Kokubo et al. (1982) developed apatite–wollastonite (A–W) (β-CaSiO₃) system based on SiO₂–P₂O₅–CaO–MgO–CaF₂, also known by its commercial name Cerabone^®. The β-wollastonite phase in the A–W system enhances mechanical performance of the glass-ceramic (Table 1). Kokubo et al. (1987) reported on the A–W GC ability to resist failure by fatigue. It was estimated that the life-time of an A–W GC, under a constant loading of 65 MPa in simulated body fluid is 10 years, compared to a sintered HAp, which under the same loading can sustain the loading before fracture for only 1 min. Both, apatite and β-wollastonite phases in the A–W system crystallize through a surface mechanism (Kokubo et al., 1982). As such, the A–W GC cannot be cast to shape by the “lost wax” technique and is processed through powder sintering route.

It was often claimed that the apatite phase in the A–W systems is FAp; however, Clifford and Hill (1996) noted that the A–W glasses are very deficient in fluorine content with regard to the FAp stoichiometry. Clifford and Hill (1996) further suggested that the apatite formed in A–W systems is therefore more likely to be a mixture of fluor- and oxyapatite. Based on the pioneering electron spin resonance (ESP) studies on fluor/oxyapatites, it has been long known that O⁻ can occupy F⁻ sites (Segall et al., 1962; Piper et al., 1965). Nonetheless, at present, the availability of high-resolution solid-state characterization techniques such as ¹⁷O MAS-NMR coupled with dynamic nuclear polarization (DNP) and high field ¹⁹F MAS-NMR could provide fast and accurate elucidation of oxygen and fluorine environments in the A–W GC; however, such work is yet to be published.

Calver et al. (2004) reported that modified Kokubo et al. (1982) A–W glass with higher metal fluoride content (AW3, Table 2) resulted in a completely changed apatite crystallization behavior. Calver et al. (2004) found that a base glass with the highest CaF₂ content favored volume FAp nucleation and crystallization. Calver et al. (2004) also found that A–W systems showed reduced T_g and FAp crystallization exotherms with increasing calcium fluoride content. Therefore, low calcium fluoride content in the original A–W system is actually suppressing crystallization of FAp. Filho et al. (1996) suggest that fully crystallized GCs, which are otherwise bioactive, will not exhibit any further bioactivity through the release of ions, such as Ca²⁺ and HPO42−, once such ions take up higher energy coordination states in the apatite lattice. Therefore, from a bioactivity point of view, it may not always be desirable to consume therapeutic ions for the formation of apatite within the glass matrix, as opposed to the release of such ions and subsequent formation of apatite on the surfaces of the material, that potentially lead to the formation of a strong and chemically stable implant–bone interface. Furthermore, the molar Ca to P ratio in a stoichiometric apatite crystal is 10 to 6 (1.67); therefore, GC systems containing a stoichiometric Ca to P ratio, or in other words, glass formulations with smaller compositional differences between the glass and the crystal phase preferentially crystallize in bulk, which is unfortunately not the case in the A–W system. On the other hand, the A–W system is inherently aluminum free; thus, risks associated with aluminum neurotoxicity, summarized by Kumar and Gill (2009) and reported by Reusche et al. (2001), are completely absent.

It could be argued as to why the original A–W system developed by Kokubo et al. (1982) contains magnesium and a non-stoichiometric Ca:P:F ratio that could otherwise aid bulk nucleation and crystallization of apatite phases. There is a considerable and long-standing proof, for example, as found by X-ray diffraction analyses of precipitated apatites by LeGeros et al. (1980), which demonstrates that the presence of Mg²⁺ ions in an aqueous solution cause a strain on the apatite structure causing it to collapse and, therefore, suppress its growth. In view of the biological apatite found in bone as opposed to tooth, Mg²⁺ ions alongside osteocalcine and proteoglycan proteins play a significant role in the development of a nanoscale apatite (Blumenthal et al., 1975), providing the bone tissue with a fine microstructure and the excellent properties that come with it. Therefore, on that basis, it could be argued that the A–W system is highly biomimetic in view of the elemental composition of the human bone.

In vivo animal studies on the implanted A–W glass-ceramic provide evidence for excellent osseointegration around the A–W implant and a chemical calcium phosphate-based interface (Kitsugi et al., 1989, 1990) between the implant and bone, with high bending and compressive strengths of 157 and 1060 MPa, respectively (Nakamura et al., 1985). Kokubo et al. (1990) found that A–W GC immersed in tris(hydroxymethyl)aminomethane (TRIS) buffer did not show bioactivity through the formation of apatite and these findings were contradictory to animal studies conducted previously, where the A–W GC was found to form a strong chemical interface with the living bone. Therefore, Kokubo et al. (1990) developed a simulated body fluid (SBF), an alternative immersion medium for in vitro assessment of bioactivity of the A–W GC. Kokubo et al. (1990) argued that TRIS buffer does not mimic the actual body environment because it is completely deficient in ions, such as Ca²⁺ and HPO42− naturally found in the bodily fluids and, therefore, argued that the lack of apatite formation in TRIS buffer, as opposed to high bioactivity of the A–W GC in SBF, also supports the view that the apatite phase on the surfaces of the A–W GC forms by a chemical reaction between the A–W GC and the ions present in the body fluid. In view of this, it can be further postulated that apatite crystals within the A–W GC act as nuclei on which ions, such as Ca²⁺ and PO43− in the solution nucleate and feed apatite formation until the eventual fusion between the apatite crystals in the A–W and the apatite in the living bone, whereby a chemical interface is formed.

Kokubo et al. (1992) reported that if aluminum is included in the A–W parent glass composition and then subsequently crystallized, the A–W GC does not show any bioactivity in SBF as opposed to the same A–W material without aluminum. The addition of aluminum results in a more chemically stable residual glass phase, which reduces the release of Ca²⁺ and PO43− ions, thereby affecting apatite formation (Strnad, 1992). Later, Blades et al. (1998) conducted in vivo animal study on aluminum-containing glass ionomer cements (GICs) as potential bone cements, where as low as 1 ppm of aluminum released was found to inhibit mineralization of the newly formed osteoid in rabbit bone. It is important to distinguish that the role of aluminum in Blades et al. (1998) study may be attributed to aluminum toxicity to bone-forming cells (Rodriguez et al., 1990) and direct inhibition of crystal growth as opposed to reduced bioactivity involving structural parameters of the residual glass phase, as proposed in Strnad (1992) study.

Good oseointegration through the formation of apatite on the implant surfaces (Neo et al., 1993) combined with good mechanical properties (Kokubo et al., 1985, 1986) perhaps explains why the A–W glass-ceramic has found promising applications in bone and vertebra replacement (Kokubo, 2008) and it is reported that over 50,000 successful bone implants have been made using the A–W glass-ceramic system (Zanotto, 2010). However, from a manufacturing point of view, a bulk crystallizing A–W material, such as proposed by Calver et al. (2004), could provide a more cost-effective material. In vitro and in vivo bioactivity of the modified A–W system developed by Calver et al. (2004) still needs to be established. The original A–W GC is currently manufactured by Nippon Electric Glass Co., Ltd. (Japan) (Montazerian and Zanotto, 2016).

Grossman (1972) of Corning Glass Works, developed the first machinable mica glass-ceramic system, later marketed by Dentsply International under the name Dicor^®. Dicor^® was seen as a very significant development since the new GC material could be easily machined to shape without a critical failure. Machinability of the mica glass-ceramics is attributed to the eminent cleavage of the mica-type crystals as a result of anisotropic crystal growth. This facilitates crack propagation in the direction of cutting without causing a critical failure of the material.

Although mica glass-ceramics initially did not contain any apatite phases, Vogel et al. (1986) developed two GC systems (Table 3) with FAp and tri/tetrasilicic Mg₃(AlSi₃O₁₀F₂)Na/K/Mg₃(Si₄O₁₀F₂)Na/K mica phases and an additional silica-free GC with apatite and aluminum phosphate phases.

Commercially available Bioverit I^® and Bioverit II^® systems crystallize to FAp and mica phases in bulk. Bioverit II^® GC crystallizes to a smaller fraction of FAp, which can be explained by the very low P₂O₅ content in the composition (Table 3). Both systems have undergone prior APS into two droplet phases and a glassy matrix phase (Höland and Beall, 2012). One droplet phase is rich in apatite elements whereas the second droplet phase is closer to the mica composition. This may explain why both phases, apatite and mica, crystallize in bulk. Bioverit III is a silica-free phosphate glass that crystallizes to FAp and an aluminum phosphate (AlPO₄) (Höland and Beall, 2012); Bioverit III material exhibits lower mechanical properties (Höland and Beall, 2012); therefore, it has not been so extensively studied.

Bioverit II contains a higher fraction of mica crystals. Therefore, Bioverit II GC exhibits better machinability but at the expense of lower mechanical properties than Bioverit I (Table 1). Based on ¹⁹F MAS-NMR experiments of mica ceramics, it was demonstrated that the fluoride ion in mica systems exists mainly in Mg(n)-F type environments, with a chemical shift at about −174 ppm for Mg(3)-F (Fechtelkord et al., 2003). As of 2016, A–FM GCs, Bioverit I and II are currently manufactured by VITRON Spezialwerkstoffe GmbH (Germany).

Mullite is a rare naturally occurring aluminosilicate mineral. Mullite-reinforced matrices exhibit enhanced mechanical properties. FAp phases can act as nucleation and crystal growth sites for new apatite phases between the implant and living bone. Both crystal phases in the fluorapatite–mullite system show elongated needle-like microstructure and exceptional mechanical properties, particularly flexural strength and fracture toughness of up to 330 MPa and 3.3 MPa √m², respectively (Table 1). The A–M GC exhibits spherulitic crystallization, which enhances fracture toughness properties of the material (Stanton et al., 2010). The system nucleates in bulk; therefore, it is readily castable by the “lost wax” route.

The first melt-derived castable FAp (Ca₅(PO₄)₃F)—mullite (Al₆Si₂O₁₃) glass-ceramics were developed by Hill et al. (1991) and were based on the SiO₂–Al₂O₃–P₂O₅–CaO–CaF₂ system. Additionally, Samuneva et al. (1998) were also able to produce an A–M glass-ceramic by a sol–gel route, rather than a melt-quench route.

Hill et al. (1991) noted that base glasses with relatively low CaF₂ content (A–C, Table 4) surface crystallized to apatite and mullite phases upon heat treatment, whereas base glasses with metal fluoride content (D and E, Table 4) bulk crystallized to anorthite with only a small fraction of FAp. The work by Hill et al. (1991) and also subsequent work by Hill and Wood (1995); Clifford and Hill (1996); Hill et al. (2000); Rafferty et al. (2000a); Clifford et al. (2001a,b), and later by Stanton and Hill (2005) suggests that metal fluoride as well as phosphorus content in fluoro-phospho-aluminate systems will likely influence prior LLPS and will therefore determine whether the glass crystallizes via the homogenous bulk route (aided by the LLPS composition) or the heterogeneous surface route.

Stanton and Hill (2005) postulated that once crystallization of FAp crystal has begun, the surrounding glass becomes depleted in F, Ca, and P and moves closer to the mullite composition, whereupon, it crystallizes to mullite by the homogenous mechanism. Stamboulis et al. (2004) explained crystallization of FAp and mullite phases in the A–M system in even greater detail by analyzing heat-treated A–M samples with magic angle spinning-nuclear magnetic resonance (MAS-NMR) spectroscopy. Spectra from ²⁷Al MAS-NMR in Figure 5 (left) show how aluminum resonance is relatively unchanged until the GC reaches second crystallization temperature (T_p2) whereupon it crystallizes to mullite. It can be observed from Figure 5 (left) that a broad peak seen at around 50 ppm, assigned to Al(IV), remains unchanged even after FAp crystallization (T_p1). During second phase crystallization, an additional peak, at around 12 ppm, which is assigned to an Al(VI) in mullite is observed. Stamboulis et al. (2004) further explain that, during first and second crystallization processes, charge balancing cations for maintaining aluminum in a fourfold coordination state, Al(IV) are consumed during FAp formation; therefore, at higher temperatures, the lack of charge balancing cations to keep aluminum in a IV coordination state forces aluminum to take up higher coordination states, in this case, Al(VI). ¹⁹F MAS-NMR spectra as shown in Figure 5 (right) show how the fluorine environment changes with increasing heat-treatment temperature. The lowermost ¹⁹F spectrum of the untreated LG120 glass in Figure 5 (right) demonstrates two broad peaks at −90 and −150 ppm that can be attributed to the amorphous fluorine environments in the untreated glass, F–M(n) and Al–F–M(n) (Zeng and Stebbins, 2000). However, at T_p1, a sharp peak at around −103 ppm develops at the expense of F–Ca(n) peak, which is assigned to fluorine in a F–Ca(3) environment in FAp. At higher temperatures, such as T_p2, fluorine from Al–F–M(n) peak is also fully consumed and Stamboulis et al. (2004) further proposed that F–M(n) such as F–Ca(n) species preferentially charge balance the non-bridging oxygens in the phosphorus locality. It can be additionally postulated that such preference to balance the non-bridging oxygens in the phosphorus locality reduces kinetic energy barrier to FAp nucleation and crystal growth, since the local environment of Ca, P, and F in the glass is similar to that present in FAp.

Relatively recently, Stamboulis et al. (2006), Hill et al. (2007), and O’Donnell et al. (2010) conducted real-time SANS and neutron diffraction (ND) experiments on the A–M systems developed in the early 1990s. They postulated that amorphous glasses that form the A–M system may have undergone phase separation by spinodal decomposition during the casting process on a scale of 25–27 nm (Hill et al., 2007). Similarly, the same cast A–M system isothermally heat treated at 740°C and 750°C initially showed neutron scattering at higher q, which then moved to lower q with increasing temperatures where the scale of the LLPS corresponded to about 35 nm. This provides evidence that undercooled liquids can undergo LLPS by nucleation and growth, whereupon the chemical system can overcome the thermodynamic and kinetic energy barriers to nucleation and subsequent amorphous phase growth. It is suggested by Hill et al. (2007) that the A–M cast glass is initially phase-separated by spinodal decomposition, whereby rapid cooling creates a barrier to nucleation and growth. However, it may be argued that scattering at lower q in the as-cast glasses may be attributed to a finer LLPS by nucleation and growth. SANS experiments by Hill and coworkers provide evidence that finer scale phase separation in the apatite–mullite system increases in size as a function of temperature, which can be either attributed to the fact that the cast A–M system is initially spinodally decomposed and then undergoes LLPS by nucleation and growth or more likely LLPS growth by an OR mechanism.

Further studies of an A–M glass-ceramic analyzed by heat treating an amorphous precursor glass up to 1,200°C with in situ TOF-ND (O’Donnell et al., 2010) explain the crystallization behavior of the A–M system in more detail. O’Donnell et al. (2010) reported that FAp and mullite crystallized on heating until 1,130°C followed by the partial dissolution of both phases at higher temperatures. It was also found that on the subsequent cooling of the A–M system, recrystallization occurred and additional new phases were produced, namely berlinite (AlPO₄) and cristobalite (SiO₂), which formed at around 1,025°C in addition to the FAp and mullite. This indicates that the subsequent cooling of the GC can produce additional crystal phases, which may be undesirable but can be avoided by introducing higher cooling rates to create an energy barrier to nucleation and growth of the undesirable crystal phases.

Stanton and Hill (2005) found that apatite phases in the A–M system grow as dendrites and spherulites. Generally, the microstructure of a GC is strongly influenced by the conditions of the heat-treatment of the parent glass, this namely includes duration, temperature, and cooling rate and whether or not the base glass compositions is doped with additional nucleants, for example, such as reported with niobium-doped FAp GCs by Denry et al. (2012). Mechanical properties, particularly fracture toughness of the A–M GC developed by Hill and coworkers, surpass mechanical properties of the alternative glass-ceramic systems discussed in this review.

In vivo animal studies conducted by Freeman et al. (2003) show that a fully crystallized A–M glass-ceramic osseointegrates with bone as shown in Figure 6B; however, its amorphous precursor base glass implant does not osseointegrate, which is evident from the fibrous tissue around the glass implant as shown in Figure 6A and the lack of implant/bone interfaces. Goodridge et al. (2007) report on both, in vitro and in vivo properties of the porous A–M glass ceramic system produced through selective laser sintering (SLS) method, using cast A–M and commercial A–W glass-ceramics as positive controls. In the 4 weeks study, Goodridge et al. (2007) reports that no sign of inflammation or adverse tissue reaction was observed around all the implants. Analyses of the implant–bone interfaces by scanning electron microscopy (SEM) showed evidence for bone ingrowth into the both porous materials, the A–M system produced through SLS method and the A–W sintered glass-ceramic. Although Goodridge et al. (2007) report that the laser sintered apatite–mullite system developed by Hill et al. (1991) does not form apatite in SBF, previous studies indicate that SBF studies are not adequate in the assessment of bioactivity of aluminum-containing GCs (Strnad, 1992).

Stanton et al. (2009) assessed the interfacial chemistry between the A–M system and a titanium alloy, for potential applications of the A–M system for orthopedic implant coatings. Stanton et al. (2009) enameled the A–M glass-ceramic to titanium by heat treatment and thereafter analyzed the interfacial reaction zone between the A–M glass-ceramic and the titanium alloy by high-angle annular dark field TEM (HAADF-TEM). Stanton et al. (2009) found that titanium diffused into the intermediate layer of the glass-ceramic and postulated that complex titanium silicides and titanium phosphides were formed based on the elemental analysis of the interfacial zones by energy dispersive X-rays (EDX), which produced characteristic photons for Ti, Si, and P elements, but not O elements. This study provides evidence that the A–M system can chemically adhere to titanium. This can be useful in overcoming the problem of coating detachment observed with micromechanical surface retention of plasma sprayed HAp coatings (Filiaggi et al., 1991). The A–M glass-ceramic coating may enhance osseointegration at the bone–implant interface and provide long-term stability of the A–M coated implants.

Wood and Hill (1991) produced cements from the A–M glass-ceramic ionomer-type systems with varying degrees of crystallinity for potential application as bone cements. Wood and Hill (1991) found that the degree of crystallinity of the A–M glass-ceramic can influence the properties of the cements, such as working and setting times of the cement pastes and the mechanical properties of the set cements.

Hill et al. (2004) developed strontium-substituted FAp glass-ceramics for potential orthopedic applications, in the system SiO₂–Al₂O₃–P₂O₅–CaO/SrO–CaF₂/SrF₂ (Table 5). Since strontium has a higher atomic number than calcium, strontium-substituted materials exhibit higher radiopacity, which enables the clinician to distinguish between the implant and bone on a radiograph.

Hill et al. (2004) observed that substituting strontium for calcium has little effect on the parent glass structure. However, the crystallization behavior of the glasses as a function of strontium content was markedly altered. Base glasses without any strontium exhibited complete bulk crystallization with the first crystallization temperature being independent of the particle size, whereas glasses with strontium and no calcium exhibited predominantly surface nucleation. Glasses with equimolar proportions of calcium and strontium crystallized through bulk and surface. Hill et al. (2004) demonstrate that increasing strontium substitution hinders bulk crystallization of apatite, which is reflected in an increase in crystallization temperature, and promotion of surface nucleation of apatite at the expense of bulk nucleation, as evidenced by particle size dependence on the crystallization temperature of the FAp phases. Hill et al. (2004) also observed that equimolar strontium–calcium composition resulted in a glass with a reduced second exotherm associated with mullite crystallization. This can be attributed to an increased mobility of the glass network and a lower glass transition temperature, which produces a more dominant heterogenous crystallization effect. Such a reduction in crystallization temperature would suggest a new crystalline phase, and not simply a reduction in crystallization temperature Hill et al. (2004).

In another study, Hill et al. (2006) elucidated Ca and Sr sites in mixed FAp through ¹⁹F MAS-NMR. Results showed the F to be present as F–Ca(3) in the all calcium glass and as F–Sr(3) in the all strontium glass. In the mixed glasses, fluorine was present as mixed sites: F–Ca(3), F–Ca(2)Sr, F–CaSr(2), and F–Sr(3). Ca had a higher tendency to occupy the F–M(3) sites than Sr, which may reflect the higher charge to size ratio of Ca²⁺ relative to Sr²⁺ and its greater affinity for F⁻ ions.

An in vivo animal study by Sabareeswaran et al. (2013) on a glass-ceramic that surface crystallized to a strontium FAp phase and a Sr-celsian phase (feldspar) as identified by XRD was found to be highly ossteoconductive and biocompatible. X-ray microtomography (XMT) analyses of the synthetic–organic interfaces showed highly mineralized newly formed bone adjacent to the synthetic implant surface.

It is known that, at low concentrations, strontium can promote mineralization of bone (Verberckmoes et al., 2003). It is also known that strontium-containing bio-glasses show increased osteoblast proliferation and alkaline phosphatase activity (Gentleman et al., 2010). Therefore, it would be desirable to establish ion release (particularly Sr²⁺) profile for the strontium FAp GC systems developed by Hill et al. (2004).

Recently, Chen et al. (2014a,b) developed novel alkali-free FAp and chlorapatite (ClAp) glass-ceramics from bioactive glasses. Chen et al. (2014a) produced a series of bioactive glasses of varying metal fluoride content and found that bioactive glasses with high fluoride content crystallized to FAp and crystalline CaF₂ and SrF₂ on quenching. This demonstrates that metal fluoride content can determine crystallization window (T_x) between the glass transition (T_g) and the crystallization onset (T_onset). It is desirable to obtain initially amorphous base glass so that the crystallization of, say a FAp phase, can be controlled. Chen et al. (2014a) also found that the amorphous base glass powder without any fluoride content crystallized to a wollastonite phase through a surface mechanism. Additionally, all fluoride-containing glasses bulk crystallized to FAp via a homogenous nucleation mechanism.

Chen et al. (2014c) were able to develop novel bioactive chloroapatite, ClAp glass-ceramics in the system of SiO₂–P₂O₅–CaO–CaCl₂. It is known that ClAp completely converts to HAp in the presence of water (Elliott and Young, 1967), hence making ClAp glass-ceramics attractive for both, medical, and dental applications. Chen et al. (2014c) emphasize that ClAp is less stable than FAp. This is attributed to the chloride ion being larger than hydroxyl or fluoride ion. The fluoride ion is small enough to fit in the center of the Ca(II) triangle in the FAp lattice, whereas larger ions such as hydroxyl and chloride do not fit in the center of the Ca(II) triangle but are rather displaced above the plane of the Ca(II) triangle. The chloride ion is larger than hydroxyl ion, and therefore, it is displaced further away from the Ca(II) triangle. This intrinsic apatite lattice instability brought about by the chloride ion allows the rapid exchange for a smaller ion, such as a hydroxyl ion. On increasing the fluoride or chloride content in the bioactive glass systems, Chen et al. (2014b,c) found that there was an increasing tendency of the glasses to crystallize. The halogen-free glass surface crystallized to a pseudowollastonite (α-CaSiO₃) and an apatite, presumably an oxyapatite. Pseudowollastonite induces apatite formation in SBF (Siriphannon et al., 2000) and is highly resorbable in vivo (De Aza et al., 2000). ClAp phases in the Chen et al. (2014b) crystallized via the homogenous route, which from a material processing point of view is highly desirable. FAp is largely insoluble in vivo so ClAp GCs offer the potential for producing highly resorbable GC implants. Nonetheless, there is a need for further characterization of these ClAp GCs, including characterization of the in vivo activity and the relationships between the composition, heat-treatment, microstructure, and mechanical properties.

Höland et al. (1994) developed FAp-leucite (KAlSi₂O₆) (A–L) glass-ceramics in the SiO₂–Al₂O₃–Na₂O–K₂O–P₂O₅–F system, for potential application in restorative dentistry. Research led by Höland led to the development of the commercial A–L glass-ceramic IPS d.SIGN (Ivoclar Vivadent, Liechtenstein). Needle-like apatite phase in the A–L system crystallized through a homogenous crystallization mechanism, which is a likely indication of prior LLPS. Additionally, upon DSC analyses of the A–L system, Höland et al. (2000) observed a unique phenomenon; the A–L system exhibits two endothermic reactions, first at 565°C and second endothermic reaction at 634°C. Previously, the two endothermic reactions were assigned to a phase transformation into two amorphous phases; one glassy, silica-rich phase and a droplet-like phase rich in Ca and P elements. Höland et al. (2000) suggests that the second endothermic reaction at 634°C is a transformation to a crystal phase, namely a sodium–calcium orthophosphate (NaCaPO₄), which was confirmed by XRD analysis. Furthermore, Höland et al. (2000) observed another interesting phenomena; at higher temperatures (640°C), NaCaPO₄ crystals dissolved and recrystallized to a new crystal phase. However, this new crystal phase could not be matched to any known phases when checked against the International Center of Diffraction Data (ICDD) database. Additionally, Höland et al. (2000) found that once the new phase is formed, apatite crystallization proceeds at 700°C. Heat treatment of the A–L system at 700°C for 8 h does not result in a material with needle-like microstructure; it requires additional heat treatment at 1,050°C for 2 h for the development of needle-like FAp crystals (Höland et al., 2000). This demonstrates that the thermal treatment of the base glass can strongly influence both, the appearance and other properties of the GC material.

Höland et al. (2000) suggest that the morphology of the needle-like apatite is comparable to that of the apatite in natural teeth; therefore, such needle-like FAp morphology imparts the GC restoration with exceptional esthetics. The A–L glass-ceramic exhibits good chemical durability, solubility of the GC being only at 60–70 µg/cm². Translucency of the A–L system varies between 5.8 and 10.4%, depending on the mode of processing. The A–L shows thermal expansion of 14.1–14.8 × 10⁻⁶ K⁻¹, which is close to that of titanium; therefore, the A–L glass-ceramic is highly suitable for direct sintering on metal abutments. Since 1998, there have been more than 60 million dental restorations performed using commercial IPS d.SIGN^® A–L glass-ceramic (Höland and Rheinberger, 2008), making it the most commercially successful apatite glass-ceramic developed to date.

Höland et al. (2015) developed radiopaque strontium FAp (Sr-FAp) containing glass-ceramics (Table 6) for dental applications where Sr–FAp phase crystallized through homogenous mechanisms. Höland et al. (2015) were also able to obtain secondary and tertiary crystal phases, in addition to Sr–FAp, including leucite (KAlSi₂O₆), rubidium leucite (RbSi₂O₆), cesium pollucite (CsAl₂S₂O₆), and sodium strontium orthophpshate (NaSrPO₄). However, both leucite phases and pollucite phases showed surface crystallization. Höland et al. (2015) found that some of the base glasses that were rapidly quenched into water (to prevent crystallization) did not avoid crystallization completely. Based on XRD analyses of all parent glasses, Höland et al. (2015) found that all glasses, except reference glass No. 5, were nanocrystalline. XRD analyses of the base glasses No. 1 and No. 2 showed the presence of nanoscale Sr₅(PO₄)₃F, whereas base glasses No. 3 and 4 contained nanocrystalline phases of NaSrPO₄. Höland et al. (2015) also report that only two of the “as-quenched” base glasses were optically clear, namely, No. 1 and 5, as opposed to base glasses No. 2, 3, and 4, which were opalescent in the visible light. Höland et al. (2015) showed that Sr–FAp glass-ceramics are radiopaque; therefore, these dental GCs may become more clinically relevant once commercialized.