Glass-ceramics are fine-grained polycrystalline materials formed when glasses of suitable compositions are heat treated and thus undergo controlled crystallisation to the lower energy, crystalline state. It must be emphasised here that only specific glass compositions are suitable precursors for glass-ceramics due to the fact that some glasses are too stable and difficult to crystallise whereas others result in undesirable microstructures by crystallising too readily in an uncontrollable manner. In addition, it must also be accentuated that in order for a suitable product to be attained, the heat-treatment is critical for the process and a range of generic heat treatment procedures are used which are meticulously developed and modified for a specific glass composition.
A glass-ceramic is formed by the heat treatment of glass which results in crystallisation. Crystallisation of glasses is attributed to thermodynamic drives for reducing the Gibbs’ free energy, and the Amorphous Phase Separation (APS) which favours the crystallisation process by forming a nucleated phase easier than it would in the original glass. When a glass is melted, the liquid formed from the melting might spontaneously separate into two very viscous liquids or phases. By cooling the melt to a temperature below the glass transformation region it will result in the glass being phase separated and this is called liquid-liquid immiscibility. This occurs when both the phases are liquid. Hence a glass can simply be considered as a liquid which undergoes a demixing process when it cools. The immiscibility is either stable or metastable depending on whether the phase seperation occurs above or below the liquidus temperature respectively. The metastable immiscibility is much more inmportant and has two processes which then cause phase seperation and hence crystallisation; nucleation and crystal growth and spinodal decomposition.
The first APS process has two distinguished stages; Nucleation (whereby the crystals will grow to a detectable size on the nucleus) and Crystal growth. Nucleation can either be homogeneous; where the crystals form spontaneously within the melt or heterogeneous; crystals form at a pre-existing surface such as that due to an impurity, crucible wall etc. Many a time the parent glass composition is specifically chosen to contain species which enhance internal nucleation which in the majority of cases is required. Such species also called nucleating agents can include metallic agents such as Ag, Pt and Pd or non-metallic agents such as TiO2, P2O5 and fluorides. The second process is spinodal decomposition which involves a gradual change in composition of the two phases until they reach the immiscibility boundary. As both the processes for APS are different, the glass formed will clearly result in having different morphology to each other.
A glass-ceramic is usually not fully crystalline; with the microstructure being 50-95 volume % crystalline with the remainder being residual glass. When the glass undergoes heat treatment, one or more crystalline phases may form. Both the compositions of the crystalline and residual glass are different to the parent glass. In order for glass-ceramics having desirable properties to be developed, it is crucial to control the crystallisation process so that an even distribution of crystals can be formed. This is done by controlling the nucleation and crystal growth rate. The nucleation rate and crystal growth rate is a function of temperature and are accurately measured experimentally (Stookey 1959; McMillan 1979, Holand & Beall 2002)
The aim of the crystallisation process is to convert the glass into glass-ceramic which have properties superior to the parent glass. The glass-ceramic formed depends on efficient internal nucleation from controlled crystallisation which allows the development of fine, randomly oriented grains without voids, microcracks, or other porosity. This results in the glass-ceramic being much stronger, harder and more chemically stable than the parent glass.
Glass-ceramics are characterised in terms of composition and microstructure as their properties depend on both of these. The ability of a glass to be formed as well as its degree of workability depends on the bulk composition which also determines the grouping of crystalline phases which consecutively govern the general physical and chemical characteristics, e.g. hardness, density, acid resistance, etc. As mentioned earlier, nucleating agents are used in order for internal nucleation to occur so that the glass-ceramic produced has desirable properties. Microstructure is the key to most mechanical and optical properties; it can promote or diminish the role of the key crystals in the glass-ceramic. The desirable properties obtained from glass-ceramics are crucial in order for them to have applications in the field of biomaterials.
Glass-ceramics are used as biomaterials in two different fields: First, they are used as highly durable materials in restorative dentistry and second, they are applied as bioactive materials for the replacement of hard tissue. Dental restorative materials are materials which restore the natural tooth structure (both in shape and function), exhibit durability in the oral environment, exhibit high strength and are wear resistance. In order for dental restorative materials to restore the natural tooth structure, it is crucial to maintain the vitality of the tooth. . However non-vital teeth may also be treated with restorative materials to reconstruct or preserve the aesthetic and functional properties of the tooth.
In order for glass-ceramics to be used for dental applications, they must possess high chemical durability, mechanical strength and toughness and should exhibit properties which mimic the natural tooth microstructure in order for it to be successful as an aesthetic. Glass-ceramics allow all these properties to be united within one material. As mentioned previously, for a glass-ceramic to have the desired properties, the glass is converted into a glass-ceramic via controlled crystallisation to achieve the crystal phase wanted and hence the desired properties it could possibly have. Hence, the glass-ceramic developed allows it to have properties such as low porosity, increased strength, durability, toughness etc which are crucial in the field of dental restorations as it prevents restorative failures which are mainly due to stress and porosity which causes cracks and hence failures.
It took many years of research in order to get a material strong enough to be initially used as a dental reconstructive material. However over the past 10-15 years, research has progressed vastly and now glass-ceramics demonstrate good strength, high durability and good aesthetics. The development and processing of glass-ceramics has been focused on particular clinical applications, such as dental inlays, crowns, veneers, bridges and dental posts with abutments.
Glass-ceramics are divided into seven types of materials:
Mica apatite glass-ceramics
Leucite apatite glass-ceramics
Lithium Disilicate glass-ceramics
Apatite containing glass-ceramics
The first commercially usable glass ceramic products for restorative dentistry were composites of mica glass ceramics. Dicor® and Dicor® MGC were products based on these. According to the mechanism of controlled volume crystallisation of glasses, tetrasilicic micas, Mg2.5Si4O10F2, showing crystal sizes of 1 to 2 μm in the glass ceramic were produced. Dicor® being amongst them was shaped by means of centrifugal casting methods to produce dental restorations such as dental crowns and inlays. Depending on the different crystal sizes and the corresponding microstructure of the glass ceramic, it was also possible to manufacture glass ceramics for machining applications. , Dicor® MGC being amongst them. This resulted in the characteristic of good machinability in this type of glass-ceramic to be exploited and results concluded that crystals upto only 2 μm in length in the material improved mechanical strength over other materials.
Mica-apatite glass-ceramics have been produced in the SiO2-Al2O3-Na2O-K2O-MgO-CaO-P2O5-F system. The main crystal phases are phlogopite, (K,Na)Mg3(AlSi3O10)F2
and fluorapatite, Ca5(PO4)3F. The base glass consists of three glass phases: a large droplet-shaped phosphate-rich phase, a small droplet-shaped silicate and a silicate glass matrix. Mica is formed during heat treatment, as in apatite-free glass-ceramics, by in-situ crystallization via the mechanism of volume crystallization. Apatite is formed within the phosphate-rich droplet phase. Astonishingly, every single apatite crystal possesses its own nucleation site in the form of a single phosphate drop. The glass-ceramic is biocompatible and suitable for applications in head and neck surgery as well as in the field of orthopaedics.
Leucite glass-ceramics can be formed by applying the advantage of the viscous flow mechanism. IPS Empress® is of this type of glass-ceramic. The material is processed by using the lost wax technique, whereby a wax pattern of the dental restoration such as an inlay, onlay, veneer or crown is produced and then put in a refractory die material. Then the wax is burnt out to create space to be filled by the glass-ceramic. As the glass-ceramic has a certain volume of glass phase, the principle of viscous flow can be applied and hence the material can be pressed into a mould. Surface crystallisation and surface nucleation mechanisms were controlled in order for this type of glass-ceramic to be formed. [42, 54] Consequently, the manufacturing of inlays and crowns developed due to the application of viscous flow mechanism of glass-ceramics in different shapes. The resulting leucite glass-ceramic restorations transluceny, colour and wear resistance behaviour can then be adjusted to those of natural tooth. Additionally, the leucite glass-ceramic restorations can be produced by machining with CAD/CAM. IPS ProCAD® and IPS Empress® CAD are glass ceramics produced via this method. All leucite glass-ceramic restorations are bonded to the tooth structure with a luting material, preferably an adhesive bonding system. The retentive pattern produced on the glass-ceramic surface is particularly advantageous in this respect.
It was possible to develop a leucite apatite glass-ceramic derived from the SiO2-Al2O3-Na2O-K2O-CaO-P2O5-F system by combining two different mechanisms, i.e. controlled surface nucleation and controlled bulk nucleation. IPS d.SIGN® is amongst these. The glass-ceramic was prepared according to the classic method of glass-ceramic formation: melting, casting to prepare a glass frit, controlled nucleation and crystallization. A two-fold reaction mechanism leads to the precipitation of fluoroapatite, Ca5(PO4)3F and leucite, KAlSi2O6 . SEM pictures show the two-phase crystal content of apatite and leucite in this type of glass-ceramic. Fluoroapatite phase takes the form of needle-shaped crystals whereas the oval areas are the leucite crystals. The clinical application of this glass-ceramic has been proven to be suitable for clinical application as veneering material on metal frameworks for single units as well as for large dental bridges involving more than three units.
The first glass-ceramic to be developed was by Stookey et al (1959) which contained Lithium disilicate. . Further research into this field allowed for IPS Empress®2 to be developed. This glass-ceramic was developed in order to extend the range of indications of glass-ceramics from inlay and crowns to three-unit bridges, by offering high strength, high fracture toughness and at the same time, a high degree of translucency. Both the flexural strength and fracture toughness of lithium disilicate glass-ceramics are almost three times of those of leucite glass-ceramics. Lithium disilicate glass-ceramic ingot are utilizied to produce the crown or bridge framework in combination with the viscous flow process. To further improve the aesthetic properties, i.e. translucency and shade match, and to optimally adjust the wear behaviour to that of the natural tooth, the lithium disilicate glass ceramic is veneered with an apatite-containing glass-ceramic using a sintering process.
In order to meet the demanding requirements of CAD/CAM applications, a lithium metasilicate glassceramic, IPS e.max®was developed. This material, which is supplied in a typically blue colour, is adjusted by thermal treatment in order to demonstrate a characteristic tooth colour.
The range of IPS e.max®products also encompasses various apatite-containing glass ceramics that are suitable for both layering material on lithium disilicate glass-ceramic and veneering material on ZrO2 sintered ceramic. The apatite crystal phase of the Ca5(PO4)3F type acts as a component that adjusts the optical properties of the restoration to natural tooth. For this reason, the crystallites are of nanoscale dimension.
ZrO2 containing glass-ceramics was the first glass-ceramic developed to be fused to high strength ZrO2 ceramic dental posts. The glass-ceramic contains Li12ZrSi6O15 crystals as the main phase; however different types of crystals are also precipitated in the glassy matrix. ZrO2 has become very interesting not only in the field of medicine but also in dental applications. High-strength and high toughness dental posts, crowns and bridges can be prepared from this material.
In order for a dental restorative material to be of clinical success, their most important properties include; high strength, high toughness, abrasion behaviour comparable to natural teeth, translucency, colour, durability) and the processing technologies (moulding, machining, sintering).  Furthermore, the material should have good marginal fit with the tooth, biocompatibility, good mechanical properties and low porosity. In addition to the aforementioned properties, the recent requirement for dental restorative materials is for its appearance to be similar to that of a natural tooth.
Glass-ceramics have been researched immensely in order to fulfil high standards of function and aesthetics from an early stage. The trend for metal free dental restorations began from the 1970’s whereby metal free feldspathic ceramics were reinforced with additional components. Since then, increasing the strength of these materials progressed rapidly by controlling the nucleation and crystallisation of glasses, as discussed earlier. These developments have now led to the introduction of a trend which is focused on achieving exceptional aesthetic results with glass ceramics as metal free dental restorations.
Although glass-ceramics exhibit the desired properties for dental restoration, their main drawback is that they are brittle which the main cause of failure is. This is due to either fabrication defects; which are created during production of the glass-ceramic or secondly, surface cracks; which are due to machining or grinding. Therefore when processing the glass-ceramic, care needs to be taken in addition to choosing the suitable method for production for specific compositions of the glass-ceramic in order to improve their mechanical properties.
Apart from the use of glass-ceramics for dental restorations, they can also be applied as bioactive materials for the replacement of hard tissue. Bone is a complex living tissue which has an elegant structure at a range of different hierarchical scales. It is basically a composite comprising collagen, calcium phosphate (being in the form of crystallised hydroxyapatite, HA or amorphous calcium phosphate, ACP) and water. Additionally, other organic materials, such as proteins, polysaccharides, and lipids are also present in small quantities. Because bone is susceptible to fracture; there has always been a need, since the earliest time, for the repair of damaged hard tissue.
Many years of research has attempted to use biomaterials to replace hard tissue, ranging from using bioinert materials, to bioactive materials such as ‘Bioglass’ (Hench et al) to ‘Apatite-wollastonite (A-W) glass-ceramics (Kokubo et al) and to calcium phosphate materials. Calcium phosphate based materials have received a great deal of attention in this field due to their similarity with the mineral phase of bone.
1.2 Calcium Phosphate Glasses
The application of calcium phosphate material as a bone substitute began by Albee (1920), who reported that a tricalcium phosphate compound used in a bony defect promoted osteogenesis. Many years later, Levitt et al (1969)  and Monroe et al (1971) were the first to suggest the use of calcium phosphate ceramics for dental and medical implant materials. Subsequently in 1971, Hench et al developed a calcium phosphate containing glass-ceramic, called Bioglass® and demonstrated that it chemically bonded with the host bone through a calcium phosphate rich layer. Furthermore the advantageous properties of calcium phosphate ceramics arose when Nery et al (1975) used a calcium phosphate ceramic for implants in surgically produced infrabony defects in dogs. This demonstrated that the calcium phosphate ceramic was nontoxic, biocompatible, and caused no significant haematological changes in the calcium and phosphorus levels. Since then, a great deal of research into calcium phosphate glass-ceramics has been conducted as potentially bone substitutes in dentistry.
Calcium phosphate based ceramics can be characterised accordingly;
Hydroxyapatite (HA, Ca5(PO4)3OH)
β-tricalcium phosphate (β-TCP, β-Ca3(PO4)2)
Biphasic calcium phosphates, BCP; mixture of HA and β-TCP
β-calcium pyrophosphate (β-CPP, β-Ca2P2O7)
Fluorapatite (FAP, Ca5(PO4)3F)
Calcium phosphate based ceramics and their properties have been characterised according to the proportion of calcium to phosphorus ions in the structure. One of the most widely used synthetic calcium phosphate ceramics is hydroxyapatite, Ca5(PO4)3OH, HA and this is due to its chemical similarities to the inorganic component of hard tissues. HA, has a Ca:P molar ratio of 1.67. It has higher stability in aqueous media than other calcium phosphate ceramics.
Tricalcium phosphate (TCP) is a biodegradable bioceramic with the chemical formula, Ca3(PO4)2. TCP dissolves in physiological media and can be replaced by bone during implantation. TCP has four polymorphs, the most common ones being α and β-forms, of which β-TCP has received a lot of attention in the field of bone substitutes. Slight imbalances in the ratio of Ca:P can lead to the appearance of extraneous phases. If the Ca:P ratio is lower than 1.67, then alpha- or beta tricalcium phosphate may be present after processing. If the Ca:P is higher than 1.67, calcium oxide (CaO) may be present along with the HA phase. These extraneous phases may adversely affect the biological response to the implant in-vivo. A TCP with a Ca:P ratio of 1.5 is more rapidly resorbed than HA. Hence, β-TCP has been involved in recent developments aimed to improving its biological efficiency and its mechanical properties in order for it to be successful as bone substitutes.
Mixtures of HA and TCP, known as biphasic calcium phosphate (BCP), have also been investigated as bone substitutes and the higher the TCP content in BCP, the higher the dissolution rate.
The crystal structure of HA can accommodate substitutions by various other ions for the Ca2+, PO43− and OH− groups. The ionic substitutions can affect the lattice parameters, crystal morphology, crystallinity, solubility and thermal stability of HA. Anionic substitutions can either occur in the phosphate- or hydroxyl positions. Fluorapatite and chlorapatite are common examples of anionically substituted HA. They display a similar structure to HA, but the F− and Cl− ions substitute for OH−. A lot of research has gone into carbonate substituted HA and it has shown to have increased bioactivity compared to pure HA, which is attributed to the greater solubility of the carbonated substituted HA. Thus, recent work has been in progress in order to optimise the production and sintering behaviour of carbonated substituted HA in order for use in biomedical applications.
Materials which are bioactive i.e. the ability to bond to living tissue and enhance bone formation, have the following characteristic compositional features: (i) SiO2 contents smaller than 60 mol%, (ii) high Na2O and CaO content, and (iii) high CaO:P2O5 ratio . Although silica based bioactive materials have shown great clinical success in many dental and orthopaedic applications, its insoluble properties has resulted in it as a potential for a long term device and the long term reaction to silica, both locally and systematically is still unknown.  Therefore, silica free, calcium phosphate glasses have attracted much interest due to their chemical and physical properties. They offer a more controlled rate of dissolution compared to silica containing glasses, they are simple, easy to produce, biodegradable, biocompatible, bioresorbable due to their ability to completely dissolve in an aqueous environment and have excellent bioactivity, osteoconductivity as well as not causing an inflammatory response. Due to their properties, especially due to it being bioresorbable, calcium phosphate glasses have been under investigation for several applications in the dental field, particular as implants. However only certain calcium phosphate compounds are suitable for implantation in the body, compounds with a Ca:P ratio less than 1 are not suitable for biological implantation due to their high solubility.
The structural unit of phosphate glasses is a PO4 tetrahedron. The basic phosphate tetrahedra form long chains and rings that create the three-dimensional vitreous network. All oxygens in the glass structure are bridging oxygens (BO), and the non-bridging oxygens (NBO) can be formed by including other species such as CaO and Na2O or MgO. Do to the effects of Ca2+, Na2+ and Mg2+ in the glass structure; they are defined as glass network modifiers, which form the glassy state and are called ‘invert glasses.’ Hence the structure of phosphate glasses can be described using the Qn terminology, where n represents the number of bridging oxygen’s that a PO4 tetrahedron has in a P2O5 glass, every tetrahedron can bond at three corners producing layers of oxygen polyhedra which are connected together with Van der Waals bonds. When the PO4 tetrahedron bonds with three bridging oxygens, giving the Q3 species, it is referred to as an ultraphosphate glass, which usually consists of a 2D network. When it bonds to two bridging oxygen’s, usually in a 3D-network it gives the Q2 species, it is referred to as metaphosphate glass. Further addition gives Q1 species, also called pyrophosphate glass, which bonds only to one bridging oxygen. Finally, the Q0 species do not bond to any bridging oxygen and hence is known as an orthophosphate glass. 
A large number of calcium phosphate glass compositions have been studied in order to exhibit suitable properties for use in biomedical applications until now, and they can be categorised into four groups:
Calcium phosphate glasses containing Potassium
Calcium phosphate glasses containing Magnesium
Calcium phosphate glasses containing Sodium and Titania
Calcium phosphate glasses containing Fluorine and Titania
1) Calcium phosphate glasses containing Potassium:
Dias et al (2003)  conducted a study and prepared bioresorbable calcium phosphate glass-ceramics between the metaphosphate and pyrophosphate region based on the composition 45CaO-45P2O5-5K2O-5MgO (Ca:P = 0.5). XRD results showed that addition of nucleating agents, K2O and MgO forms bioactive: β-CPP and biodegradable phases: KCa(PO3)3, Ca4P6O19 as well as β-Ca(PO3)2 which is considered to be non-toxic. DTA results showed two crystallisation peaks, Tp at 627°C and 739°C and two melting temperatures, Tm at 773°C and 896°C which was thought to be due to the partial melting of the crystalline phases or residual glass matrix. The glass transition temperature, Tg was observed at 534°C. FTIR results showed functional groups corresponding to metaphosphate and pyrophosphate, (PO3)- and (P2O7)4-. These results are in accordance with functional groups of the crystalline phases identified by XRD: β-CPP, KCa(PO3)3, Ca4P6O19 and β-Ca(PO3)2. Results from degradation studies of these glass-ceramics confirmed that by controlling the overall composition of the O:P in the glass, glass ceramics with high degree of degradability can be obtained. The level of chemical degradation observed for these materials is well-above that reported in literature for bioactive ceramics that are clinically used, namely HA and TCP. It was therefore concluded that the incorporation of K2O in glass ceramics increases the solubility and also these calcium phosphate glass ceramics makes them potentially clinically helpful for promoting the regeneration of soft as well as hard connective tissue by allowing the degradability to be controlled.
A study by Knowles et al (2001)  investigated the solubility and the effect of K2O in the glass-system based on the general composition: K2O-Na2O-CaO-P2O5. The exchange of a mono or divalent ion with another of a similar charge was therefore investigated. The P2O5 and CaO content were fixed, at 45 mol% and the CaO content at 20, 24 or 28 mol% and the ratio of K2O to Na2O was varied from 0 to 25mol %. Results showed, firstly an increase in CaO content caused the solubility to decrease, as expected and confirmed from previous studies. [81,94] Secondly, for all CaO contents there was an increase in solubility, when K2O content was increased.  In a recent study by Marikani et al (2008), based on the same general composition, they demonstrated that the addition of K2O caused a decrease in both density (from 2.635 g cm-3 to 2.715 g cm -3 and microhardness measurements (from 257 to 335 HV) and hence weakens the structure. These findings are attributed to the replacement of lighter cation (Na2O) by a heavier one (K2O). The ionic radius of potassium is larger than the ionic radius of sodium so, the addition of K2O has a larger disrupting effect on the structure and hence weakens the glass-network. The decrease of melting point with the addition of K2O content indicates that K2O increases network disruption by producing non-bridging oxygens. And the low value of Tg indicates that the glass samples are thermally unstable. Additionally, the elastic modulus, decreases when the concentration of K2O is increased, which implies the weakening of the overall bonding strength, as more cross linking is degraded. The increase of the internal friction and the decrease of the thermal expansion coefficient with the addition of the K2O content are due to the formation of non-bridging oxygen ions. The SEM micrographs of the glass samples recorded before immersion in SBF indicates the amorphous nature of the materials and when glasses were immersed in SBF solutions for 10 days, the glass-samples showed bioactivity.
Although the addition of K2O to the ternary Na2O-CaO-P2O5 based system offers greater flexibility in terms of tailoring the solubility to suit potential biomedical applications, only little research has been conducted in using K2O in calcium phosphate glasses, probably because it has shown to increase network disruption which was confirmed by decrease in Tm, addition of K2O causes a decrease in density and microhardness measurements, it weakens overall bonding strength confirmed by a decrease in the elastic modulus, causing it to be less rigid as well as producing thermally unstable glasses which was confirmed by the low values of Tg. These mechanical properties are not desirable in the long run and due to it being less rigid, it would not withstand stress in biomedical applications and consequently result in failure.
2) Calcium phosphate glasses containing Magnesium:
Research into calcium phosphate glasses which produce biocompatible and bioactive phases has generated a lot of interest.
–An attempt to induce β-TCP was undertaken by Zhang et al (2000) on calcium phosphate glass-ceramics in the pyrophosphate region based on the composition 50CaO-40P2O5-7TiO2-1.5MgO-1.5Na2O (Ca:P molar ratio = 0.625). XRD results showed that the β-TCP phase was not detected and the main crystalline phase precipitated was β-CPP with smaller amounts of soluble Calcium titanophosphate, CaTi4(PO4)6 CTP, and Sodium titanophosphate, NaTi2(PO4)3. Kasuga et al (1998) reported a similar occurrence in the structure of glass-ceramics which contained TiO2 (wt 3 %). . SEM observations demonstrated light areas which were confirmed by EDS analysis to be β-CPP, grey areas was thought to correspond to Na- containing phases and dark areas were composed of lower CaO contents compared to the other two areas and contained MgO and Na2O. These results were identical to Kasuga et al’s study (1999). The undetectable β-TCP phase was possibly due to the low content of MgO and TiO2 added and the low Ca:P ratio of the glass. Although bioactive and biosoluble phases were precipitated in the glass-ceramic, no continuous apatite layer was formed even after 8 weeks of immersion in SBF solution.
–A study by Brauer et al (2007) observed the solubility of several phosphate glasses in the system P2O5-CaO-MgO-Na2O-TiO2. The glass compositions ranged from ultraphosphate glasses (with phosphate contents over 50 mol %) to polyphosphate glasses (containing 50 mol% P2O5 or less which are formed by phosphate chains or rings possessing different chain lengths) to invert glasses (pyrophosphate glasses- P2O5 concentrations of around 34 mol %.). Results showed that the phosphate glasses showed a uniform dissolution. No selective alkali leaching, which is known from silica based glasses, was observed. Also that the solubility of the glasses strongly depend on the glass-composition. The higher the phosphate content resulted in an increase in solubility; According to Vogel et al , this is due to the polymerisation of the phosphate chains and the Q1 end units being more susceptible to hydration and subsequent hydrolysis than Q2 middle groups. Also it was observed that the higher the concentration of Na2O resulted in an increase in solubility too due to the effect Na+ has on the glass structure. Addition of titanium oxide resulted in a decrease in both the solubility and the tendency of the glasses to crystallise by forming cross links between phosphate groups and titanium ions. Invert glasses showed a considerably smaller solubility than polyphosphate glasses and offer an alternative to polyphosphate glasses, since they are more stable to moisture attack. However, decreasing the P2O5 content makes glasses not only more stable to hydrolysis but also restricts the glass forming area. Hence, glasses in the pyrophosphate region show a larger tendency to crystallize than polyphosphate glasses . However invert glasses in the system P2O5-CaO-MgO-Na2O showed that properties such as solubility and crystallization tendency can be controlled by adding small amounts of metal oxides . Results of solubility experiments showed that the glass system investigated enabled adjustment of solubility with only minor chemical changes. This ability to control the solubility is very promising for medical application where the coordination of implant degradation and bone formation are a key issue.
A study by Dias et al (2005) studied the crystallisation of the glass-system: 37P2O5-45CaO-5MgO-13TiO2 (Ca:P=0.6)in the pyrophosphate and orthophosphate region, by using TiO2 as a nucleating agent and MgO as a network modifier. Results showed that they contained four different crystalline phases; two of them, β-CPP and CTP are reported to be biocompatible and bioactive, respectively [88,97,98]. The biocompatibility of the other two phases Titania pyrophosphate: TiP2O7 and α-CPP has not been clearly reported so far in the literature. It appeared from XRD studies that the formation of α and β-CPP was a simultaneous process. Usually, α-CPP is observed on heating β-CPP to temperatures greater than 1150°C however for the glass composition used in the current study, it crystallised at the same time as β-CPP at only 620°C. A further study by Dias et al (2007) based on the same general composition confirmed these four crystalline phases by XRD patterns. Furthermore, they reported the chemical degradation of these glass-ceramics are in between that reported in literature for bioactive ceramics that are clinically used such as HA and TCP. Hence opportunities exist to use these glass-ceramics as bone-graft in certain clinical applications; i.e. cranioplasty, where relatively slow resorption of the implant and replacement by bone is required.
The addition of magnesium and titania to calcium phosphate invert glasses was compared by Kasuga et al (1999)  using the compositions; 60CaO-30P2O5-7Na2O-3TiO2 and 60CaO-30P2O5-7Na2O-3MgO in mol %. (Ca: P = 1). Results had shown that when both compositions were immersed in SBF, only the composition containing titania produced a calcium phosphate phase on the surface of the glass-ceramic whereas the magnesium composition did not. This was further confirmed in a later study by Zhang et al who also used TiO2 as well as Na2O and MgO in the composition.
Calcium phosphate glasses containing magnesium have the advantage of control over it’s solubility, as well having good chemical degradation which is reported to be in between those of bioactive ceramics clinically used, such as TCP and HA. Furthermore they have shown to produce bioactive and biocompatible phases, [9,23,24,26] However, there has been only little research in addition of magnesium to calcium phosphate glasses when compared with the addition of titania, due to the fact that magnesium containing calcium phosphate glasses do not produce a calcium phosphate phase when immersed in SBF and hence are not useful as bioactive materials for bone implants. Furthermore, titania has been shown to enhance mechanical properties of calcium phosphate glasses as well as improving it’s chemical durability and shows other useful properties which will be discussed later on.
3) Calcium phosphate glasses containing Sodium (Na2O) and Titania (TiO2):
Kasuga et al, has extensively researched into calcium phosphate glasses in the pyrophosphate region. Kasuga, prepared new types of phosphate invert glasses with compositions of high CaO and low P2O5 contents to obtain bioactive materials; the glasses consist of PO43- (orthophosphate) and/or P2O74- (pyrophosphate) ions without PO32- (metaphosphate) ion and the phosphate groups are connected through Ca2+ ions, which acts as a network modifier.
Kasuga (1998) for the first time, obtained SiO2-free calcium phosphate glasses in the pyrophosphate region by introducing small amounts of Na2O and TiO2 (totalling 10 mol %.) The majority of his research on calcium phosphate glasses has been based on the composition: 60CaO-30P2O5-7Na2O-3TiO2 (Ca:P = 1).
Kasuga’s XRD analysis has shown that calcium phosphate glasses containing sodium without titania based on the composition: 60CaO-30P2O5-10Na2O does not produce a β-TCP phase and only forms β-CPP, β-NaCaPO4 and 4CaO-3P2O5 (tromelite) phases. [3,10] It has been shown that the addition of titania to the glass 60CaO-30P2O5-7Na2O-3TiO2 (Ca:P = 1) resulted in the formation of large crystals of both β-CPP and β-TCP. Furthermore, it was shown that calcium phosphate glasses with the composition: xCaO-(90-x)P2O5-3TiO2-7Na2O that had a CaO content of 60-62 mol % contained β-CPP and β-TCP phases however glasses with CaO content of 55 and 58 mol % formed only β-CPP without β-TCP nor a PO3 (metaphosphate) group. Hence, calcium phosphate glasses must have a high CaO content in order to contain orthophosphate and pyrophosphate groups without metaphosphate groups and the addition of titania to these glasses induced formation of β-CPP and β-TCP crystals. Further XRD analysis by Kasuga et al (2001), on machinable calcium pyrophosphate glass-ceramics, had shown that when glass powder compacts of 60CaO-30P2O5-5TiO2-5Na2O (Ca:P = 1) were heated between 850-1000°C, both β-CPP and β-TCP were observed and with increasing temperature, the peak intensities of β-CPP was found to decrease slightly, even though β-CPP crystal could not be melted below ≈1350°C.  Kasuga et al’s study on calcium phosphate glass-ceramics for bioactive coating on a titanium alloy revealed that the surface of the coating of a titanium alloy consisted predominantly of β-TCP and β-CPP crystalline phases. It also showed that 60CaO-30P2O5-7Na2O-3TiO2 glass-ceramic layer has a greater amount of β-TCP than in the 50CaO-40P2O5-7Na2O-3TiO2 glass-ceramic and that (Ca0.5,Na)Ti2(PO4)3 crystals also precipitated. 
Differential thermal analysis (DTA) from Kasuga’s studies (1998,1999) based on the composition 60CaO-30P2O5-7Na2O-3TiO2 (Ca:P = 1) observed a slight shrinkage of the glass powder compacts at 600-800°C which was believed to be due to the viscous flow of the glass powders. Also, a large shrinkage occurred when the temperature was elevated over ≈800°C and stabilised around 850-870°C. Two crystallisation peaks, Tp’s were observed between 630-750°C and a melting temperature, Tm at ≈800°C which was thought to be due to the partial melting of the crystalline phase.[9,10] Kasuga’s further studies (2002 and 2003) [101, 102] on calcium phosphate glasses and their coating on titanium alloys based on the same composition and another, (50CaO-40P2O5-7Na2O-3TiO2) revealed very similar DTA results in both exothermic and endothermic peaks and also had shown that both powder compacts sinter at 800°C. [101, 102]A study on machinable calcium pyrophosphate glass ceramics, (2001)  based on the composition: 60CaO-30P2O5-5TiO2-5Na2O (Ca:P = 1), showed a slight shrinkage of several per cent when the temperature was elevated above 780°C and large shrinkage of the powder compacts began at 800°C and was stabilized around 850-870°C. These results may possibly suggest that crystallisation of the glasses occur at 640-730°C resulting in the formation of β-CPP and β-TCP crystalline phases with a residual phosphate glassy phase containing TiO2. When the temperature was elevated over 780°C the CPP phase precipitated in the crystallised product was partially melted by a reaction with the glassy phase. The formed melt flowed viscously around 800°C, which resulted in the sintering of the powders. Although the densification was practically completed at 850-870°C, the partial melting of the CPP phase proceeded gradually. At temperatures above 900°C, no crystallisation of new phases and no phase transformation of the precipitated crystals were observed and the gases involved in the melt may be emitted to form pores. Therefore the optimum heating temperature for the glass was determined to be 850°C.
Results from Kasuga’s studies have shown peaks at 960-1060cm-1 which corresponds to PO43- group of calcium orthophosphates such as HA and peak at ≈3500cm-1 corresponding to the OH group, both peaks increased after soaking in SBF. [9, 10] Based on another study on the composition, xCaO-(90-x)P2O5-3TiO2-7Na2O, when x = 45 (Ca:P = 0.5) strong peaks were observed at ≈700cm-1 and ≈1170cm-1, which is due to the PO3 (metaphosphate) group. . In the glass x =55, peaks at 748cm-1 and 1044cm-1 were observed, which was due to the motion of bridging oxygen and the latter peak is due to non bridging oxygens in the Q1 tetrahedra which demonstrate the existence of the pyrophosphate group. It was also found that when the CaO content in the glass increased above 58 mol% an additional peak appeared at 950-960cm-1 which is assigned to the symmetric stretching mode of non-bridging oxygen atoms in the Q0 tetrahedra, orthophosphate group. In all of the glasses (x= 45-62) weak peaks at 890-900cm-1 were observed which are ascribed to the Ti-O stretching vibrations in TiO4 units that contain non-bridging oxygens. Furthermore, at ≈630cm-1 shoulders were observed in all the glasses, which may be due to the TiO6 group. Therefore the study confirmed that calcium phosphate glasses containing CaO ≥ 55 mol % with small amounts of Na2O and TiO2 had no long chain phosphate structure and were invert glasses with orthophosphate and pyrophosphate groups. Increasing the amount of TiO2 with decreasing Na2O in the composition: 60CaO-30P2O5-yTiO2-(10-y)Na2O caused an increase in the intensity of peaks ≈890cm-1 which is due to TiO4 and ≈650cm-1 due to TiO6. Furthermore, in the composition, 60CaO-30P2O5-7Na2O-3TiO2 a peak at ≈960 was observed due to PO4: orthophosphate group as well as P2O7 peaks due to the pyrophosphate group and no metaphosphate group was detected.
In another study by Navarro et al (2003),  on the glass-system, P2O5-CaO-Na2O (Ca:P = 0.5), with titania addition, showed that titania free glasses displayed the characteristic shifts of metaphosphate glasses: the PO2 asymmetric stretch at ≈1260cm‑1, the PO2 symmetric stretch at ≈1170cm-1 and the P-O-P symmetric stretch near 690cm-1 both of which corresponded to metaphosphate groups, and the PO3 symmetric stretch at 1040cm-1 that belongs to pyrophosphate groups. The rest of the glasses which had titania, contained the peaks aforementioned as well as new bands; the TiO5 stretch at 900cm-1 and the TiO6 units stretch at 630cm-1. It was also observed that these new bands increased in intensity as the amount of titania added to the glasses increased. 
Kasuga’s SEM observations based on the composition: 60CaO-30P2O5-7Na2O-3TiO2 showed that due to the different atomic number contrast of phases precipitated, there exist bright and dark portions in the images for the glass-ceramic. Phases A and B are bright portions whereas C is the dark portion. A and B phases consist of calcium phosphates with small amounts of Na. Phase A is larger in Ca content than phase B and hence, phase A corresponds to β-TCP and phase B to β-CPP. Phase C, the dark portion contains larger amounts of Ti with Na with lower Ca content than phases A and B. Phase C is believed to be caused by the partial melting due to the densification of the glass powders. Therefore the bright portions observed are the calcium phosphate crystalline phases and the dark portions are the glassy phases containing TiO2 and Na2O.
In another of his studies (2001)  based on the same composition, in an attempt to improve the mechanical strength of the glass-ceramics, he heated the glass compacts which had high density (93.5%) to 850°C and low density glass compact (89.5%) to 950°C. SEM observations showed that the glass compact heated to 850°C (high density), densified relatively well however contained numerous small pores, which was detected from one of his earlier studies when it was too heated to 850°C. [6, 10] However the glass compact heated to 950°C (low density) showed less pores (round in shape) but they were larger in size than when heated to 850°C. It is thought that the formation of the pores may be related to flow of a liquid phase formed by partial melting of the crystals precipitated in the glass ceramics at ≈ >800°C. When the temperature is elevated to >850°C, gas bubbles in the glass-ceramic may be taken in the liquid phase. Since the viscosity of the phase decreases with increasing heat temperature the bubbles would gather to reduce their surface area; resulting in the formation of large pores. Therefore, the bending strength, σf of brittle materials is related to the fracture toughness, KIC and the critical flaw size, C in the materials. The size C, for semicircular surface flaw is estimated from σf and KIC to be ≈250μm and is ≈1.7 times higher than the value for 850°C (≈150μm). The glass-compact heated to 850°C showed fine sized calcium phosphates embedded in the glassy matrix phase which can be seen in the glass-ceramic, whereas at 950°C the SEM images showed each phase being extremely developed.
In order to eliminate the pores, a hot-pressing technique was applied to the sintering of the glass-ceramic which resulted in an increase in relative density by 98.7%. As well as this, residual pores were drastically reduced in number and in size compared to before. The bending strength, σf also increased by ≈160MPa and this value is close to that of natural cortical bone. The KIC was not improved (≈1.7MPa.m0.5) Therefore, as the mechanical strength and toughness of the glass-ceramics is relatively larger than β-TCP and HA ceramics, it is promising results for the future and it was found that one of the optimum conditions to prepare high strength glass-ceramics was by a hot-pressing technique.
Kasuga’s study (2001)  whereby he first reported machinable silica free glass-ceramics has shown that SEM observations based on the composition: 60CaO-30P2O5-5TiO2-5Na2O of the fracture face heated at 850°C showed almost no pores whereas above 900°C many large pores were observed in the samples. At 850°C, plate shaped products of several ten nanometers in thickness were piled up and were interlocked with one another. This observed morphology is similar to that of fracture face of mica-based machinable glass-ceramics; which have been shown to have good machinability. SEM observations of the polished face showed presence of bright and dark portions due to the difference in the atomic number contrast of the formed phases. When heated for 3 hours, slightly dark portions of < 1μm in size were distributed. When heated for 8 hours, contrast difference was clearly seen and microstructure developed exceedingly, the grains were grown into an interconnecting 3D-network. EDS analysis confirmed that the bright grains are a calcium phosphate phase, indicating the presence of β-CPP crystal. Although the presence of β-TCP crystal was not found in the photo, it was however confirmed by XRD. The amount of β-TCP crystal may have been relatively small and titanium was included in dark portions in the photo.
4) Calcium phosphate glasses containing Fluorine and Titania
It is understood that in order to improve the chemical durability of calcium phosphate glass-ceramics, one of the ways is to anticipate the formation of fluoroapatite, which has the lowest solubility among various calcium phosphate crystals. Kasuga et al (2007)  investigated a novel calcium phosphate glass-ceramic for a dental filler based on the composition: 40CaO-25TiO2-30P2O5-5CaF2 [CTP-F] (Ca: P = 0.67) and for a comparative study he also prepared a glass-ceramic without fluorine: 45CaO-25TiO2-30P2O5 [CTP]. He found that the glass ceramic, CTP-F had good chemical durability. He later prepared (2009) a calcium titanium phosphate glass-ceramic with improved chemical durability than CTP-F and this glass-ceramic was prepared by including a small amount of fluorine based on the composition: 35CaO-10CaF2-30P2O5-25TiO2 in mol% [CFPT] (Ca:P = 0.58) and he used CTP again for comparison.
Kasuga’s XRD analysis on calcium titanium phosphate glass-ceramic with and without fluorine, CTP-F and CTP respectively showed peaks assigned to Nasicon-type crystal: CaTi4(PO4)6 and peaks assigned to Ti(PO3)3 and (TiO)2P2O7 crystals. Once CaF2 was added to the composition, apatite crystal was formed preferentially and CTP-F showed a weaker peak intensity of CaTi4(PO4)6 crystal than in the CTP glass-ceramic.  A later study by Kasuga et al (2009) showed that the glass-ceramic CTP contained crystalline phases such as α-Ca2P2O7, CaTi4(PO4)6, TiO2 (anatase), and (TiO)2P2O7, whilst CFPT contained only Ca-Ti4(PO4)6 and apatite phases. The XRD peaks of CFPT glass-ceramic were sharp and their intensities were strong compared to the CTP glass-ceramic.
DTA curves from Kasuga’s study (2009) showed the glass transition temperatures, Tg of CTP and CFPT glasses to be 675°C and 615°C, respectively. The crystallisation temperatures were 785°C and 685°C of CTP and CFPT glasses, respectively and the lower temperature was due to fluorine inclusion.  In many cases, the chemical durability of glasses and glass-ceramics deteriorates upon including fluorine. However in Kasuga’s study the crystallisation behaviour of the CTP glass was drastically influenced by including a small amount of fluorine, and the resulting glass-ceramic showed excellent chemical durability.
SEM micrographs of CTP-F glass-ceramic surface demonstrates small-sized pits which are believed to originate from dissolution of crystalline and/or glassy calcium phosphate phases in the glass ceramics.
SEM images of CTP and CFPT glass-ceramics show different morphologies. In the CPT glass-ceramic, the characteristic morphology, which is suggested to result from a spinodal-type phase separation of the glass before crystallisation, is observed. Earlier studies revealed that, in this kind of glass-ceramic, a calcium phosphate phase and a Nasicon-type crystalline phase such as CaTi4(PO4)6 are three-dimensionally intertwined, and the calcium phosphate phase is easily dissolved and removed by acid treatment. [108-109] On the other hand, in the CFPT glass-ceramic numerous independent pits several tens of nanometers in size are observed just as in the previous study. The composition of the surface from EDS analysis was suggested to be the CaTi4(PO4)6 phase and it was suggested that the phase that was leached out is the apatite, which is embedded in the CaTi4(PO4)6 phase. 
During the heat treatment of CFPT glass, apatite is considered to form preferentially, and subsequently, or simultaneously, a large amount of the CaTi4(PO4)6 phase forms around the apatite. As a result, apatite particles several tens of nanometers in size would be apparently embedded in the CaTi4(PO4)6 phase. The orthophosphate group in the glass is consumed in the formation of not only the CaTi4(PO4)6 phase but also the apatite phase. Because of the formation of these phases, the amount of phosphate would decrease, resulting in the controlled formation of crystalline phases such as titanium phosphate and calcium pyrophosphate in the CTP glass-ceramic. In Kasuga’s earlier studies, no apatite formation was observed in fluorine containing calcium-phosphate-based glasses consisting of the meta- and/or pyrophosphate groups without the orthophosphate group. Therefore preferential apatite formation was suggested to occur in glasses containing the orthophosphate group.
Furthermore, EDS analysis had shown that in the glass-ceramic CFPT, fluorine was partially evaporated and 2.8% of residual fluorine was contained in the glass out of the 6% of fluorine initially used in the composition.
Kasuga’s FTIR analysis demonstrated that the calcium phosphate glass-ceramics with and without fluorine, CTP and CTP-F respectively show peaks corresponding to orthophosphate, pyrophosphate, and metaphosphate groups. The glass-ceramic containing fluorine (CTP-F) show sharper peaks than the CTP glass-ceramic which may imply that there is a larger amount of crystalline phases in CTP-F than there is in the CTP glass-ceramic. Furthermore, the CTP-F glass ceramic has larger intensities of peaks that correspond to the orthophosphate group than CTP. Generally the intensities of the peaks observed due to pyrophosphate group in both the spectrums are very small. In the CTP-F glass-ceramic, the peak corresponding to anatase is observed as supported by XRD. In the CTP glass-ceramic spectrum, a shoulder peak around 900cm-1 due to Ti-Onb group (Onb; non-bridging oxygen), is observed which is often seen in the glassy phase. It is thought that even after crystallisation, a relatively large amount of titanium constituent is included in the glassy phase.  In his later study (2009), it was shown that both CTP and CFPT glasses showed peaks corresponding to the orthophosphate group (PO43-), the pyrophosphate group (P2O74- group), and the metaphosphate group (PO3- group), which is in accord with Kasuga’s earlier study of 2007. [106-107]
Furthermore, it was confirmed from the 19F MAS-NMR spectra of CFPT, that the apatite in the glass-ceramic was a fluorine-containing oxyapatite, Ca10(PO4)6(O, F2).
Kasuga et al (2007)  found that the glass ceramic, CTP-F had good chemical durability which was suggested to originate from an increase in the amount of crystalline phases and a high content of titanium constituent in the residual glassy phase. It was observed that the incorporation of fluorine to the glass induced the apatite formation in the glass and the apatite phase was an oxyapatite crystal containing fluorine. (fluoro-oxyaptite) . He later prepared a novel CFPT glass-ceramic which showed even more improved chemical durability than CTP-F. It was shown that the structure of the glass-ceramic was considerably influenced by a small amount of fluorine included in the calcium phosphate glass. Fluorine-containing oxyapatite, Ca10(PO4)6(O, F2) and Nasicon type, CaTi4(PO4)6 phases were preferentially formed during heat treatment and apatite particles several tens of nanometers in size were embedded in the Nasicon-type phase. Titanium ions were also included in the residual phosphate glassy phase. The microstructure of this glass-ceramic was suggested to result in the excellent chemical durability and is promising results in order to design novel dental fillers.
Benefits of Titania addition to calcium phosphate glasses:
A lot of interest has been shown on the role of titania in calcium phosphate glasses. The titania content serves as an effective bulk nucleating agent in calcium phosphate glass (TiO2 above 4 wt %). Studies have confirmed that when TiO2 is added to calcium phosphate glasses, it improves its glass-forming ability, it improves the chemical durability of the glass, [13,14] its modulus elasticity,  it raises the glass viscosity, which was shown by an increase in the glass-transition temperature, Tg and consequently strengthens the glass network. Furthermore, studies have shown that when TiO2 is added into calcium phosphate glasses, it increases the bulk density of the glass, it decreases the solubility of the glasses well as decreasing its tendency to crystallise.
It is understood that the incorporation of titanium ions into the phosphate network causes a structural change. The phosphate network changes from chain extending metaphosphate tetrahedra to chain-terminating pyrophosphate groups, which is expected when the O:P ratio increases in these types of glasses. (Based on the general composition: CaO-P2O5-Na2O)  Therefore, TiO2 incorporation is associated with the formation of TiO5 or TiO6 titanate polyhedra between the phosphate glass chains. As, the titanium ion has a small ionic radius and a large electrical charge compared to Na+, it has the ability to penetrate into the glass network in an interstitial position. This means that titanium has the capacity to enter into the vitreous arrangement and place itself between the phosphate chains and rings, creating cross-linking between Ti4+ ions and the phosphate tetrahedra. Therefore addition of TiO2 into phosphate glasses increases the O:P ratio, creating stronger Ti–O–P cross-links in the glass network. As a result, the glass becomes rigid, more resistant to hydration subsequently increasing its elastic modulus and chemical durability.
A study by Neel at al (2008)  studied the effect of increasing titanium dioxide on calcium phosphate glasses based on the general composition: CaO-P2O5-Na2O. Glasses of 15 mol % TiO2 were successfully obtained however, glasses with 20 mol.% could not be prepared as they crystallised easily during cooling. By incorporating 15 mol % TiO2 into the glasses, the bulk density increased which was due to the replacement of the low-density element Na+ with the high-density element Ti4+. Analysis of DSC data showed that the addition of TiO2 results in a significant increase in Tg values. DTA results showed that the glass containing 5 mol % TiO2 showed two crystallisation peaks. However the glass-free TiO2, and the glasses containing 10 and 15 mol % TiO2 showed only one crystallisation peak, in the region 490-580°C, 705-810°C and 725-815°C respectively. Addition of 10 and 15 mol.% TiO2 produced glass structures with two broad melting peaks ranging from 790 to 890°C and 835 to 935°C, respectively, instead of a single sharp melting peak ranging from 705 to 775°C as seen for the ternary glass with no TiO2. 
Ternary glass with no TiO2 produced a NaCa(PO3)3 phase which was identified from XRD. Glasses with 5 mol.% TiO2, produced two phase: NaCa(PO3)3 was the main phase, and TiP2O7 was the second phase. The glass with 10 and 15 mol.% TiO2, another phase, β-CaP2O6 was also identified. For glass containing 15 mol.% TiO2, TiP2O7 became the main phase, instead of NaCa(PO3)3 which was the main phase for all other compositions.
The glasses containing TiO2 (5, 10 and 15 mol %) showed one order of magnitude reduction in degradation rate compared to TiO2 free glass which suggests that addition of TiO2 beyond 5 mol % does not produce further significant reduction in the degradation rate compared to 5 mol.% TiO2-containing glasses. Hence we can propose that there is an optimal amount of Ti4+ ions to be incorporated into the glass network to reduce the degradation, beyond which addition of more Ti4+ ions has no effect on degradation rate. 
Furthermore, addition of titania to calcium phosphate glasses has shown to induce formation of β-CPP and β-TCP during heat treatment. and is known to produce a calcium phosphate phase (apatite layer) on the surface of the glass ceramic and hence has shown bioactivity. Addition of titania has shown to improve its chemical durability and it is thought that the apatite formation is influenced by the durability of the glassy phase in the glass-ceramics.Apart from this, the glass-ceramic containing titania has shown relatively high fracture toughness, KIC (1.9 MPa m0.5) and bending strength (100-120 MPa). The toughness is approximately twice larger than that of TCP or HA ceramics. The glass-ceramic containing titania has enhanced mechanical properties as well as toughness, which is promising as a novel glass-ceramic with bioactivity for biomedical applications.
Benefits of Fluorine addition to calcium phosphate glasses:
Due to the properties of glass and glass-ceramics being determined critically by their composition, the crystallisation process and the microstructure, both the design and characterisation of these glasses play very important role on the development of new medical glasses and will be discussed in this chapter.–FEI
Some studies have reported that using appropriate heat treatment the microstructure can be controlled by controlling the phases and sizes of crystals that are precipitated in the glasses [g]
Middle of Review?
Calcium phosphates are now used for a variety of different applications covering all areas of the skeleton including spinal fusion, craniomaxillofacial reconstruction, treatment of bone defects, fracture treatment, total joint replacement (bone augmentation) and revision surgery.
Calcium phosphate implants (and hydroxyapatite, in particular) are used in the form of coatings on metallic implants, as fillers in polymer matrices, as self setting bone cements, as granules or as larger, shaped structures.
Disadvantages of calcium phosphate glasses:
Brittleness- cannot be used in load bearing applications. As a result they are mostly used as coatings on metallic substrates. Potential to fail catastrophically, difficulty to machine.
§ MOST COMMON RESEARCHED TPYES OF CaP glasses—-Among the bioactive ceramics most commonly investigated for bone substitution are β-TCP, HA and bioactive glasses. TALK ABIT ABOUT BONE COMPOSITION! NOT ALL Ca:P are suitable for implants only > 1!!!
§ Talk about compositions of HA, b-TCP, b-CPP in relation to their properties i.e-biocompatiblity etc
§ Talk about applications of each type of CaP materials, i.e HA, etc [LIT REVIEW ON Ca:P glasses having app’s in dental field— and relating to what was found in terms of solubility, hardness, biocompatibiltiy etc [3-4JOURNALS MAX!!!]
Since the discovery by Hench et al (1971), the most commonly investigated types of calcium phosphate ceramics are HA, β-TCP, β-CPP which have been found to be bioactive and are already in clinical use for bone substitutes. . However not even these bioactive materials can replace all bone autografts. The need to develop bioactive materials with a higher bonding ability as well as different mechanical properties is essential.
Different calcium phosphates exhibit widely different resorption properties. The resorption property of calcium phosphates depends on Ca/PO4 ratio, degree of crystallinity and crystal structure. Chin et al (1996) reported the development of resorbable Ca-P-O glass prepared from CaO and P2O5 for tooth and bone implant applications. In the present study, a highly resorbable Ca-P-O glass was prepared and this was evaluated for its biocompatibility.
But the potential of any ceramic material to be used as an implant in vivo depends upon its ability to withstand complex stresses at the site of application and its compatibility with the biological environment.
Among different forms of calcium phosphates, particular attention has been placed to
tricalcium phosphate (Ca3(PO4)2, TCP) and hydroxyapatite (Ca10(PO4)6(OH)2, HAp) due to their outstanding biological responses to the physiological environment. The contemporary health care industry uses calcium phosphate ceramics in various applications, depending upon whether a resorbable or bioactive material is ideal. The recent trend in bioceramic research is focused on overcoming the limitations of calcium phosphates, precisely hydroxyapatite ceramics and in improving their biological properties via exploring the unique advantages of nanotechnology.
Calcium phosphate based materials have also received a lot of attention due to their similarity with the mineral phase of bone, the absence of antigenicity and the excellent osteoconductivity. Such materials are hydroxyapatite (HA, Ca5(PO4)3OH ), β-tricalcium phosphate (TCP, β-Ca3(PO4)2), β-calcium pyrophosphate (CPP,β-Ca2P2O7) and fluorapatite (FAP, Ca5(PO4)3F) [34-37]. These materials have found application as fillers for bone defects, bone substitutes and as scaffolds for drug delivery due to their unique solubility.
The requirement for bioactivity
All the bioactive ceramics mentioned above, apart from beta-tricalcium phosphate form an apatite layer on their surface in the living body and hence bond to living bone through this apatite layer. This apatite layer is very similar to bone mineral in its composition and structure. Therefore osteoblasts, which are responsible for bone formation preferentially, proliferate and differentiate to produce apatite and collagen on this apatite layer and consequently the surrounding bone can come into direct contact with the surface apatite layer. When this occurs, a strong chemical bond is formed between the bone minerals and the surface apatite layer to reduce the interface energy between them. This indicates that the essential requirement for an artificial material to bond to living bone is the formation of a biologically active bone-like apatite layer on its surface in the living body.
The requirement for apatite formation
The fact that apatite formation only occurs in bone tissue in the human body is attributed to the high activation energy for the homogeneous nucleation of apatite in human body fluid. Therefore if an artificial material has a functional group that could be an effective apatite nucleation site on its surface, then it will easily form apatite nuclei on its surface. Once the apatite nuclei are formed, they would spontaneously grow by consuming calcium and phosphate ions from the surrounding body fluid. The requirement for apatite formation on an artificial material in a living body is the presence of a type of functional group that could be an effective site for apatite nucleation on its surface. These types of functional groups can include TiO2 or SiO2 gels which have shown to form bone-like apatite on their surfaces. Hence, abundant Ti-OH and Si-OH groups on the gels are effective for apatite nucleation.
1.5 Aims and Objectives
To make a Fluorine containing Calcium Phosphate Glass i.e. a simple Fluoroapatite stoichiometric glass. An early realisation demonstrated that a glass having a Ca:P of > 0.5 could not be made and fluorine was not retained in the structure. Hence the aim of my research was to make a fluorine containing calcium phosphate glass with a Ca:P of 0.5 and less, in the hope to retain fluorine by the addition of titania. As previous research has shown that calcium phosphate ceramics having a Ca:P of less than 1 are not suitable as implants as they are not compatible with bone, thus the glass-ceramic I have produced has a Ca:P of 0.5 and less, so instead it may possibly have applications as a dental material.