevated temperature is between 70 and 95°C, and is preferably 70°C. The mixing time can be seconds to hours, and is preferably 30 seconds. The heated solution is then cooled to room temperature and the catalyst solution added. The resulting catalysed liquid is then drawn into a syringe and injected subcutaneously into the defected site.

In one embodiment to prepare porous scaffolds, the method preferably comprises mixing bioglass and polymer sols at elevated temperature obtain a homogenous and well-dispersed solution. Preferably the elevated temperature is between 50 and 95°C, and is preferably 70°C. The mixing time can be seconds to hours, and is preferably 30 seconds. The hybrid is then cooled to room temperature and sodium bicarbonate is added a foaming agent to create porosity within the scaffold. Other foaming agents are well known to the skilled person. The temperature of the solution is then elevated to form bubbles from degradation of sodium bicarbonate and to stabilize the porous structure during gelation of hybrid composite. Preferably the elevated temperature is between 50 and 95°C, and is preferably 70°C. The resulting porous scaffold (hybrid bioglass-polymer composite) is then dried, preferably at room temperature.

According to a sixth aspect, the present invention provides the use of a hybrid composite of the invention as a bone filler, as an injectable formulation, or as a porous scaffold for biomedical purposes. Preferably the hybrid of the invention is shaped into the desired application by controlling the gelation time using process parameters such as temperature and solvent and also, the catalyst (sodium bicarbonate) for the condensation reaction of compounds.

Preferably the composite material forms a hydroxyapatite layer upon exposure to bodily fluids.

Use of the composite material of the invention as a biocompatible material.

An implantable medical device comprising the composite material of the invention. Preferably the medical device is formed into a device chosen from: a 3D implantable scaffold, an orthopaedic implant for reconstructive surgery, a dental implant/prostheses, a spine implant, implants for craniofacial

reconstruction and alveolar ridge augmentation, for cartilage regeneration, an osteochondral defect implant, a strut, a stent and a stent-graft. The medical device may be permanently implanted or temporarily implanted. In some embodiments the medical device is substantially biodegradable. In some embodiments the medical device has a porosity of between about 10 to about 80%, a pore size is between about 20 to about 500 micron. In one embodiment the medical device is coated with at least one resorbable polymer material selected from polyglycolides, polydioxanones, polyhydroxyalkanoates, polylactides, alginates, collagens, chitosans, polyalkylene oxalate,

polyanhydrides, poly(glycolide-co-trimethylene carbonate), polyesteramides, and polydepsipeptides. In one embodiment the medical device is coated with at least one healing promoter selected from thrombosis inhibitors, fibrinolytic agents, vasodilator substances, anti-inflammatory agents, cell proliferation inhibitors, and inhibitors of matrix elaboration or expression.

Bone implant or biocement comprising the hybrid composite of the invention.

A method for producing an implantable medical device comprising: transferring the hybrid composite of the invention onto a substrate thereby forming said implantable medical device. An implantable drug delivery device comprising the hybrid composite of the invention.

A method for improving the long term stability of an implantable medical device comprising the step of: coating said device with the hybrid composite of the invention.

Use of the hybrid composite of the invention in the regeneration or resurfacing of tissue, comprising contacting the tissue with a quantity of the hybrid composite of the invention for a sufficient period to at least partially effect said regeneration or resurfacing.

A method for regenerating or resurfacing tissue, comprising the step of: contacting said tissue with the hybrid composite of the invention.

A method for forming osseous tissue on an orthopaedic defect, comprising the step of: contacting said defect with the hybrid composite of the invention.

Brief Description of the Figures

Preferred embodiments of the invention will now be described with reference to the accompanying Figures, in which:

Figure 1 A is a schematic molecular structure of a copolymer of PMMA- co-MPMA-bioglass class II hybrid material, and Figure IB is a schematic representation of Fig. 1 A

Figure 2 show the 400 MHz 1H NMR spectra of a) P _Co nt, b) PLOW, c) P _Co nt expansion in the region of 3.5-4.1 ppm, and d) P _LOW expansion in the region of 3.5-4.1 ppm in deuteriochloroform (CDC1 ₃ ). The presence of peaks at 3.59 ppm (a, -OCH3), 1.014 ppm (b, -CH ₂ ) and 1.621 (c, -CC¾) in 1H NMR spectrum confirmed the complete synthesis of Pc ₀ nt (PMMA);

Figure 3 is the ¹ C NMR spectrum of the P _Me d PMMA-co-MPMA copolymer of the present invention. The methyl and carbonyl carbons have been assigned as indicated. The spectrum included the characteristic peaks of carboxyl group (178 ppm), methylene (56 ppm), methoxy (52 ppm), quaternary (45 ppm) and methyl carbons (17 ppm). These data confirm the successful copolymerisation of PMMA and MPMA; Figure 4 is the FTIR spectra for a) P _Co nt, b) P _L0 w, c) P _Me d and d) P _H i _g h- The peaks at 1722 cm ^"1 and 1452 cm ^"1 were assigned to a C=0 symmetric stretching and CH ₂ bending, respectively. The peaks at 2944 cm ^"1 and 2983 cm ^{" 1} were assigned to the stretching mode of a C-H bond. The peak observed at 1070 cm ^"1 was attributed to the Si-O-C bending and C-0 vibrations were detected at 1150 cm ^"1 ;

Figure 5 depicts optical images of hybrid samples a) Hc _on t6o, b) H _H i _g h6o and c) Hcont4o- It was observed that by increasing the molar ratio of MPMA in the copolymers structure, firmer and more transparent monoliths were formed;

Figure 6 is the ¹ C CPMAS NMR spectra of a) neat PMMA (P _Co nt), b)

PMMA with high coupling agent (PHigh) and c) H _H i _g h6o hybrid. These spectra were acquired with high power 1H decoupling. The spinning side bands are marked with and asterix (see, Fig.6(a)). Broad yet well-resolved C=0, OCH ₃ , CH ₂ and CH ₃ resonances of PMMA backbone are visible. In addition, in the case of functionalised PMMA in Pffigh, the -0-CH ₂ - peaks of the coupling agent MPMA chains were also visible at 67 ppm;

Figure 7(a) is the simplified schematic structure of hybrid in the absence of Ca ²⁺ . Q and T are parameters that describe the hybrid network connectivity. Q ⁿ shows a silicon bonds to n other silicons via "bridging" oxygen atoms, and the T sites indicate a silicon atom bonds to one carbon atom (Si-C) and to other silicons;

Figure 7(b) shows the ²⁹ Si CPMAS NMR spectra of hybrid samples set against bioglass. The Q- and T-site silicons are assigned as indicated. Three main resonances are observed at -110, -100 and -90 ppm that correspond to the Q ⁴ , Q ³ and Q ² species, respectively. T ⁿ sites are only visible in the H _H i _g h6o and H _H i _g h8o hybrids at -50, -60 and -65 ppm. The presence of these peaks confirmed the formation of Si-C bonds, and hence demonstrate the contribution of polymer into the bioglass network structure. Therefore, it may be concluded that at least a molar ratio of MPMA:PMMA=0.1 is required for fabrication of PMMA- bioglass Class II hybrid;

Figure 8(a) is a Directly-Polarised ²⁹ Si HPDMAS NMR of PMMA- bioglass hybrids. The acquired spectra are shown in bold while deconvoluted peaks fits are shown as thin lines. The Si peaks were deconvoluted to quantify the bioglass structure in the composites;

Figure 8(b) shows the relative populations for the different Q sites in the hybrid materials. As can be seen, in the neat bioglass and hybrids in the absence of coupling agents (bioglass and Hc _on t6o), the relative amount of Q ⁴ species was c. 15-22 % higher as compared to hybrids with high coupling agent (H _H i _g h6o and H _H i _g h8o), while hybrids with low and medium coupling agents (H _LoW 6o and H _Me d6o) possessed intermediate concentrations of Q ⁴ species. These data demonstrate that the coupling agent has a direct impact on the condensation reaction of silicate framework;

Figure 9 shows the 2D 1H- ¹ C and 1H- ²⁹ Si HetCor for H _H i _gh 6o. The ID 1H projections are plotted to the left of the 2D spectra while the ¹ C and ²⁹ Si projections are plotted at the top. As expected from the nature of bioglass, one observes strong ²⁹ Si correlations to 1H species at c. 0.5 ppm and 4.5 ppm.

These signals are assigned to isolated silanol OH species and weakly-adsorbed H ₂ 0 moieties. Additionally, a strong correlation of 1H species at c. 2.5 ppm corresponding to the methylene protons of PMMA with the bioglass ²⁹ Si was observed. The observation of a molecular correlation between the PMMA and the bioglass by MR implies close molecular contact between a significant fraction of the PMMA molecules and the bioglass; this confirms the formation of a nanoscale composite in the hybrid material;

Figure 10 shows the FTIR spectra for a) Hc _on t6o, b) H _H i _g h6o, c) Pcont, and d) bioglass. The peaks at 1030-1040, 930 and 790 cm ^"1 corresponded to Si-O-Si network structure and were detected only for bioglass and H _H i _g h6o- The peaks at 1625 cm ^"1 and 3370 cm ^"1 correspond to the presence of C0 ₃ ^2" and OH groups, respectively, of the bioglass composition. The absence of C=0 peak at 1722 cm ^"1 and the presence of small Si-O-C peak at 1110 cm ^"1 for Hc _on t6o confirmed the formation of weak van der Waals force between the oxygen in PMMA (C=0) and silicon ion owing to high electropositivity of Si ⁴⁺ ions. This intermolecular interaction resulted in the formation of a Class I hybrid of PMMA and bioglass;

Figure 11 shows SEM images for a) H _Co nt6o and b) H _ffig h6o. The H _Co nt6o morphology was non-homogeneous and the phase separation between the bioglass and polymer was visible on the micron scale. However, in the chemically-coupled hybrid (H _H i _g h6o) a homogenous and smooth surface was detected without any indication of large scale phase separation. These images confirm that phase separation was avoided by formation of covalent bonding between PMMA and bioglass;

Figure 12 shows the TGA (a, top) and DTG (b, bottom) curves of a) PMed, b) bioglass, c) H _Co nt6o, d) H _Me d6o, e) H _H i _g h6o. In the TGA profile of neat PMMA-co-MPMA copolymer, one observes three distinct peaks at 165 °C, 270 °C and 360 °C. These peaks are attributed to the presence of head-to-head linkages, end-chain unsaturation and random scission within the polymer chain, respectively. In the DTG profile of bioglass, a single peak at 100 °C is due to the loss of water;

Figure 13 shows the DSC curves of a) Pcont, b) Pffigh, c) pure bioglass, d) Hcont6o, e) H _H i _g h6o- The peak at 120 °C which was observed in bioglass DSC profile was attributed to the condensation of bioglass silica network. The same trend was also observed in Hc _on t6o and H _H i _g h6o curves. However, this peak was shifted toward higher temperatures (i.e., to 150 °C) in H _H i _g h6o due to the covalent bonding of polymer chains to silica network. Therefore, inorganic moieties reduced the polymer chains mobility and the covalent bonds between the phases reduced the free movement of both organic and inorganic chains. Upon increasing the coupling between phases, more energy (heat) was required in hybrid samples to approach the same degree of freedom. It could then be concluded that covalent bonding of silica to PMMA chains may have enhanced the thermal stability of the final product.

Figure 14 is a table showing optical transparency and gelation time of various samples.

Figure 15 shows SEM images of a) bioglass, b) HO, c) HI . Phase separation between ceramic and polymer was evident for the physical mixture of PMMA and bioglass (HO). However, the absence of phase separation and homogenous surface structure for hybrid HI underlined the chemical conjugation and structural integrity of these samples; Figure 16 shows the EDS results of a) HO and b) HI including the EDS mapping results for carbon, silicon and calcium elements and the statistical results of the presence of elements. Three random points were selected in each sample and the average amount of calcium was quantified. The results demonstrate that while the distribution of calcium in HO was dramatically changed, it possessed homogeneous distribution with narrow range variation for HI;

Figure 17 shows STEM images of a) HO and b) HI, with a 1 μπι scale ruler shown at the bottom left hand corner of each image. As shown, the creation of separate phases is clearly noticeable for physical mixture (HO), while HI acquired from covalent bonding was homogenous even in the nanoscale range. These data suggest that the presence of MPMA and covalent bonding between PMMA and bioglass had a paramount impact on physical integrity and creation of a homogenous composition of these two compounds;

Figure 18 shows the Atomic Force Microscopy (AFM) results of hybrids including topographic images, phase imaging and roughness (Ra) of the surface of the HO and HI for 5 x 5 μπι area. As shown, the physical mixture (HO) possessed a non-uniform surface corroborating the appearance of separate phase of PMMA and bioglass. However, the surface of the HI hybrid was

homogenous and there is no evidence of phase separation both at low and high resolutions. In addition, the surface roughness of HO was six- fold higher than HI samples. The results of AFM also confirmed that addition of MPMA as a coupling agent is efficient in creation of a homogeneous hybrid from a polymer and bioglass;

Figure 19 relates to the mechanical properties of bioglass and HI samples, a) A stress-strain diagram, b) a microhardness analysis (p<0.001). The Young's (elastic) modulus of the HI hybrid obtained from the stress-strain diagram is 40-times higher compared to the bioglass samples;

Figure 20 shows images of a) bioglass compared to b) the HI hybrid after free-fall from 45 cm height on a wooden surface. The formation of cracks in the structure of the bioglass was clearly visible, but the integrity of hybrid HI was maintained; Figure 21 depicts the bioactivity of the a) bioglass, b) HI samples after 7 days incubation in SBF at 37°C. A 10 μιη scale ruler is shown to the bottom left hand corner of each image. The results show that the apatite layer was formed on the surface of both bioglass and hybrid. These data demonstrate that the presence of 80 vol.% PMMA in the structure of PMMA-bioglass hybrid had a negligible impact on bioactivity of the hybrid. Therefore, these hybrids potentially show similar biological properties to bioglass, yet, are mechanically stable enough to be used as bone replacement materials;

Figure 22 shows SEM images of (a, b) pure PMMA (P ₀ ); (c, d) HO; (e, f) HI; and (g, h) bioglass. Both the physical mixture (HO) and chemically conjugated samples (HI) supported cell attachment. However, this attachment was enhanced significantly in HI . A homogenous filopodia cell migration was observed on the surface of HI, which promotes cells anchoring to the biomaterial surface and enhances cell spreading. Furthermore, for HI samples, cells were spread well on the surface and base lamellipodia were well- developed;

Figure 23 is an MTS assay of the HI (H _H i _g h6o) hybrid sample compared with the PMMA and bioglass. As shown, a significant difference ( <0.001) was observed in the absorbance of hybrid samples (HI) compared to the pure PMMA and bioglass. These data demonstrate that covalent bonding between

PMMA and bioglass had a positive impact on cell interaction and provide a superior environment for cell adhesion and proliferation; and

Figure 24 is a degradation profile of the a) bioglass, b) HI sample.

Dashed lines indicate the linear model fitted to the curve. It was observed that the presence of PMMA covalently bonded to the bioglass decreased the rate of degradation of the bioglass. As shown, the degradation rate of the hybrid sample was 1.5 times less than the bioglass.

Figure 25 are AFM results including topographic images, phase imaging and roughness of the surface of the HO (physical mixture) and HI (hybrid).

Figure 26 are the mechanical properties of bioglass and HI samples, a)

Stress-strain diagram, b) Ultimate strain, c) Microhardness values (*** represents p < 0.001). Figure 27 shows degradation profiles of hybrid compared with pure PMMA and bioglass (The data was extrapolated for BG until day 83).

Figure 28 shows FTIR spectra of a) HI, b) residues of degraded HI. Figure 29 are bioactivity test shows the formation of HA on the surface of a) bioglass, b) HI, c) HO and d) PMMA samples after 7 days incubation in SBF.

Figure 30 are SEM images of a) Pure PMMA, b) HO, c) HI, d) bioglass. White arrows show the individual narrowed cells on the surface of samples.

Figure 31 is MTS assay of HI hybrid sample compared to the PMMA and bioglass.

Figure 32 shows ALP staining on materials showing ALP+ cells (blue) after 4 days of osteogenic differentiation. Positive staining was seen on HI hybrid and tissue culture plastic, but not on PMMA, bioglass, or HO material.

Figure 33 shows tissue reaction to the bioglass implant at day 10 after implantation, a) Azan- staining, the vessel-rich (arrows) multi-layered tissue wall (double arrow) covering the bulk-like structural segments of the implanted material (BG). A layer of mononuclear cells (squares) was adherent to the material surface, b) Macrophage-specific F4/80-immunostaining, macrophages (pentagons) were located in all regions of the peri-implant tissue (double arrow), c) TRAP-staining, only a low number of mononuclear cells within the peri-implant tissue expressed TRAP (diamonds), while the majority of these cells were TRAP-negative (black arrow heads).

Figure 34 shows tissue reaction to bioglass fragments at day 10 after implantation, a) H&E-staining, an overview of the implantation bed of fragments (BG), which were embedded within a vessel-rich (arrows) granulation tissue. At their surfaces mononuclear (pentagons) and

multinucleated cells (squares) were observable, b) Macrophage-specific F4/80- immunostaining, the involvement of phagocytic mono- and multinucleated cells (pentagons/squares) in the tissue reaction to the bioglass fragments (BG). c) TRAP-staining, the marked presence of multinucleated giant cells involved in the degradation of the bioglass fragments. These cells showed higher TRAP- enzyme expression potential. Figure 35 Tissue reaction to the HI implant at day 10. a) Macrophage- specific F4/80-immunostaining, an overview of the distribution of mononuclear cells within of the peri-implant tissue (double arrow) within the subcutaneous CT of the CDl mouse. Most of the cells at the surface as well as the periphery were identified as macrophages (pentagons), b) H&E-staining, a granulation tissue was located at the surface of the implant, c) TRAP- staining, only a low amount of cells at the surface of the material as well as within the material- adherent tissue (CT) showed signs of TRAP activity (diamonds). Most of the cells were TRAP-negative (black arrow heads).

Figure 36 shows the optical transparency of PMMA-co-MPMA and

PMMA-bioglass hybrid solutions below and at 70 °C.

Figure 37 shows the effect of temperature and solvent on the gelation time of PMMA-bioglass hybrids.

Figure 38 shows the effect of concentration of SB on the gelation time of HlEtOH at room temperature.

Figure 39 shows the effect of gelation time and network structure on the mechanical properties of hybrids, a) stress-strain curves, b) toughness value, c) ultimate stress, and d) ultimate strain.

Figure 40 are SEM images of hybrids after 7 days soaking in SBF, a) HI, b) HlEtOH, c) EDS analysis of crystals formed on the surface of hybrids.

Figure 41 show injectability of the hybrid solution (HlEtOH) by addition of SB.

Figure 42 provides the degradation profiles of HI and HlEtOH for the period of 100 days.

Figure 43 shows images of porous PMMA-bioglass scaffold, a) SEM micrograph, b) μ-CT scan of PMMA-bioglass porous scaffold in different cross- sections and the ratio of closed and open pores.

Figure 44 shows the mechanical properties of HlEtOH monoliths and porous structures - a) stress-strain curves, and d) ultimate strain.

Figure 45 provides the viability assay of HlEtOH samples in the form of monoliths and porous scaffolds (PS). Figure 46 provides the alizarin red S staining images of monoliths and porous scaffolds.

Figure 47 provides the quantified analysis of mineralization study of HIEtOH samples in the form of monoliths and porous scaffolds (PS).

Definitions

In describing and claiming the present invention, the following terminology will be used in accordance with the definitions set out below. It is also to be understood that the terminology used herein is for the purpose of describing particular embodiments of the invention only and is not intended to be limiting. Unless defined otherwise, all technical and scientific terms used herein have the same meaning as commonly understood by one having ordinary skill in the art to which the invention pertains.

Unless the context clearly requires otherwise, throughout the description and the claims, the words 'comprise', 'comprising', and the like are to be construed in an inclusive sense as opposed to an exclusive or exhaustive sense; that is to say, in the sense of 'including, but not limited to'.

Other than in the operating examples, or where otherwise indicated, all numbers expressing quantities of ingredients or reaction conditions used herein are to be understood as modified in all instances by the term 'about'. The examples are not intended to limit the scope of the invention. In what follows, or where otherwise indicated, '%' will mean 'weight %', 'ratio' will mean 'weight ratio' and 'parts' will mean 'weight parts'.

Notwithstanding that the numerical ranges and parameters setting forth the broad scope of the invention are approximations, the numerical values set forth in the specific examples are reported as precisely as possible. Any numerical value, however, inherently contains certain errors necessarily resulting from the standard deviations found in their respective testing measurements.

The terms 'predominantly' and 'substantially' as used herein shall mean comprising more than 50% by weight, unless otherwise indicated. The recitation of a numerical range using endpoints includes all numbers subsumed within that range (e.g., 1 to 5 includes 1, 1.5, 2, 2.75, 3, 3.80, 4, 5, etc.). The terms 'preferred' and 'preferably' refer to embodiments of the invention that may afford certain benefits, under certain circumstances.

However, other embodiments may also be preferred, under the same or other circumstances. Furthermore, the recitation of one or more preferred

embodiments does not imply that other embodiments are not useful, and is not intended to exclude other embodiments from the scope of the invention.

Headings of sections provided in this patent application and the title of this patent application are for convenience only, and are not to be taken as limiting the disclosure in any way.

Reference throughout this specification to "one embodiment" or "an embodiment" means that a particular feature, structure or characteristic described in connection with the embodiment is included in at least one embodiment of the present invention. Thus, appearances of the phrases "in one embodiment" or "in an embodiment" in various places throughout this specification are not necessarily all referring to the same embodiment, but may. Furthermore, the particular features, structures or characteristics may be combined in any suitable manner, as would be apparent to one of ordinary skill in the art from this disclosure, in one or more embodiments.

The terms used in this application, if not otherwise defined, are those agreed on at the consensus conference on biomaterials in 1987 and 1992, see Williams, OF (ed.): Definitions in biomaterials: Proceedings of a consensus conference of the European Society for Biomaterials, Chester, England. March 3-5, 1986. Elsevier, Amsterdam 1987, and Williams OF, Black J, Doherty PJ. Second consensus conference on definitions in biomaterials. In: Doherty PJ, Williams RL, Williams OF, Lee AJ (eds). Biomaterial-Tissue Interfaces.

Amsterdam: Elsevier, 1992. In this application, by bioactive material is meant a material that has been designed to elicit or modulate biological activity.

Bioactive material is often surface-active material that is able to chemically bond with the mammalian tissues. A biodegradable material is a material that breaks down in vivo, but with no proof of its elimination from body.

The term bioresorbable in this context means that the material is disintegrated, i.e. decomposed, upon prolonged implantation when inserted into mammalian body and when it comes into contact with a physiological environment. The by-products of a bioresorbable material are eliminated through natural pathways either because of simple filtration or after their metabolisation. The terms bioresorbable and resorbable can be used

interchangeably, but is it is clear that bioresorption is meant in this description. Especially, the term resorbable glass means silica-rich glass that does not form a hydroxyl-carbonate apatite layer on its surface when in contact with a physiological environment. Resorbable glass disappears from the body through resorption and does not significantly activate cells or cell growth during its decomposition process. By the term bioabsorbable it is meant a material that can dissolve in body fluids without any molecular degradation, and then excreted from the body.

By biomaterial is meant a material intended to interface with biological systems to evaluate, treat, augment or replace any tissue, organ or function of the body. By biocompatibility is meant the ability of a material used in a medical device to perform safely and adequately by causing an appropriate host response in a specific location, causing no foreign-body reactions and being non-toxic. By resorption is meant decomposition of biomaterial because of simple dissolution. By composite is meant a material comprising at least two different constituents, for example a polymer and a ceramic material, such as glass.

In the present context the term medical devices relates to any kind of implant used within the body, as well as devices used for supporting tissue or bone healing or regeneration. An implant according to the present context comprises any kind of implant used for surgical musculoskeletal applications such as screws, plates, pins, tacks or nails for the fixation of bone fractures and/or osteotomies to immobilize the bone fragments for healing; suture anchors, tacks, screws, bolts, nails, clamps, stents and other devices for soft tissue-to-bone, soft tissue-into-bone and soft tissue-to-soft tissue fixation; as well as devices used for supporting tissue or bone healing or regeneration; or cervical wedges and lumbar cages and plates and screws for vertebral fusion and other operations in spinal surgery. Detailed Description of the Invention

Experimental Details

1. Materials and Methods

1.1 Materials

Precursors required for the synthesis of PMMA-co-MPMA copolymer including MPMA, α,α'-Azoisobutyronitrile ("AIBN") and Ν,Ν'- dimethylformamide ("DMF") were purchased from Sigma- Aldrich and used as received. Methyl methacrylate ("MMA") purchased from Sigma was used after distillation under reduced pressure. Hydrochloric acid (Merck), tetraethyl orthosilicate ("TEOS"; Sigma-Aldrich), calcium chloride dihydrate

(CaCl ₂ .2H ₂ 0; Ajax Finechem Pty Ltd), tetrahydrofuran ("THF"; Merck) and deionised water were used for the fabrication of inorganic solution and hybrid.

1.2 Preparation of the inorganic "bioglass" solution

TEOS was mixed with deionised water and HC1 and stirred for 30 minutes followed by addition of calcium chloride dihydrate and stirred for another 30 minutes to yield a homogenously-mixed solution. The precursors were used with the molar ratio of TEOS: water :HC1: calcium carbonate dihydrate = 1 : 8 : 0.01 : 0.2 and this solution is referred to hereinafter as "sol (A)".

A common calcium source for the preparation of sol-gel derived bioglasses is calcium nitrate tetrahydrate. However, in this study, calcium chloride was used to minimise the risk of toxicity resulted from the nitrate byproduct.

1.3 Preparation of the organic "PMMA-co-MPMA " solution

Poly (methyl methacrylate)-co-3-(trimethoxysilyl)propyl methacrylate ("PMMA-co-MPMA") copolymer was synthesised by a free-radical polymerisation technique using AIBN as an initiator. Precursors were mixed in a Schlenk flask (MMA: AIBN=200 (molar ratio); DMF (20 mL)) and degassed by three sequential freeze-pump-thaw cycles. Polymerisation was conducted in an oil bath at 70 °C over 12 hours. The solution was then cooled to room temperature and the polymer was purified by precipitation in diethyl ether followed by filtration and drying in vacuum for a period of 2 days. The yield in each run was ~1 g (-70%). PMMA-co-MPMA was then dissolved in THF at a concentration of 10 wt.% and the solution is hereinafter referred to as "sol (B)".

1.4 Preparation of PMMA-bioglass hybrid monoliths using THF as the solvent for polymer

Different volume ratios of Sol(A) and Sol(B) were prepared and mixed by magnetic stirrer at room temperature to form a uniform mixture. The solutions were then kept sealed at 25 °C prior to gel formation followed by drying unsealed for 7 days at the same conditions. The samples were incubated at 37°C for 7 days and finally the monoliths were vacuum dried for 2 days at 40°C. This drying procedure was conducted to preserve the monolith structure during solvent evaporation and also to remove residues of solvents. The preliminary study demonstrated that using fast drying procedure for these samples resulted in disintegrated structure.

1.5 Preparation of PMMA-bioglass hybrid monoliths using ethanol as the solvent for polymer

Sol (B) (PMMA-co-MPMA dissolved in ethanol 10 wt%) and Sol(A) were mixed in the volumetric ratio of Sol(B):Sol(A) = 60:40 vol% by magnetic stirrer at 70 °C to form a transparent solution. The solutions were then immediately cooled down to room temperature by water bath. Sodium bicarbonate (SB) was added as a solution in miliQ water with a specific concentration to the hybrid solution. Gelation occurred at room temperature and dried for 24 hours at ambient temperature unsealed. It should be noted that fast gelation did not damage the monolithic structure of these samples.

2. Characterisation of the hybrid materials

2.1 ¹ H & ¹³ C nuclear magnetic resonance (NMR)

Selected synthesised polymers were characterised by 1H and ¹ C NMR on a Bruker Ultra Shield Avance (400 MHz) spectrometer using CDC1 ₃ as the solvent. The concentration of the samples was 20 g.L ^"1 .

2.2 ^} H, ¹³ C and ²⁹ Si solid state NMR

The prepared gel was characterised by solid state NMR analyses to determine the structure of solid network. All the NMR experiments were carried out on a Bruker Biospin Avance III solids-300 MHz and Avance II standard bore, 700 MHz spectrometers (Bruker-Biospin, Rheinstetten, Germany). A Bruker 4-mm double resonance magic-angle spinning ("MAS") probehead was used with MAS frequencies of 5 kHz for ¹ H- ²⁹ Si HetCor and 12 kHz for other experiments. The 90 degree pulse length was 4 for 1H and 4.5 for ²⁹ Si and ¹ C nuclei.

A 300 second recycle delay was used in direct-polarisation ²⁹ Si MR experiments. The SPINAL-64 ¹ H- ²⁹ Si and 1H- ¹ C heteronuclear decoupling strength was 80 kHz during ²⁹ Si and ¹ C detection. Hartman-Hahn cross- polarisation (HHCP) was employed for experiments requiring polarisation transfer from 1H to ¹ C or ²⁹ Si with a contact times of 1 ms and 4 ms, respectively. ¹ H- ²⁹ Si and 1H- ¹ C HetCor experiments were carried out with frequency- switched Lee-Goldburg (FSLG) for 1H-1H homonuclear decoupling and 1H chemical shifts were scaled by 0.571 accordingly during data processing. Also attempted were 160 ti increments of 51 μβ. The glycine ¹ C CO resonance at 176 ppm, the Kaoline ²⁹ Si resonance at -91 ppm and hydroxyapatite 1H resonance at 0.2 ppm were used to calibrate the ¹ C, ²⁹ Si and 1H chemical shifts, respectively. The neat bioglass, PMMA and hybrids were dried, powdered and c. 60 mg were packed into the NMR rotors for the respective measurements. 2.3 Gel Permeation Chromatography (GPC)

The synthesised copolymers were characterised for the weight average molecular weight (Mw), number average molecular weight (Mn) and

polydispersity index (PDI) (Mw/Mn) by GPC with a Shimadzu Prominence HPLC (Kyoto, Japan) equipped with a GPC column (Jordi Gel DVB mixed bed, Alltech) diode array detector (SPD-M20A, Shimadzu) and refractive index detector (RID- 1 OA, Shimadzu). The copolymers were dissolved at a

concentration of 1 mgmL ^"1 in HPLC grade THF and 10 vol.% DMF. The mobile phase was 10 vol.% DMF in THF and delivered at 1 mL.min ^-1 through the column that was kept constant at 50 °C. The GPC system was calibrated with PEG standards and the relative molecular weight of polymers was determined by comparing the retention times to that of standards using Waters Empower software. 2.4 Attenuated total reflection Fourier transform infrared (A TR-FTIR) The functional groups of the samples were analysed by ATR-FTIR spectroscopy (FTIR; Nicolet 6700, Thermo Fisher Scientific Inc.). The synthesised polymers and hybrid samples were characterised for their molecular structures at a scan speed of 32 scan/min.

2.5 Scanning electron microscopy (SEM)

The surface microstructures and phases formed on the specimens were scanned by field emission scanning electron microscopy (FE-SEM; Zeiss ULTRA plus). Samples were mounted on aluminium stubs using conductive carbon paint, then gold coated by using Emitech K7550X instrument prior to SEM analysis.

2.6 Scanning electron microscopy-Energy Dispersive Spectrometry (SEM- EDS)

The surface microstructures and crystal phase formed on selected specimens were analysed by field emission scanning electron microscopy (FE- SEM; Zeiss ULTRA plus). This instrument was equipped with Bruker XFlash 4010 EDS detector with high speed acquisition and hypermapping capability. Samples were mounted on aluminium stubs using conductive carbon paint, then gold coated by using Emitech K7550X instrument prior to SEM analysis.

2. 7 Scanning transmission electron microscopy (STEM)

STEM analyses were conducted to investigate the interaction between the phases in nanoscale. Powders were embedded in epoxy resins and microtomed with Leica Ultracut ultramicrotomes (UC7) for 100 nm layers. The layers were harvested and seated on carbon grids prior to STEM analysis (Zeiss ULTRA plus).

2. 7 Thermogravimetric Analysis (TGA)

Thermogravimetric analysis (TA Q-500) was carried out to study the thermal decomposition of the various samples. Analyses were conducted over the temperature range of 30 °C to 600 °C at heating rate of 20 ⁰ C.min ^_1 under a nitrogen atmosphere. 2.8 Differential scanning calorimetr (DSC)

The thermal properties of the polymers were analysed by DSC (TA Q- 1000, USA). The samples had an average weight of 4 mg and were heated from room temperature to 400 °C under constant nitrogen flow (50 mL.min ^"1 ).

Samples were sealed in the aluminum pans and heated at a rate of 10 "C.min ^"1 . Standard aluminum sample pans were used as the reference.

2.9 Atomic Force Microscopy (AFM)

The surface characteristics of the prepared hybrid samples were investigated by atomic force microscope (AFM; Asylum Research, MFP-3D- BIO) in AC mode. Phase imaging was used to identify different phases. A thin film was required to examine the surface properties.

Table 2 - Prepared copolymers of PMMA-co-MPMA

3. Results and Discussion

3.1 Synthesis of the PMMA-co-MPMA copolymers ("P")

PMMA-co-MPMA copolymers were synthesised with MPMA:PMMA molar ratios of 0, 0.004, 0.02 and 0.1 which were coded as control polymer (Pcont), polymers with low (PLOW), medium (PMed) and high (Pffigh) ratio of functional groups, respectively (see, Table 2). The suffix defines the ratio of functionalisation of a copolymer that corresponds to the ability for covalent bond formation with bioglass.

1H MR analysis was used to characterise the molecular structure of the synthesised copolymers. Fig.2(a,b) shows the 1H NMR results for Pc _ont and P _LOW in deuteriochloroform (CDC1 ₃ ). The presence of peaks at 3.59 ppm (a, - OCH3), 1.014 ppm (b, -CH ₂ ) and 1.621 (c, -CC¾) in 1H NMR spectrum confirmed the complete synthesis of Pc ₀ nt PMMA). The occurrence of a peak at 3.92 ppm (d, -(COOCH ₂ )) in Fig.2(b) demonstrated the successful copolymerisation of MPMA with MMA even at the lowest molar ratio of MPMA:PMMA = 0.004 (i.e., P _L0 w)- The respective absence and presence of peaks that are associated with MPMA in Pc _on t and P _LOW were confirmed on MR spectra and are shown in Fig.2(c, d).

The 1H NMR spectra of PMed and Pmgh showed the same characteristic resonances of MMA (3.59 ppm) and MPMA (3.92 ppm). The molecular weights of the synthesised copolymers, percentage of functional groups and number of moles of these groups per gram of each copolymer were calculated and are listed in Table 2.

The ¹ C NMR spectrum of PMMA-co-MPMA copolymers synthesised in this study is provided in Fig.3. The ¹ C NMR spectrum of PMed included the characteristic resonances of carboxyl group (178 ppm), methylene (56 ppm), methoxy (52 ppm), quaternary (45 ppm) and methyl carbons (17 ppm). These data confirm the successful copolymerisation of PMMA and MPMA. A similar spectrum was obtained for P _Low and Pmgh samples.

The molecular structures of the synthesised copolymers of PMMA-co- MPMA with different ratios of MPMA were confirmed by FTIR analysis. As depicted in Fig.4, peaks at 1722 cm ^"1 and 1452 cm ^"1 were assigned to a C=0 symmetric stretching and CH ₂ bending, respectively. The peaks at 2944 cm ^"1 and 2983 cm ^"1 were assigned to the stretching mode of a C-H bond. The peak observed at 1070 cm ^"1 was attributed to the Si-O-C bending and C-0 vibrations were detected at 1150 cm ^"1 . It was observed that by increasing the

MPMA:PMMA molar ratio from P _Low (i.e., 0.004) to P _H i _g h(/ ^' .e., 0.1), the Si-O-C peak at 1070 cm ^"1 was increased, which was in agreement with the 1H NMR spectral results, thereby confirming the contribution from the higher content of MPMA to the entire copolymers structure.

3.2 Synthesis of the organic-inorganic hybrids ("H")

Bioglass (sol (A)), having a composition of TEOS : water : HC1 :

calcium carbonate dehydrate = 1 : 8 : 0.01 : 0.2 and copolymer (sol (B)) solutions in THF (at 10 wt.%) were mixed in three different volumetric ratios of sol (B) : sol (A) = 40 : 60, 60 : 40 and 80 : 20 that was followed by co- condensation of the silanol groups of organic and inorganic components. The respective resultant solutions were coded as H40, ¾o and ¾o representing the volumetric ratios of polymer to bioglass solutions.

As an example, Hc _on t4o is the hybrid sample prepared from Pc _on t polymer and the H ₄₀ mixture of sol (B) : sol (A) (i.e., 40 : 60). These hybrids were examined visually, to determine the effect of composition of polymer solution to bioglass solution and the fraction of MPMA on the transparency, phase separation and gelation behaviour of samples.

Transparency is an indicative parameter for the absence of phase separation below the scale of 400 nm (see, e.g., Wei, et al., J. Appl. Polym. Sci., 1998, 70, 1689-1699).

The ratio of MPMA : PMMA played an important role in gelation and the formation of the transparent products. The samples that had their transparency and gelation time examined are listed in Table 3. It was observed that by increasing the molar ratio of MPMA in the copolymers, firmer and more transparent monoliths were formed. However, at low MPMA : PMMA molar ratio (e.g., <0.1) samples were less transparent and gelation time was greater than 20 days at ambient temperature (see, Fig.2(a,b)).

The composition of the bioglass (sol (A)) was another factor that elicited a significant impact on the phase separation, transparency and gelation time of the hybrids. As shown in Fig.5, a one-phase solution was acquired for Hc _on t6o and H _H i _g h6o- However, when using high proportion of bioglass (e.g., Hc _on t4o >50 vol.%), two separate phases were formed due to the agglomeration of the polymer chains (see, Fig.5(c)).

Table 3— The effect of composition on gelation time and optical properties

a gelation time after one hour mixing

OP: one phase formation; TP: two phase formation; T: transparent

NT: not transparent; NG: no gel formation before 20 days

The functionalisation of PMMA with MPMA had a significant impact on increasing the concentration of bioglass that can be used for hybrid formation with no phase separation {e.g., polymer agglomeration). For example, at low MPMA : PMMA molar ratio, the hydrophobic effect of PMMA inhibits covalent bonding and the creation of a class II hybrid with uniform distribution of the phases. Therefore, two separate zones of polymer and bioglass were formed for Hc _on t4o samples. However, the same composition of bioglass could be used in the fabrication of a Class II hybrid (H _H i _g h4o) with no phase separation at high MPMA concentration (MPMA : PMMA = 0.1).

Furthermore, it was observed that the gelation time of the H _H i _g h samples was tuned by varying the composition of the bioglass. The gelation time of the hybrid material was decreased from 20 days to about 2 hours at room temperature when the composition of polymer solution was increased from 40 vol.% to 80 vol.%.

The microhardness identification number for a hybrid sample containing the highest MPMA molar ratio (H _H i _g h8o) was more than 15 times higher than the pure bioglass. This result shows that addition of covalent bonding between organic-inorganic compounds not only accelerates the gel formation, but also dramatically enhances the mechanical properties of the resultant hybrids.

3.3 Molecular scale analyses

¹³ C CPMAS NMR results of the polymers P _Co nt and P _H i _g h; and of the hybrid H _H i _g h6o samples are shown in Fig.6. The spectra clearly demonstrate the ¹ C NMR resonances of the PMMA scaffolding material. Broad yet well- resolved C=0, OCH ₃ , CH ₂ and CH ₃ resonances of the PMMA backbone are visible. The broadening of ¹ C NMR peaks in the solid state can be directly assigned to the effects of conformational disorder within the amorphous PMMA polymer. In addition, in the case of functionalised PMMA in Pffigh, the -0-CH ₂ - peaks of the coupling agent MPMA chains were also visible at 67 ppm. Also observed was an increased broadening of the backbone -CH ₂ signal in the H _H i _g h6o hybrid material. Without wishing to be bound by theory, it is suggested that this effect may be due to polymer chain interaction with the relatively rigid bioglass components that underpin bioglass-polymer hybrid formation interaction at the nanoscale level.

The schematic illustration of the molecular structure of the proposed hybrid is shown in Fig.7(a). "Q" and "T" are parameters that describe the hybrid network connectivity. Q ⁿ shows silicon bonds to n other silicons via "bridging" oxygen atoms, and the T sites indicate a silicon atom bonding to one carbon atom (Si-C) and to other silicons {see, Mahony, et al., Adv. Funct.

Mater., 2010, 20, 3835-3845).

For example, T ³ shows that a silicon atom bonds to three silicon and one carbon atom. Three main peaks are observed at -110, -100 and -90 ppm in the qualitative ²⁹ Si CPMAS NMR in Fig.7(b), that correspond to the Q ⁴ , Q ³ and Q ² species, respectively. T ⁿ sites are only visible in the H _H i _g h6o and H _H i _g h8o hybrids at -50, -60 and -65 ppm for T ¹ , T ² and T ³ , respectively. The presence of these peaks confirms the formation of Si-C bonds - and hence, demonstrates the contribution of the polymer into the bioglass network structure.

Therefore, it may be concluded that at least a molar ratio of MPMA : PMMA = 0.1 is required for fabrication of PMMA-bioglass class II hybrid.

It is, therefore, anticipated that in H _H i _g h materials, there is a relatively high degree of interaction at a molecular level between PMMA and bioglass (e.g., a nanoscale interaction).

Directly polarised, single-pulse ²⁹ Si MR spectra with high-power 1H decoupling (HPDMAS) are shown in Fig.8. The resulting ²⁹ Si peaks were deconvoluted to quantify the bioglass structure in the composites using the DMfit software (see, e.g., Massiot, et al., Magn. Reson. Chem., 2002, 40, 70- 76). In comparison to the CPMAS spectra (see, Fig.7(b)), the ²⁹ Si single-pulse spectra exhibited a higher intensity of Q ⁴ sites, which was due to the nature of two experiments. In CPMAS, only those ²⁹ Si species at a close proximity of c. 0.5 nm from 1H sites were cross-polarised and yielded an NMR signal.

Consequently, the most condensed Q ⁴ silicon sites, which are the furthest away from 1H species, are under-represented in the CPMAS spectrum. However in the single-pulse experiment, the ²⁹ Si NMR signal was not dependent on proximity to 1H species and therefore yielded the full quantitative signal of all Si species.

The deconvolution of spectra in Fig.8(a) yielded the relative population of different silicon Q sites in the materials, which are tabulated in Fig.8(b). Q ⁴ sites are indicative of the degree of condensation of bioglass within the hybrids - and therefore, the micro-architecture of inorganic phase in these materials. As can be seen from Fig.8(b), in the neat bioglass and hybrids in the absence of coupling agents (i.e., bioglass and Hc _on t6o), the relative amount of Q ⁴ species was c. 15-22 % higher as compared to hybrids with high coupling agent (i.e., H _H i _g h6o and H _H i _g h8o), while hybrids with low and medium coupling agents (i.e., H _LoW 6o and H _Me d6o) possessed intermediate concentrations of Q ⁴ species.

These data demonstrate that the coupling agent has a direct impact on the condensation reaction of silicate framework. The lower populations of Q ⁴ species in the H _H i _g h6o and H _H i _g h8o hybrid materials demonstrate smaller domain sizes of the bioglass phase and potentially, a more open network structure as compared to the H _Low , Hc _on t and neat bioglass samples.

Again, without wishing to be bound by theory, it is thought that while the coupling agent acts as a site for the nucleation of bioglass condensation, the presence of PMMA polymer chains serves to hinder the formation of very large domains of bioglass. This mechanism is anticipated to yield a structure with closely interacting polymer and bioglass phases on the nanometer scale.

The 2D 1H- ¹ C and 1H- ²⁹ Si HetCor NMR spectra of the H _H i _gh 6o material conclusively identified the nanometer-scale interaction between the polymer and bioglass phases. In these 2D NMR spectra, the 1H chemical shifts correlating to ¹ C- ²⁹ Si chemical shifts were spread in a 2D map, with the intensities plotted as contour levels as shown in Fig.9. The 2D ¹ H- ¹ C HetCor was acquired with a short contact time of 100 of cross-polarisation transfer from the 1H to the ¹ C nuclei. As a result, this spectrum primarily shows the localised correlations of ¹ C species to their directly bonded 1H species. In the case of 2D ¹ H- ²⁹ Si HetCor, the spectrum was acquired with a relatively long contact time of 4 ms of cross polarisation transfer. This means that in addition to the two-bond Si-O-H correlation, this spectrum also shows correlations to protons within c. 1 nm proximity. As expected from the nature of bioglass, one observes strong ²⁹ Si correlations to 1H species at c. 0.5 ppm and 4.5 ppm.

These signals are assigned to isolated silanol OH species and weakly-adsorbed H ₂ 0 moieties.

Additionally, a strong correlation of 1H species at c. 2.5 ppm

corresponding to the methylene protons of PMMA with the bioglass ²⁹ Si was observed. Considering the fact that the hybrid material was not isotopically labelled and that NMR is an inherently insensitive technique, the observation of a molecular correlation between the PMMA and the bioglass by NMR implies a close molecular contact between a significant fraction of the PMMA molecules and the bioglass. This confirms the formation of a nanoscale composite in the hybrid material.

FTIR was also used to investigate the molecular structure of the prepared hybrids. The FTIR spectra of neat bioglass, PMMA {i.e., Pcont), hybrid of bioglass and highly functionalised copolymer {i.e., H _H i _g h6o), and the physical mixture of bioglass and polymer {i.e., Hc _on t6o) are shown in Fig.10 {i.e., (a) Hconteo, (b) H _H igh6o, (c) Pcont, and (d) bioglass). The peaks at 1030-1040 cm ^"1 , 930 cm ^"1 and 790 cm ^"1 corresponded to Si-O-Si network structure and were detected only for bioglass and H _H i _g h6o- These peaks were not observed in Hc _on t6o due to the lack of gel formation in this sample. The peaks at 1625 cm ^"1 and 3370 cm ^"1 correspond to the presence of C0 ₃ ^2" and OH groups, respectively, of the bioglass composition.

The absence of a C=0 peak at 1722 cm ^"1 and the presence of a small Si- O-C peak at 1110 cm ^"1 in the FTIR spectrum of Hc _on t6o confirmed the formation of weak van der Waals force between the oxygen in PMMA {i.e., a C=0) and silicon ion owing to the high electropositivity of Si ⁴⁺ ions. This intermolecular interaction resulted in the formation of a class I hybrid of PMMA and bioglass. In this method, components were mixed in the form of liquid, hence, the steric effect was decreased and the likelihood of inducing van der Waals forces between molecules was increased. On the other hand, the presence of C=0 in the H _H i _g h6o hybrids showed that the silicon ions contributed in Si-O-Si covalent bonds.

The FTIR results show that the amount of MPMA coupling agent played a critical role on the molecular interactions of compounds from weak to strong bonding that may result in the production of materials with unique physical properties such as morphological features and thermal stability.

3.4 Morphology

The surface morphology of Hc _on t6o and H _H i _g h6o hybrids was investigated by scanning electron microscopy (SEM). As shown in Fig. 11(a), the Hc _on t6o morphology was non-homogeneous and the phase separation between the bioglass and polymer was visible on the micron scale. However, in Fig. 11(b), for the chemically-coupled hybrid (H _H i _g h6o) a relatively homogenous and smooth surface was detected without any indication of large-scale phase separation. These images confirm that phase separation may be avoided by the formation of covalent bonding between PMMA and bioglass phases. 3.5 Thermal behaviour

The thermal behaviour of the synthesised material is an important characteristic from the standpoint of its potential applications - especially when a high temperature process is required, for example, in a sterilisation processes for bioimplants. It was anticipated that the nanoscale interaction between bioglass and polymer would impact significantly upon the characteristics of a hybrid material. As shown in Fig. 12(a), the TGA profile of neat PMMA-co- MPMA copolymer showed three distinct peaks at 165 °C, 270 °C and 360 °C. These peaks are attributed to the presence of head-to-head linkages, end-chain unsaturation and random scission within the polymer chain, respectively. The polymer was degraded completely by 400 °C due to the breakdown of its backbone structure.

The TGA profile of bioglass depicted in Fig. 12(b) shows a single peak at 100 °C, which is due to the loss of water. A steady state profile was observed for bioglass up to 600 °C because of thermal stability of the silica structure. .

The onset thermal decomposition temperature ("OTDT") of Hc _on t6o was increased compared with that of the pure polymer due to the addition of silica, as listed in Table 4. The formation of covalent bonding between organic and inorganic components had the same impact upon the thermal characteristics of the H _Me d6o and H _H i _g h6o hybrids. The degradation peaks were at 360 °C, 381 °C, 390 °C and 400 °C for P _Me d, H _Co nt6o, H _Me d6o, H _H i _g h6o, respectively. Thus, the OTDT of these samples was tuned to higher temperatures by increasing the MPMA composition which greatly impacted the molecular structure and level of chemical interaction.

The degree of covalent bonding showed a significant impact upon the thermal behaviour of products as shown in the deconvoluted thermal gravimetry (DTG) curve of the H _H i _g h6o (see, Fig.12). Hc _on t6o and H _Me d6o possessed three major distinct weight losses in TGA and DTG curves, which were in agreement with the degradation profile of the synthesised polymer. However, for the H _H i _g h6o sample, weight losses were detected only at two distinct temperatures (249 °C and 400 °C). This is thought to have occurred due to the presence of strong covalent bonds between PMMA chains and silica structure that inhibit thermal degradation within the range of 200 °C to 300 °C. The OTDT of H _H i _g h6o was also higher than other hybrid products having lower levels of coupling, such as H _Me d6o and H _Low 6o.

The degradation profile of the hybrid was also influenced by the presence of covalent bonding. A similar trend was observed for the degradation profile of the hybrid H _con t6o {see, Fig.12A(c) and Fig.12B(c)) and the polymer (P _Med ) {see, Fig.12A(a) and Fig.12B(a)) at temperatures below OTDT of the polymer {i.e., 360 °C). In this temperature range, the degradation of Hc _on t6o was greater than H _H i _g h6o- However, the amount of silica maintained at 600 °C for Hcont6o was higher than that of H _H i _g h6o- These results demonstrate that the polymer and silica in the physical mixture (Hc _on t6o) were degraded separately. However, in the Class II hybrid (H _H i _g h6o), degradation was uniform due to covalent bonding between PMMA and silica. These results were in agreement with the data obtained by Differential Scanning Calorimetric analysis (DSC) which are presented in Fig.13.

DSC analysis was conducted for pure bioglass, synthesised polymers {i.e., Pcont, Pffigh) and the hybrids {i.e., H _Co nt6o, H _H i _g h6o). The results in

Fig.13(a,b) show that the glass-transition temperature (T _g ) of pure PMMA was 105 °C, whereas no thermal effect was observed in DSC profile of highly functionalised PMMA (P _H i _gh ) below 250 °C. This effect might be due to the covalent bonding of MPMA-PMMA that limited the mobility of PMMA polymer chains. This effect decreases the degree of freedom of the PMMA chains and results in the absence of a T _g peak for the P _ff ig _h copolymer. Table 4 - Peaks of the thermal decomposition of samples and their S1O ₂ content

The peak at 120 °C which was observed in the bioglass DSC profile {see, Fig.13(c)) was attributed to the condensation of the bioglass silica network. The same trend was also observed in Hc _on t6o and H _H i _g h6o curves {see, Fig.13(d, e), respectively). However, this peak was shifted toward higher temperatures {i.e., to 150 °C) in H _H i _g h6o, presumably due to the covalent bonding of the polymer chains to the silica network. Therefore, inorganic moieties reduced the polymer chains mobility and the covalent bonds between the phases reduced the free movement of both organic and inorganic chains. Upon increasing the coupling between phases, more energy (heat) was required in the hybrid samples to approach the same degree of freedom. It could therefore be concluded that the greater the degree of covalent bonding of silica to PMMA chains, the greater the thermal stability of the final hybrid product.

4. Bioactivity analyses

4.1 Preparation of samples

Selected samples were characterised for their bioactivity by incubation in simulated body fluid ("SBF") at 37 °C {i.e., body temperature) for the periods of 1, 3, 7, and 14 days. After such time, the samples were washed twice with deionised water to remove any residue of minerals absorbed on the surface and dried in the oven at 37 °C for 24 hours. The dried monoliths were examined to determine their morphological characteristics and investigate the formation of any apatite layer on the surface.

4.2 Materials and methods

Primary human osteoblast ("HOB") was used to assess the cell interaction properties of the prepared samples. Complete Osteoblast Growth Media was used to culture the HOB cells which were incubated (Incubator Thermo Fisher Heracell 150i) at 37 °C in the presence of 5% C0 ₂ and 95% humidity. The media were refreshed every three days until the cells approached confluence. HOBs were detached from flask by trypsin treatment and collected for cell seeding.

The monolith samples (each, 12 mm diameter x 1 mm height) were placed into a well plate and kept in 70% ethanol for one hour to sterilise followed by rinsing with phosphate buffered saline ("PBS"), triple. Samples were then exposed to UV light for 30 minutes and were washed with fresh medium at 37 °C overnight. The substrates were placed into well tissue- culturing polystyrene plates and cells with the density of 2 ^χ 10 ⁵ cells/mL were seeded onto the monoliths and kept in a C0 ₂ incubator at 37 °C for four days.

Cell attachment was examined by investigating the morphology of the samples in SEM. Samples were fixed in 2.5% glutaraldehyde in 0.1 M PBS for 1 h and rinsed with PBS, triple. Samples were stained with 1% osmium tetroxides in 0.1 M PBS and incubated at room temperature for 1 h. After such time, they were dehydrated by replacing different ethanol solutions, (30%, 50%, 70%), 95%), 100%), sequentially). The specimens were dried in

hexamethyldisilizane overnight before gold-coating and SEM analysis for morphological determination.

Osteoblast cell proliferation was determined by the 3-(4,5- dimethylthiazol-2-yl)-5-(3-carboxymethoxyphenyl)-2-(4-sulfop henyl)-2H- tetrazolium (MTS) assay. Cell cultured samples were incubated for 4 days followed by rinsing three times with PBS and placed in a sterilised well plate. Fresh medium (250 μΐ.) and MTS agent (50 μΐ.) were added to each sample. The samples were then kept in a C0 ₂ incubator (Thermo Fisher Heracell 150i) at 37 °C for one hour to allow the MTS to react with the metabolically-active cells. Subsequently, water soluble formazan product produced by cell activity was quantified by the absorbance at 490 nm using a microplate reader (BioRad 680). Tissue culture polystyrene (TCPS) was used as a control in MTS assay. 4.3 Results and discussion

In this study, hybrid samples were prepared by the sol gel method in which an aqueous solution of bioglass (sol (A)) was mixed with a solution of PMMA (sol (B)) at 70 °C. As listed in Table 5, samples containing MPMA are coded as "HI" {i.e., "hybrids") and those without coupling agents represent the physical mixture and are coded as "HO". Pure bioglass was also considered as the control sample.

The transparency of the samples was examined since it is an indicative parameter for the absence of phase separation below the scale of 400 nm. Both HI and HO solutions were transparent at 70 °C. However, the HO sample became opaque upon decreasing the temperature below 65 °C. It was observed that the hybrid fabricated in presence of MPMA {i.e., HI) was transparent even at room temperature due to the formation of covalent bonds to the bioglass. Furthermore, it was observed that the gelation time of this hybrid was 30 minutes at 50 °C, whereas the physical mixture without MPMA (i.e., HO) did not form any gel even after 2 days. Another factor that played a critical role in gelation of the hybrid was the concentration of polymer. At low concentration such as 5 wt.% ("HI -5 wt.%"), no gel was formed due to the low number of covalent bonding in the hybrid samples. However, by increasing the concentration of polymer to 10 wt.% ("Hl-10 wt.%") a transparent gel was created.

One of the obstacles for the in situ application of hybrids is the low pH of the media. The present Inventors have now demonstrated that pH can be adjusted to that of biological body fluids (i.e., pH 7) by addition of 4 wt.% sodium bicarbonate to the hybrid solution (see, Fig.14). In addition, in the presence of a small quantity of sodium bicarbonate, the gelation time decreased dramatically to less than two minutes; this result renders the hybrid relatively suitable for in situ applications.

Table 5 - Physical properties including gelation time and optical status

NG: no gel formation; T transparent; NT: not transparent; OP: one-phase solution TP: two-phase solution

4.4 SEM-EDS and STEM analyses of hybrids as biocompatible materials

The surface morphology of neat bioglass, HO and HI hybrids were compared in Fig.15. Phase separation between the ceramic and polymer was evident in Fig.15(b) for the physical mixture of PMMA and bioglass (HO). However, the absence of phase separation and the presence of a relatively homogenous surface structure in Fig.15(c) for the hybrid (HI) underlined the chemical conjugation and structural integrity of these samples. This result was further confirmed by mapping the SEM image by EDS detector for determining different elements such as calcium, silicon and carbon.

As shown in Fig.16, calcium was distributed homogenously in the HI hybrid throughout entire surface. However, for HO, calcium was detected only on a few regions of the surface and the distribution was relatively non-uniform. In addition to the EDS mapping, three random points were selected in each sample and the average amount of calcium was quantified. The results shown in Fig.16 demonstrate that while the distribution of calcium in HO was changed dramatically, it possessed homogeneous distribution with narrow range variation for HI . These data were further supported by the results achieved from STEM analysis, demonstrating a significant difference between HO and HI samples.

As shown in Fig.17, the creation of separate phases is clearly noticeable for physical mixture (HO), while HI acquired from covalent bonding was relatively homogenous even in the nanoscale range. These data suggest that the presence of MPMA and covalent bonding between PMMA and bioglass had a paramount impact on the physical integrity and creation of a homogenous composition in these two compounds.

Atomic Force Microscopy (AFM) was used to compare the surface topography of hybrids fabricated in this study with a physical mixture of PMMA-bioglass. As shown in Fig.18, the physical mixture (HO) possessed a non-uniform surface corroborating the appearance of separate phases of PMMA and bioglass. However, the surface of the HI hybrid was homogenous and there was no evidence of phase separation both at low (5 μπι) and high (1 μπι) resolutions. In addition, the surface roughness of HO was six- fold higher than that of the HI samples (see, Fig.18(c, f, j)). The results of AFM also confirmed that the addition of MPMA as a coupling agent is efficient in creating a homogeneous hybrid from a polymer and bioglass and merging these two phases in the nanoscale range. 4.5 Mechanical testing of hybrid materials

Selected mechanical properties of the prepared hybrids were compared to neat bioglass samples to identify any potential applications. The Young's (elastic) modulus of the HI hybrid obtained from the stress-strain diagram (see, Fig.19) is 40-times higher compared to that of the bioglass samples.

Furthermore, a microhardness test showed that there is a significant (p< 0.001) increase in the numbers obtained for hybrid as opposed to bioglass. These results were clearly in agreement with visual observation (see, Fig.20) in which samples were dropped from a distance of 45 cm without any force (i.e., freefall) onto a wooden surface. The formation of cracks in the structure of the bioglass was clearly visible (see, Fig.20(a)); however, the integrity of hybrid HI was maintained during this test.

4.6 Biological activity of PMMA-bioglass hybrids

Prior to conducting biological assays, the bioactivity of various samples was assessed by formation of apatite layer when they were incubated in simulated body fluid (SBF). The results in Fig.21 show that the apatite layer was formed on the surface of both bioglass and hybrid after seven days of incubation. These data demonstrate that the presence of 80 vol.% PMMA in the structure of PMMA-bioglass hybrid had a negligible impact on the bioactivity of the hybrid. Therefore, these hybrids potentially show similar biological properties to bioglasses - yet, on the other hand, are mechanically stable enough to be used as bone replacement materials. To further clarify the cell-interaction properties, a cell study was conducted for this sample.

4. 7 Cell attachment, morphology and proliferation

In vitro tests were performed to determine the biological activity of the hybrids synthesised in this study. Primary human osteoblast cells were selected for conducting the biological assays to demonstrate the potential application of these hybrids for bone implant as filler. Previous studies had shown that cellular growth of osteoblasts on the interface of biomaterials is a function of the physical properties, microstructure and chemical composition of materials.

Samples HO and HI were selected for an in vitro cell study. The positive and negative controls (pure PMMA (P ₀ ) and bioglass samples) were used to determine the effect of a chemically-conjugated hybrid on osteoblast adhesion and spreading. As shown in Fig.22, both physical mixture (HO) and chemically conjugated samples (HI) supported cell attachment. However, this attachment was enhanced significantly in HI .

A homogenous filopodia cell migration was observed on the surface of

HI, which promotes cells anchoring to the biomaterial surface and enhances cell spreading. Furthermore, for HI samples, cells were spread well on the surface and base lamellipodia were well-developed, thereby corroborating the adopted flattened morphology of the cell growth.

Cell viability was measured by MTS analysis on the cell cultured samples after four days post-seeding and reported as the O.D. value. As shown in Fig.23, a significant difference ( <0.001) was observed in the absorbance of hybrid samples (HI) compared to the pure PMMA and bioglass. Tissue culture polystyrene (TCPS) was used as the control sample. The results of the MTS assay as the quantitative analysis highlights the superior cell interaction properties of the hybrid material; the absorbance at 490 nm increased from 0.3 ± 0.04 for pure PMMA to 0.5 ± 0.01 for hybrid specimens. These data demonstrate that covalent bonding between PMMA and bioglass had a positive impact on cell interaction and provide a superior environment for cell adhesion and proliferation. The presence of covalent bonds between the polymer and bioglass in hybrids leads to nanoscale interaction and emphasis of certain advantageous physico-chemical properties compared with those of the physical mixture.

4.8 Degradation

The degradation profile of the prepared hybrid was compared to that of the bioglass control samples. It was observed that presence of PMMA covalently bonded to the bioglass decreased the rate of degradation of the bioglass. As shown in Fig.24, linear lines were fitted to the degradation curves of the samples and it was observed that the degradation rate of the hybrid sample was 1.5 times less than the bioglass.

5. Conclusions

The structure of PMMA-co-MPMA-bioglass hybrids was studied at the molecular level using various techniques such as ¹ C, ²⁹ Si and 1H MR (ID and 2D), FTIR, TGA and DSC analyses. The results of these analyses corroborated the covalent bonding between the organic and inorganic moieties as a result of MPMA. A more open network was obtained by increasing the molar ratio of MPMA into the structure. The compositions of MPMA and the ratio of polymer to bioglass solutions were modified to decrease the gelation time from days to only two hours at room temperature in the absence of any initiator. A transparent and firm composite structure was produced with uniform

distribution of the compounds at optimum composition when using a high amount of MPMA and less than 40% bioglass. In addition, the thermal stability of hybrid at optimum composition was dramatically higher than the value reported in previous studies, underlining the superior physical characteristics of material produced in the present study. This material is expected to have broad applications, such as gate dielectric, optics, and coating - as well as biomedical applications including dental and bone replacement. 6.0 Additional studies

In this further study, the physical properties of hybrid (HI), were compared with physical mixture of PMMA-bioglass (HO). Both pure bioglass and PMMA homopolymer were considered as control samples, as shown in Figure 14. It was observed that only HI formed a one-phase, transparent three dimensional (3D) monolith structure, which is due to the formation of covalent bonding between MPMA and silanol groups of bioglass. However, in the absence of MPMA, the sol-gel method was not efficient for the creation of uniform mixture of PMMA and bioglass and two separate phases of a thin film of polymer and a gel structure of brittle bioglass were formed. Previous investigators have shown that it is viable to create 3D structure of PMMA- bioglass at different composition by in situ polymerization of MMA monomers in the presence of bioglass as filler. However, these products are still not as uniform as hybrids and exhibit a non-homogenous degradation profile and non- uniform distribution of mechanical force due to the different load-bearing characteristics of polymer and bioglass particles. The hybrid and bioglass were formed within 5 and 120 hours, by sol gel method at room temperature, respectively. However, it was found that by increasing temperature from 25 °C to 50 °C, the gelation times of HI and bioglass (BG) were significantly decreased to 3 hours and 45 minutes, respectively. These data demonstrate the significant impact of temperature on formation of silica network. The faster gelation of HI compared to bioglass was due to the presence of strong Si-C bonds. The rapid gelation is favorable for the fabrication of in situ products in clinical applications.

6.1 Morphological analyses

The morphology of HO and HI was compared in the macro- and nano- scale to assess the risk of phase separation as a result of lack of molecular integration. As shown previously, differences in surface characteristic of bioglass and HO elucidate the appearance of phase separation between ceramic and polymer for physical mixture of PMMA and bioglass (HO). However, the surface of hybrid samples was homogenous and no phase separation occurred in the structure of HI underlining the presence of chemical conjugation resulted in structural integrity. The absence of phase separation in HI and homogenous distribution of bioglass in PMMA structure was further confirmed by mapping the SEM image by EDS detector. The results demonstrated that calcium, silicon and carbon were distributed uniformly throughout the HI sample, while for HO, calcium was detected only on a few regions of surface. The absence of phase separation in nano-scale for HI and formation of two phases for HO were elucidated by STEM analysis.

AFM analysis was used to compare the surface properties of hybrids fabricated in this study with physical mixture of PMMA-bioglass. The roughness values were non-uniform for HO sample underlining the non- homogenous distribution of applied stress. However, as depicted in Figure 25 the surface of HI hybrid was smoother and the roughness dropped 5 fold compared to physical mixture.

The results of analysis by SEM, EDS, STEM, and AFM demonstrated that the addition of 0.1 molar ratio of MPMA:MMA that led to covalent bonding between PMMA and bioglass was efficient in creating a homogeneous mixture of PMMA and bioglass. This amount of coupling agent was adequate to prevent the phase separation between these two organic and inorganic phases. 6.2 Mechanical properties

Mechanical properties are one of the key factors for selecting

biomaterial. The bulk mechanical properties of prepared hybrid monoliths were compared to neat BG. Physical mixture (HO) was excluded from this test due to the lack of 3D structure and non-uniform gel formation. As shown in Figure 26(a), the uniaxial stress-train compression curves of HI and BG showed two regions (I and II) before fracture; the linear region I from which the Young's modulus was calculated and the quasi-linear region in which the stress was increased until fracture. It was observed that region I in HI extended up to 0.1 mm/mm while this region was decreased 10 fold in BG and did not exceed 0.01 mm/mm. Furthermore, The 40 fold reduction of Young's Modulus of bioglass (229 MP a) compared to HI (6 MP a) resulted in achieving higher mechanical stability for hybrid samples. The ultimate strain of HI, also, was eight fold higher than bioglass sample (Figure 26 (b)) that was completely collapsed in the first stage of mechanical test. Moreover, the toughness of hybrid was 2.6 fold higher than bioglass sample. Without wishing to be bound by theory, it is believe that the enhancement of mechanical properties of HI was attributed to the covalent bonding (Si-O-C bonds) and integration of PMMA chains into the silica structure.

It is favorable to develop a biomaterial that mimic the hardness of host tissue. As shown in Figure 26 (c), the hardness value of HI hybrid was close to natural bone {i.e. 33-45 Vickers hardness (HV)) and significantly higher than bioglass and PMMA and even their physical mixtures (p < 0.001). This data demonstrated that HI was more stable and less deformed under stress compared to HO, bioglass and PMMA. It can be concluded that the integration of PMMA chains into the molecular structure of silica increased the toughness and ductility, while decreased the brittleness. The compression strength and

Young's modulus of the HI that was prepared from 60:40 vol % PMFS:bioglass solutions were within the range suitable for osteoblast cells adhesion and proliferation. In vitro cell study was then conducted to assess the cell adhesion to this material.

6.3 Degradation analysis

One of the challenging issues in the application of composite materials is the separate degradation profile of each component, especially when one of the components degrades before the regeneration of the tissue. Inhomogeneous degradation causes uneven stress distribution or load transmission on the composite materials that leads to adverse effects such as loosening of implant.

The degradation of HI, bioglass and pure PMMA samples were measured within 100 days. The data in Figure 27 shows that PMMA has no weight loss during this period; however, bioglass was rapidly degraded and completely dissolved after 83 days. The mechanism of degradation of bioglass is well established in the literature. During the degradation of bioglass, Si0 ₂ is gradually dissolved and results in producing silicic acid, which is excreted from the body. However, the degradation rate of Si0 ₂ matrices is a function of parameters such as composition of precursors, watenTEOS molar ratio, pH, and network connectivity of the silica matrix that is shown by Q ⁿ species. Q ⁿ shows the number of Si-O-Si bonds around each silicon. For example, high value of Q ³ and low value of Q ⁴ are indicative of less compact network, hence faster degradation rate. Previous studies also showed that the covalent bonding between a bioglass and a polymer had a negligible effect on mechanism of degradation that is based on releasing the silica ion.

The composition of bioglass in this study was

TEOS:water:HCl:CC=l :8:0.01 :0.2 with 41% Q ³ and 53% Q ⁴ species that were calculated from ²⁹ Si MR analysis. The data in Figure 27 demonstrated that these characteristics led to the degradation rate of 0.92 % per day and 0.65 % per day for bioglass and HI samples, respectively. The zero-order kinetics of bioglass degradation was in agreement with other studies. The addition of PMMA to bioglass by covalent bonding (HI) resulted in impeding the degradation rate of bioglass for 1.5 fold and maintaining nearly 30 % of its original weight after 100 days of incubation. It was demonstrated that the mechanism of the new bone formation at the implant interface and the host tissue continues until there is ion exchange between the implant and body fluids. Therefore, retarding the degradation of silica matrix in HI provides a longer period for ion exchange, hence resulting in stronger interaction with the host tissue and formation of thicker apatite for bone formation. It is assumed that during gradual degradation of bioactive glass, innate hydroxyapatite and extracellular matrix are regenerated in situ that mimic the required mechanical strength for bone regeneration. In addition, no weight loss was observed in the degradation profile of HO which was due to the fact that the hydrophobic PMMA did not facilitate the interaction of bioglass particles with the surrounding solution.

The FTIR analysis conducted on the residues of HI sample after 100 days incubation in PBS at 37 °C confirmed the presence of silica ions in the degraded HI sample. As shown in Figure 28, C=0 and C0 ₃ ^2" peaks of polymer and bioglass detected at 1722 cm ^"1 and 1625 cm ^"1 , respectively. The presence of the 1625 cm ^"1 peak in the solid residues underlines that after 100 days, bioglass still maintained in the solid phase and was not completely dissolved in media. The ratio of intensities for characteristic peaks at 1625/1722 cm ^"1 was indicative of bioglass degradation. This ratio is calculated to change from 2 to 0.66 after degradation, which was in agreement with 30% weight loss of HI . This slower degradation rate is suitable for bone implant due to the fact that at least 90 days implantation time is required for bone tissues to regenerate. The visual observation and the presence of strong peak at 1722 cm ^"1 confirmed the presence of PMMA in the degraded sample.

6.4 Biological activity of PMMA-bioglass hybrids

Prior to in vitro test, the bioactivity of hybrid was tested by the incubation of samples in SBF that is a standard method for selection of biomaterials for bone and dental applications. In this method, formation of an apatite layer is indicative of bioactivity of samples. As shown in Figure 29, an apatite layer was clearly observed on the surface of both bioglass and hybrid samples. On the other hand, No apatite layer was formed on PMMA surface due to the inertness of this polymer. In physical mixture (HO), formation of apatite was observed only in some regions of the surface and the full coverage similar to the hybrid was not detected. This result confirmed the formation of apatite only on the surface of bioglass. One of the main advantages of the hybrids of the invention compared to physical mixtures or conventional composites is that the inorganic layer is well-distributed over the surface and it is more accessible for precipitation and formation of calcium-phosphate layer. However, in physical mixtures the organic phase covers the surface of samples and decreases the interface between inorganic phase and calcium-phosphate ions.

Furthermore, the formation of apatite on the surface of HI implies that the presence of inert PMMA polymer in the composition of 60:40 PMFS:bioglass had negligible impact on bioactivity.

6.5 In vitro attachment, proliferation, and osteoblastic differentiation

Composite materials representing bulk mixing (HO) and MPMA-coupled hybrid (HI) of PMMA:bioglass were analyzed in vitro for their cell attachment and osteoblast growth properties. These materials were compared with their constituent materials, pure PMMA and bioglass. The bioglass was markedly brittle and prone to thermal and mechanical shock induced fracturing.

Cells were found to adhere to all surfaces and supported attachment after 4 days in culture as visualized by scanning electron microscopy. Cells showed a more flattened morphology indicative of increased attachment on bioglass compared to PMMA. Notably this morphology was better maintained on the hybrid HI compared to HO, suggesting the improved homogeneity of the coupled PMMA-bioglass gave more consistent attachment.

Cell growth was measured by viability assay on cultured primary osteoblasts at day 1 and 7 after cell seeding. No significant difference was seen between the samples after 24 hours, however the HI hybrid showed a remarkable increase after 7 days which was superior to all other samples including the physical mixture in terms of cell viability. Cell viability number of HI was close to the control (TCPS) after 7 days of culture (p > 0.05). A trend towards increased viability was seen in tissue culture plastic controls, although this is a surface optimized for cell attachment and growth.

Finally, the alkaline phosphatase (ALP) activity of samples was assessed. Human pre-osteoblasts were grown upon HO, HI, PMMA and bioglass materials and treated with osteogenic differentiation media to induce expression of mature osteoblast markers. Robust staining was seen on cells grown on tissue culture plastic as a positive control. Notably, blue ALP+ cells were seen on the HI -treated hybrid samples but not on the HO samples or on the PMMA or bioglass alone. These data suggest that the MPMA-coupled hybrid material may have superior properties as a bone replacement material for early osteoblast differentiation.

6.6 In vivo study

6.6.1 Histological results— bioglass group

The brittleness of bioglass resulted in the formation of bulk-like structures within the implantation bed and fragments of various sizes ranging up to small particles after 10 days. As shown in Figure 33 and Figure 34, the implantation of this material provoked two different cellular reactions. Around the bulk-like structures within the implantation bed granulocytes and

mononuclear cells, mainly macrophages and lymphocytes, were embedded within a well vascularized granulation tissue. The immunohistochemical detection of the macrophage-specific F4/80 antigen showed that these cells were mainly detectable as a monolayer at the material-tissue-interface and loosely distributed within the material-adherent granulation tissue. Furthermore, only a small percentage of cells within the peri-implant tissue was shown to express tartrate-resistant acid phosphatase (TRAP).

The tissue reaction to bioglass fragments was dominated by

multinucleated giant cells in addition to the above mentioned mononuclear cells. The immunochemical detection of the F4/80 antigen additionally revealed the dominance of macrophages among all material-adherent mononuclear cells. Furthermore, TRAP staining revealed that within the implantation bed of fragments, both TRAP-positive mononuclear and multinucleated cells dominated when compared to those without TRAP-expression (control samples). Without wishing to be bound by theory it is believe that the occurrence of the multinucleated giant cells within the bioglass might be mainly size-related and less associated with the potential material impurity. 6.6.2 Histological results— HI group

Better mechanical properties and less brittleness of HI samples in comparison to the bioglass resulted in almost no material fragments within the peri-implant tissue during implantation and explantation (Figure 33). Thus, only one cellular reaction pattern was detected. The histological analysis at day 10 after implantation revealed that the HI implant was embedded within a cell- and vessel-rich connective tissue (Figure 33). This tissue was localized as a relatively thick wall along the material-tissue-interface and contained macrophages, granulocytes and lymphocytes as well as fibroblasts. The tissue reaction was comparable to that towards the bioglass bulk-like structures. It should be noted that no cell and tissue penetration into the material core was observed due to the lack of porosity. Directly at the surface of the HI implant predominantly macrophages were located (Figure 33). Furthermore, less number of multinucleated cells was detectable within the tissue adjacent to the material (Figure 33(b)). TRAP detection revealed that only a low number of cells at the surface of HI implant showed expression of this enzyme (Figure 33 (c)).

In this study an in vivo pilot study was performed in order to assess the tissue reaction to bioglass and HI samples. The histological evaluation has revealed that during the implantation phase of 10 days bioglass underwent a fragmentation most likely due to physical forces during animal movement rather than tissue penetration into the material. Consequently, this material induced two different cellular reactions. The bulk-like structures induced a mainly mononuclear inflammatory response, while the fragments induced

multinucleated giant cells. The occurrence of the latter was independent of fragment size. Overall, the implantation bed of this group was relatively extensive, which again can be attributed to the presence of the multinucleated giant cells, which are known to produce vascular endothelial growth factor (VEGF) among other substances responsible for material degradation. HI has negligible fragmentation during the observation period, which can be attributed to its higher mechanical stability. This material induced mainly mononuclear cells and a relatively low amount of multinucleated giant cells. Interestingly, the latter did not show a high TRAP-expression and it was observed that the implantation bed of HI was well vascularized.

6. 7 Conclusions

PMMA-bioglass hybrids were produced by sol-gel method in the presence of chemical bonding that integrated the organic and inorganic components. This molecular level interaction addressed the issue of phase- separation that is commonly observed in preparation of physical mixtures of bioglass with a polymer. In addition, the mechanical properties of the hybrid acquired at optimum composition of PMMA:MPMA was significantly improved compared with bioglass; the Young's modulus of hybrid was decreased 40 fold and its hardness was 16 fold higher than pure bioglass. The chemical bonding of PMMA with bioglass resulted in prolonging the degradation of bioglass that may be favorable for bone regeneration and in situ drug release applications. Additionally, the results of in vitro and in vivo studies demonstrated that the molecular level interaction had no adverse effect on biocompatibility of bioglass and in fact significantly enhanced its integrity and reduced the level of inflammation. Therefore, the fabricated hybrid in this study is a viable alternative for bioglass and PMMA-bioglass physical mixtures as bone regenerative materials.

7.0 Additional experimental studies

7.1 Materials & Methods

7.1.1 Materials

Precursors required for synthesis of PMMA-co-MPMA copolymer including MPMA, α,α'-Azoisobutyronitrile (AIBN) and Ν,Ν ^' - dimethylformamide (DMF) were purchased from Sigma and used as received. Methyl methacrylate (MMA) purchased from Sigma was used after distillation under reduced pressure. Hydrochloric acid (HCl; Merck), tetraethyl orthosilicate (TEOS; Sigma), calcium chloride dihydrate (CaCl ₂ .2H ₂ 0 (CC); Ajax Finechem Pty Ltd), tetrahydrofuran (THF; Merck) and deionized water were used for fabrication of the inorganic solution and the hybrid. 7.1.2 Preparation ofbioglass, PMMA-co-MPMA and pure PMMA solutions

TEOS was mixed with deionized water and HC1 and stirred for 30 minutes followed by addition of calcium chloride dihydrate. A common calcium source for the preparation of sol-gel derived bioglasses is calcium nitrate tetrahydrate; however, in this study, calcium chloride was used to minimize the risk of toxicity resulted from nitrate by-product. The precursors were mixed with the molar ratio of TEOS:water:HCl:CC=l :8:0.01 :0.2 and the solution was referred to as sol(A) for convenience.

Free radical polymerization technique was used for the synthesis of PMMA-co-MPMA with MPMA:MMA = 0.1 mol ratio coded as PMFS and

PMMA (without MPMA) using AIBN as an initiator. Precursors were mixed in a Schlenk flask (MM A: AIBN=200 (mol ratio), DMF (20 mL)) and degassed by three freeze-pump-thaw cycles. Polymerization was conducted at 70°C for a period of 12 hours. The polymer was purified by precipitation in diethyl ether followed by filtration and drying in vacuum. PMFS or PMMA were dissolved in THF with a concentration of 10 wt% and the solutions were labeled as sol(B) and sol(C), respectively.

7.1.3 Hybrid formation

Sol(A) and sol(B) were mixed in the volumetric ratio of

sol(B):sol(A)=60:40, then mechanically stirred for one hour to obtain a homogenous and well dispersed solution. This composition was selected due to the fact that bioglass composition was shown to have no significant impact on the network characteristics and molecular integration of hybrids.

It was then kept sealed until gel was formed, then was dried at ambient temperature and subsequently at 37°C for a period of 7 days at each

temperature. The product was then dried in vacuum oven at 40 °C for a period of 2 days. This temperature profile was developed for drying the samples to remove residues of solvents and maintain the monolith structure.

7.1.4 Physical mixture formation

Sol(A) and sol(C) were mixed (sol(C):sol(A)=60:40 vol%) and mechanically stirred for one hour; The same procedure was followed to dry these samples as hybrids. A thin film of polymer was obtained with a gel structure of bioglass that was subsequently grinded and well-mixed as powder. Afterwhich, the powder was dissolved in THF (10 wt%) and casted on a teflon container, which was then vacuum dried at 40 °C for 2 days.

7.2 Characterizations

7.2.1 Mechanical testing

Hybrid samples were prepared in the form of monoliths and uniaxial compression tests was performed in an unconfmed state with a 1000 N load cell by Instron (Model 5543). Dimensions of the samples were 6.61 ± 0.05 mm diameter and 1.2 ± 0.1 mm height. The samples were subjected to a loading and the Young's modulus was obtained as the tangent slope of the stress-strain curve between 0 % - 10 % strain level. The area under the compressive stress- strain curves was calculated for measuring the toughness of samples. Three samples were examined for each group for statistical analysis.

7.2.2 Degradation assessment

The degradation rate of samples was tested by measuring the change in sample weight over time under simulated physiological conditions. Three samples were kept in PBS at 37°C and at different time intervals they were removed from the degradation medium and rinsed three times with deionized water and dried prior to weighing. The measurement was continued for a period of 100 days.

7.2.3 Bioactivity

The samples were characterized for their bioactivity by incubation in simulated body fluid (SBF) at 37 °C for the periods of 1, 3, 7, and 14 days. Samples were washed twice with deionized water to remove any residue of minerals absorbed on the surface. The dried monoliths were examined to determine the formation of apatite layer on the surface.

7.2.4 Cell attachment, morphology and proliferation

Human pre-osteoblasts (HOB) were used to assess the cell interaction with the samples. Complete osteoblast growth media (Invitrogen) was used to culture the HOB cells, which were incubated at 37 °C in the presence of 5% C0 ₂ and 95 % humidity. The media was refreshed every three days until the cells approached confluence. The samples (12 mm diameter ^χ 1 mm height) were placed into a well plate and kept in 70 % ethanol for one hour for sterilization, followed by rinsing with PBS three times. Samples were then exposed to UV light for 30 minutes and were washed with fresh medium at 37 °C overnight. The substrates were placed into well-plates and seeded with cells at the density of 2 ^χ 10 ⁵ cells/ml and kept in culture for 7 days. Cell morphology was examined at day four of culture by SEM (FE-SEM; Zeiss ULTRA plus) analysis after being fixed with glutaraldehyde according to previously published method. ⁶⁴ Viability was measured at 1 day and 7 days post seeding. Cellular viability was assessed using the CellTitre 96 Aqueous One Solution Cell Proliferation Assay kit

(Promega) according to the manufacturer's instructions. Briefly, scaffolds were incubated with the viability solution for 30 min at 37 °C and read using a spectrophotometer at 595 nm. All samples were assayed in triplicate for statistical analysis. Tissue culture polystyrene (TCPS) was used as a control. 7.2.5 In vitro osteogenic differentiation and alkaline phosphatase activity

To initiate osteogenic differentiation, samples were transferred to media supplemented with ascorbic acid (50 μg/ml), β-glycerophosphate (10 mM) and BMP-2 (200 ng/ml) after 48 hours cultured with cells. The time of transferring the samples to this media was considered as day 0 for alkaline phosphatase (ALP) test.

ALP is an osteogenic marker that is expressed by differentiating osteoblasts. ALP staining was carried out on pure PMMA and bioglass, their physical mixture and hybrid at day 4. Samples were fixed with gluteraldehyde, washed with PBS and incubated in TRIS buffer (1 M, 9.4 pH) for 5 minutes. The scaffolds were then stained in Naphtol AS-BI phosphate (Sigma) as a substrate and Fast Blue (Sigma) as the stain. Scaffolds were washed with H ₂ 0 to remove excess staining before imaging. Cells alone on tissue culture plastic were used as a control. Images were captured using a Leica MZ6 microscope with a Qlmaging Micropublisher 5.0 camera. 7.3 In vivo study

7.3.1 Experimental design of the in vivo pilot study

The in vivo experiments were carried out after approval from the Committee on the Use of Live Animals in Teaching and Research of the State of Rhineland-Palatinate, Germany. A total of 10 female, 4-6 week-old CD1- mice (Charles River Laboratories, Germany) were reared under standard experimental conditions at the in vivo Laboratory Animal Unit at the Institute of Pathology of the Johannes Gutenberg University of Mainz, Germany. The animals were randomly divided into three experimental groups. Animals of the first two groups (n = 4 and 8 animals in total) underwent a subcutaneous implantation with the bioglass and prepared hybrid in this study in accordance with a previously described experimental setup. Two additional animals (n = 2) underwent the preparation of the subcutaneous pocket without biomaterial insertion (control group). This control group served for classification of the inflammatory response related to the operation procedures.

7.3.2 Subcutaneous implantation model

The subcutaneous implantation of bone substitutes was applied according to a previously published operation procedure. The animals were anesthetized with an intraperitoneal injection (10 ml of ketamine (50 mg ml ^"1 ) with 1.6 ml of 2% xylazine). Subsequently, a subcutaneous pocket in the subscapular region was formed by means of a scalpel and surgical scissor. The bone substitute materials were subsequently inserted under sterile conditions into the subcutaneous tissue pocket under the thin skin muscle of the subscapular region. Wound closure was performed by means of Prolene 6.0 suture material (Ethicon, Germany).

7.3.3 Explantation and tissue preparation

The tissue preparation for all of the groups was performed according to a previously described method. Experimental animals were sacrificed by an overdose of Ketamine and Xylazin at day 10 after implantation. After 10 days the bone substitute materials were resected together with the surrounding peri- implant tissue. Tissue fixation was carried out by means of 4% formalin for 24 hours. For further histological workup and (immuno-) histochemical staining, the tissue of the implant site was cut into three segments of identical dimensions containing the left margin, the center and the right margin of the biomaterial. Paraffin embedding was performed after dehydration of the biopsies in a series of increasing alcohol concentrations followed by xylol incubation. Six ongoing 3-5-μπι thick sections were made from the central segment of each animal by means of a rotation microtome.

7.3.4 Histological examination of the animal tissue

The material-tissue interaction was visualized by means of previously published histochemical and immunohistochemical staining methods. The first three slides sections were stained with haematoxylin and eosin (H&E), Movat ^' s Pentachrome and Azan, respectively. The fourth slide was used to identify osteoclast-like cells by tartrate-resistant acid phosphatase (TRAP) staining according to previously described methods. The fifth slide was used for immunochemical staining with a pan-macrophage marker (F4/80 antibody, rat monoklonal, Clon BM8, eBioScience, USA) in combination with peroxidase and diaminobenzidine (En Vision Detection System, Peroxidase/DAB, rabbit/mouse, K5007; Dako Cytomation, Hamburg, Germany). The sixth slide served as a control of the staining method in absence of the F4/80 antibody. All of the other chemicals were purchased from Sigma- Aldrich and used without further purification.

7.3.5 Histological analysis of animal tissues

The histopathological evaluation was performed by two independent investigators (SG and MB) by means of a conventional diagnostic microscope (Nikon Eclipse 80i, Tokyo, Japan). The description and the outcome of the cell- and tissue-biomaterial interactions were evaluated by examination of the total implantation bed and its peri-implant tissue as previously described.

8.0 Further studies

Poly(methyl methacryate) (PMMA)-bioglass composites were fabricated through the nano-scale interaction between the organic and inorganic components using silane coupling agent at optimized condition of

MPMA:MMA=0.1 and polymer:bioglass=60:40 vol%. The gelation time of hybrid solution was expedited from 5 hours to only 3 minutes by manipulating the process parameters such as solvent and by controlling the pH of hybrid solution. More condensed structures were produced with 23.4% decrease and 23.72% increase in the number of Q ³ and Q ⁴ species by using ethanol as the solvent and by the addition of sodium bicarbonate (SB) at optimum

concentration (0.4 wt%). Furthermore, the mechanical strength was enhanced 82 % and degradation rate was decreased 6-fold compared to the fabricated PMMA-bioglass hybrids discussed above resulting in the formation of more stable constructs. In addition, porous scaffolds were produced with 67% porosity and interconnective pores in the range of 100-300 μιη by using nontoxic SB at optimum concentration. The alizarin red S staining showed the differentiation of osteoblast cells on fabricated PMMA-bioglass hybrid underpinning its favorable characteristics for bone replacement applications.

The procedure of preparing the hybrids disussed above included dissolving MPMA-functionalized polymer in tetrahydrofuran (THF), which was subsequently mixed with bioglass sol. Fabricated transparent hybrid solution was gelled within 5 hours at room temperature and then dried for a period of 2 weeks for hardening and forming a monolithic crack-free structure. Removing the impurities such as TFIF and acid was also conducted during drying stage.

The presence of pores in 3D structure facilitates cell infiltration, oxygen, nutrients and water transfer, which improves the functionality of scaffolds. This property facilitates the integration of bone tissue into the semi-degradable porous structures. Sol-gel method was used for the fabrication of porous scaffolds from hybrid class of material. The pore formation in sol-gel derived hybrid solution was usually conducted by adding a surfactant such as Triton X- 100 to produce bubbles when agitating followed by the addition of hydrofluoric acid (FIF) as the catalyst for condensation of silica and fixative for bubbles. Fast gelation time of hybrid solution (< 10 minutes) is an important parameter, which determines the feasibility of this technique by fixation of bubbles inside the structure. This method was used to produce porous scaffolds from gelatin and poly(y-glutamic acid) with the modal pore size of 200-300 μπι. One drawback is the use of toxic chemicals such as hydrofluoric acid (FIF). In this study it was attempted to fabricate 3D porous structure from PMMA-bioglass hybrids with minimal processing time. The effect of several parameters such as temperature, solvent and reagent on the gelation time was investigated. In addition, it was examined to use sodium bicarbonate (SB) as an alternative to HF as a gas foaming agent for pore formation in these hybrids during the sol-gel method to circumvent the use of corrosive chemicals in the process. It was hypothesized that conducting the sol-gel method at 70°C resulted in the formation of C0 ₂ gas bubbles from the decomposition of sodium bicarbonate, hence the creation of porous composite scaffolds. Several analytical techniques were used to assess the impact of changing solvent and rapid gelation on the network connectivity, mechanical, degradation properties of hybrid. Furthermore, the biological activity of PMMA-bioglass hybrid as a bone replacement material was investigated through in vitro studies.

8.1 Materials and Methods

8.6 In vitro study

8.6.1 Cell culture

Mouse clonal pre-osteoblast cells (MC3T3) were used to assess the cell interaction with samples. Alpha Minimum Essential Medium (a-MEM) supplemented with 10% fetal bovine serum (FBS), 1% L-glutamine and 1% penicillin/streptomycin was used to culture the cells, which were incubated at 37°C in the presence of 5% C0 ₂ and 95 % humidity. The media was refreshed every three days until the cells approached confluence. MC3T3 cells were detached from flask by trypsin treatment and collected for cell seeding.

Two sets of samples were considered for cell study. The monolith samples (6 mm diameter x 3 mm height) and porous scaffolds (12 mm diameter x 2 mm height) were placed into well plates and kept in 70 % ethanol under UV light for one hour for sterilization followed by rinsing with PBS triple. The samples were placed into well tissue-culturing polystyrene plates (TCPS) and cells with the density of 1 ^χ 10 ⁵ cells/ml were seeded onto the monoliths and porous scaffolds, respectively. TCPS was used as control in this study. The samples were subsequently kept in a C0 ₂ incubator at 37 °C for further analyses. To initiate osteogenic differentiation, samples were transferred to media supplemented with ascorbic acid (50 μ§/ιη1), β-glycerophosphate (10 mM) and BMP-2 (200 ng/ml) after 7 days cultured with cells. The time of transferring the samples to this media was considered as day 0 for viability with osteogenic media.

8.6.2 Viability test

Viability was measured at day 7 after seeding cells and the addition of osteogenic media. Cellular viability was assessed using the CellTitre 96

Aqueous One Solution Cell Proliferation Assay kit (Promega) according to the manufacturer's instructions. Briefly, scaffolds were incubated with the viability solution for 30 min at 37 °C and read using a spectrophotometer at 595 nm. Cell growth and proliferation was measured by viability assay on cultured samples at day 1 and day 7 after cell seeding. Furthermore, the viability of cells after the addition of osteogenic media was also assessed after 7 days. All samples were assayed in triplicate for statistical analysis.

8.6.3 Mineralization test

The formation of calcium phosphate was determined by using Alizarin Red S staining after 28 days of cell culture (21 days after the addition of OM). Alizarin Red S-calcium complex is formed in this assay in a chelation process, and the end product is birefringent. Scaffolds were fixed in 4%

Paraformaldehyde (PFA) after 28 days of cell culturing and were kept in the fridge in PBS until performing test. Alizarin red solution was prepared by mixing Alizarin red S (2g) with 100 mL of deionized water and the pH of solution was adjusted to 4.1-4.3 with 10% ammonium hydroxide. After which, scaffolds were washed with deionized water and 500μ1 Alizarin Red Solution was added to each well and left for 5 minutes. Subsequently, scaffolds were washed in deionized water until all excess of red staining is removed and red nodules were observed by optical microscopy. For quantification of this assay, 500 μΐ of diluted Cetylpyridinium Chloride (10% solution) was subsequently added to each scaffold and the absorbance of this solution was then read at 540 nm. 8. 7 Results and Discussion

8. 7.1 Effect of solvent and temperature in sol-gel method on the gelation time of PMMA-bioglass hybrid

PMMA is soluble in solvents such as THF, acetic acid, l-Methyl-2- pyrrolidone (NMP) and ethanol at specific conditions. In this study, ethanol was used as an alternative to THF for the fabrication of PMMA-bioglass hybrids to assess the effect of solvent due to its lower boiling temperature {i.e. 78 °C) compared with NMP (with boiling temperature of 202 °C) and acetic acid. Furthermore, ethanol has been shown to be safe for biomedical formulations. For example, ethanol was used to dissolve sucrose acetate isobutyrate (SAIB) for injectable drug formulation for sustained delivery of Risperidone to

Schizophrenia patients with the commercial name of SABER™. Furthermore, SAIB dissolved in ethanol has also been used for the development of an injectable carrier for orthopaedic injuries and to deliver Recombinant human bone morphogenetic proteins (rhBMP-2), Zoledronic acid and HA particles to enhance bone regeneration rate. In these systems, ethanol diffused into the surrounding tissue and SAIB solidified in situ. Therefore, it can be concluded that ethanol is a suitable candidate for dissolving PMMA-co-MPMA

copolymer.

PMMA-co-MPMA copolymer was soluble in ethanol at above 70 °C. However, as shown in Figure 36, the addition of bioglass solution inhibited the precipitation of this copolymer in ethanol at below 70 °C. This effect was due to the rapid covalent bonding between this copolymer and bioglass sol. This temperature was, therefore, used for the formation of chemical bonding between silanol groups of MPMA and bioglass. The hybrid solution was then rapidly cooled down to a predetermined temperature for further processing and characterization. The optimized PMMA-bioglass hybrid with 0.1 molar ratio MPMA:MMA and 60:40 vol% polymenbioglass solutions was used in this study. The PMMA-bioglass hybrid fabricated by using ethanol as the solvent was coded as Hl _E toH and its characteristics was compared with HI samples, which possessed the same characteristics (MPMA:PMMA= 0.1 molar ratio, polymer:bioglass=60:40 vol%) and was prepared in THF. The effect of solvent on gelation time was compared when the hybrids (HI and Hl _E toH) were prepared at room temperature (25 °C). The gelation time of samples was measured by test tube tilting method. As shown in Figure 37, the gelation time of Hl _E tOH hybrid was 120 ± 7 minutes that was significantly lower than HI sample that required 300 ± 12 minutes (at 25 °C) to gel. The reduction of gelation time of Hl _E toH hybrid was attributed to the presence of ethanol, which increased the rate of condensation. This effect of ethanol on the mechanism of silica network formation is well established in literature. Briefly, ethanol decreases the rate of hydrolysis and increases the rate of condensation in sol-gel method and results in the fabrication of bulkier and more branched structures.

The temperature had a significant impact on the gelation time of hybrid. As shown in Figure 37, by increasing the temperature from 25 °C to 37 °C, the gelation time of HI and Hl _E tOH was reduced from 300 ± 12 minutes to 90 ± 7 minutes and from 120 ± 7 minutes to 37 ± 2 minutes, respectively. In addition, by increasing the temperature from 25 °C to 70 °C, the gelation time of Hl _Et o _H hybrid was reduced from 120 ± 7 minutes to 7 ± 0.03 minutes. It is important to note that it was not practical to fabricate hybrid in THF (HI) at above 55 °C due to the rapid solvent evaporation and precipitation of copolymer in sol-gel system.

8. 7.2 Effect of sodium bicarbonate on the gelation time of PMMA-bioglass hybrid

Sodium bicarbonate (SB) was used for the preparation of buffer solution and adjusting the pH of hybrid samples, hence its gelation time. SB is a basic solution and provides OH ^" ions that can promote the rate of condensation in sol- gel method. Different amount of SB (from 0 to 10 wt% SB to hybrid) was added to PMMA-bioglass hybrid solution at room temperature. Before the addition of SB, the pH of hybrid solution was between 2 and 3. However, the pH of hybrid was raised to 7 before gelation for all concentrations of SB that was examined. As shown in Figure 38, by adding 0.4 wt% SB, the gelation time of hybrid was expedited from 120 ± 7 minutes to only 3 ± 0.03 minutes. Further increasing the concentration of SB to 10 wt% resulted in decreasing the gelation time to 15 ± 0.03 seconds. Addition of SB for more than 10 wt% was not practical as the gelation time was too fast to be controlled. Lower concentrations of SB (e.g. 0.4 wt%) provided better controllability on the gelation time and resulted in the fabrication of homogenous monolithic structure from PMMA-bioglass hybrid at room temperature. These results demonstrated that SB can be used as an alternative to HF that is commonly used in sol-gel method for accelerating the condensation of silanol groups and reducing the gelation time. SB has no toxicity compared to HF that is counted as a corrosive chemical.

For further characterization, Hl _E tOH was produced at optimized condition with fastest gelation time. Briefly, the hybrid solution was prepared by mixing PMMA-co-MPMA copolymer dissolved in ethanol (10 wt%) at 70 °C and mixed with bioglass sol in the polymenbioglass = 60:40 vol% at 70 °C for 10 seconds to allow the reaction between the silanol groups of MPMA and bioglass occur. Subsequently, the solution was cooled down to room

temperature and 0.4 wt% SB was added to the solution. Gel which was formed in 3 ± 0.03 minutes was placed at 37 °C for hardening. This sample was compared with HI, which was produced at room temperature using THF in the absence of SB.

8. 7.3 Impact of fast processing on network characteristics

Characterization of the thermal behavior of Hl _E tOH and HI hybrids showed no significant change in the thermal degradation profile obtained via TGA analysis. It was observed that the onset thermal decomposition

temperature (OTDT) of both samples were at 400 °C.

The network factors (Q ⁿ ) were measured to determine the impact of gelation time on the structure of final product. Q ⁿ shows a silicon bonds to n other silicons via "bridging" oxygen atoms. The Q ⁿ species and network connectivity model (Nc) of samples (HI and Hl _E toH) were listed in Table 10. Nc, which ranges between zero to 4 describes the characteristics of silica network and is defined as the average number of bridging oxygen atoms per silicon ⁸ . It was shown that the acceleration of gelation time resulted in increasing the number of Q ⁴ species and Nc value, underpinning constructing a more condensed structure. The Nc of Hl _E tOH was enhanced 15 % compared with HI . This is in agreement with the studies that investigated the mechanism of sol-gel process of silanol groups. It was demonstrated that the rate of hydrolysis and condensation reactions plays an important role on the final shape and structure of network. When silica gel is synthesized in alcoholic solutions, number of bridging oxygens increase and condensation reaction continues for a long period. This phenomena leads to an increase in the value of Q ³ and Q ⁴ species in the silica network and formation of bulkier structure. Therefore, it is reasonable to postulate that the Hl _E toH hybrid possessed greater Nc Value compared with HI .

Table 10: The effect of gelation time on the characteristics of hybrid network

8. 7.4 Effect of gelation time on the mechanical properties of hybrid

The desired mechanical properties for bone replacement are to possess low Young's modulus while maintaining the high ultimate stress. It was demonstrated in our previous study that the integration of PMMA into the structure of bioglass resulted in 40-fold decrease in the Young's modulus and this value was decreased from 229 MPa in pure bioglass to 6 MPa in HI samples. This effect was efficient in addressing the issue of brittleness of bioglass, however, the ultimate stress of HI sample was still not sufficient for load-bearing bone replacement applications. In this study, the effect of fast network formation and more condensed structure on the mechanical properties of samples was investigated. As shown in Figure 39, reducing the gelation time in Hl _E toH had negligible effect on the Young's modulus of samples compared with HI . However, the ultimate stress of Hl _E tOH was 5.6-fold higher than HI and approached 5.7 ± 1.16 MPa. This value was in the range of the ultimate stress of cancellous bone (i.e. 5-10 MPa). However, it was lower than the ultimate stress of cortical bone (i.e. 100-230 MPa) and sintered bioglass at 800 °C (70S30C) (i.e. 32-89 MPa). The toughness value of Η1 _Ε , _0Η, however, approached 1174 Nm ^"2 , which was 1300-fold greater than the toughness of bioglass 70S30C (i.e. 0.9 Nm ^"2 ). Furthermore, it was observed that the toughness value of Hl _E toH was 1.6-fold greater than HI due to the fast gelation time and more condensed silica structure (higher Nc value). Higher toughness value of Hl _E toH is beneficial as the sample is more resistible to the applied load.

8. 7.5 Biological activity of PMMA-bioglass hybrids

Prior to conducting biological assays, the bioactivity of samples was assessed by the formation of apatite layer when they were soaked in SBF. The results in Figure Error! Reference source not found.40 show a uniform formation of apatite layer the surface of HI and Hl _E toH- However, this layer was more crystallized and thickened on the surface of Hl _E tOH compared with HI . This effect was due to the presence of OH ^" ions, which expedited the nucleation and crystallization of apatite crystals. Furthermore, the calcium phosphate characteristics of the formed crystals on the surface of hybrids were confirmed using EDS analysis (see Figure 40(c)).

8. 7.6 Injectability of PMMA-bioglass hybrid

The feasibility of using Hl _Et o _H hybrid for injectable formulation was investigated. Hybrid solution was injected through a 21G needle and the injectability profile at room temperature was shown in Figure 41. Mixing and loading the syringe were carried out in 2 minutes. Subsequently the syringe was exposed to the load (N) and the stress versus time was recorded. Solution was steadily extruded through the needle and no increase was detected in the load until 60 seconds. After which, a sudden increase was observed in the required load, due to gelation of hybrid solution. This solution can be contemplated as an alternative for injectable polymers mixed with hydroxyapatite or bioglass particles because this system reduced the risk associated with needle blockage due to the presence of bioglass particles. In addition, the hybrid also provides a more homogenous mixture of polymer and bioglass.

8. 7. 7 Degradation

The degradation property of bone implants is one of the most challenging issues in the application of composite materials. Inhomogeneous degradation of glassy/ceramic and polymer compounds results uneven stress distribution or load transmission on the composite materials that leads to adverse effects such as loosening of the implant. It was demonstrated that the formation of nano-scale interaction between organic and inorganic components addressed the issue of separate degradation and hybrids possessed superior degradation properties compared with physical mixtures and conventional composites.

The degradation mechanism of bioglass is well established in the literature. Briefly, The Si-O-Si network of bioglass gradually disrupts producing silicic acid, which is excreted from the body. This phenomena is a function of parameters such as composition of precursors, watenTEOS molar ratio, pH, and Nc. For example, low Nc value results in less compact network, hence faster degradation rate. Furthermore, it was demonstrated that the release of Si ions to body fluids is beneficial for the mechanism of new bone formation at the site of implant _.

The degradation of HI and Hl _E toH samples was measured within 100 days and was shown in Figure 42. The degradation rates of HI and Hl _E toH was 0.65% and 0.1% per day, respectively. The lower degradation rate of Hl _E toH is attributed to the higher Nc value of this sample which resulted in more condensed network. Higher Nc value resulted in impeding the degradation of Hl _E tOH for 6-fold while providing the required ion exchange to promote the interaction with the host tissue and formation of apatite around the implant. Therefore, it can be concluded that Hl _Et oH can be an appropriate candidate for long term bone implant application due to the high mechanical properties, long- term ion exchange with surrounding tissue and the ability to be shaped as required.

8. 7.8 Formation of porous scaffolds

The low concentrations of SB (e.g. 0.4 wt%) was used to form firm monolithic structures. However, higher concentrations (e.g. 10 wt%) was used for the fabrication of porous scaffolds. SB possessed dual effect: a) formation of bubbles due to the decomposition of SB at 70 °C into C0 ₂ , and b) rapid gelation (15 seconds) due to increase in the pH of hybrid solution, which fixed the bubbles. During gelation the pores were also ruptured leading to interconnectivity in the structure. As shown in Figure 43(a), the pores were irregular and their size ranged between 100-300 μηι, which is in the range suitable for osteoblast cells proliferation. Furthermore, the μ-CT image of this porous scaffold (Figure 43(b)) demonstrated 67.38% porosity in the scaffold, the majority of which (> 99%) were open pores. The presence of open pores confirms the interconnectivity, which is a crucial factor for water and nutrient transfer through 3D scaffold. The numbers obtained by μ-CT image showed small error bars that confirmed the repeatability and reliability of this technique for fabrication of porous scaffolds during sol-gel reaction. However, it must be noted that this procedure was not successful in producing pores in neat bioglass due to its low mechanical properties and high brittleness. The porous scaffolds (PS) and monoliths fabricated in this study were further considered for biological characterizations.

8. 7.9 Mechanical properties of porous hybrid scaffold

The mechanical properties of porous scaffolds fabricated from Hl _E toH formulation were also examined. Porous scaffolds (PS), similar to all other porous structures, possessed lower value of Young's modulus and compressive strength compared to monoliths. The Young's modulus of porous scaffolds was 0.087±0.006 MPa and their toughness value was 4.57±1.14 Nm ^"2 . The stress- strain diagram of these samples, however, was very steady and the samples were as resistible as the monoliths to the level of strain. As shown in Figure 44, the ultimate strain of porous samples was 0.26 mm/mm, which showed negligible difference with the value for monolithic structures {i.e. 0.32 mm/mm).

8. 7.10 In vitro Cell viability

The viability of cells in a specific period of time (at day 1 and 7) and through the addition of osteogenic media was recorded and shown in Figure 45. It was demonstrated that cells viability for monoliths and porous scaffolds (PS) was increased 4.2-fold and 4-fold from day 1 to day 7, respectively. This result demonstrates that the Hl _E tOH hybrid provided a favorable environment for cells to adhere to and proliferate. Furthermore, it was observed that the addition of osteogenic media did not have negative impact on the viability of cells and the cell numbers were not negatively impacted by the addition of osteogenicmedia. 8. 7.11 Cell mineralization assay

The ALP staining, which was the primary expression of osteoblastic functions was conducted on the PMMA-bioglass hybrids in our previous study and promising results were obtained for the mineralization ability of PMMA- bioglass hybrid. Therefore, in this study alizarin red S staining was conducted as the long term mineralization capability of this material. The results of alizarin red S staining (Figure 46) showed that morphologic differentiation to osteoblasts and formation of minerali zed nodules occurred in both Hl _E toH monoliths and porous scaffolds (PS). The mineralized regions (red spots) were shown in Figure 46 by arrows. Furthermore, the presence of red nodules even before the addition of osteogenic media confirmed the osteoconductivity of hybrid. In addition, the accumulation of red nodules in hybrids in several regions corroborated the formation of mineralized tissues on these samples. The quantification of Alizarin red S study shown in Figure 47, also, confirmed that the PMMA-bioglass hybrid fabricated in this study possessed significantly (p<0.05) higher absorbance compared with the TCPS before the addition of osteogenic media. After the addition of osteogenic media, a remarkable increase was observed in the absorbance of TCPS compared to hybrid samples. This was due to the fact that the staining was penetrated through the samples and lower amount was leached out while for TCPS the staining was easily detected.

8.8 Conclusions

New technique was developed for the fabrication of PMMA-bioglass hybrids with improved physico-chemical properties. It was demonstrated that ethanol was an effective alternative solvent for TUF, which not only improved the biocompatibility of processing procedure but also resulted in faster gelation time (3 minutes). Fabricated hybrid possessed more condensed structures with higher Nc value, higher mechanical strength and longer degradation profile. Produced hybrid was easily shaped due to the favorable injectability.

Furthermore, porous scaffolds with high pore interconnectivity were fabricated during sol-gel technique by using SB as the gas foaming and silica condensation catalyst. The differentiation of osteoblast cells on the hybrids was confirmed and opened an avenue for the application of fabricated hybrid for replacing damaged bone tissue.

In the description provided herein, numerous specific details are set forth. However, it is understood that embodiments of the invention may be practiced without these specific details. In other instances, well-known methods, structures and techniques have not been shown in detail in order not to obscure an understanding of this description. Thus, while there has been described what are believed to be the preferred embodiments of the invention, those skilled in the art will recognise that other and further modifications may be made without departing from the spirit of the invention, and it is intended to claim all such changes and modifications as fall within the scope of the invention. Any formulas given above are merely representative of procedures that may be used. Steps may be added or deleted to methods described within the scope of the present invention.