Hydrated silica layers have shown a critical role, as mentioned above, on the bioactivity or apatite formation on calcium silicate glasses and some inert materials. How about silica gel via the sol-gel route as they inevitably hold many silanol groups in the course of gelation? The silica gels from common systems consisting of tetraalkoxysilane Si(OR)₄, water, and solvent together with acid or alkali catalysts would never induce apatite nucleation in the Kokubo’s simulated body fluid (SBF). Nakanishi was the first who fabricated a single oxide gel with proven bioactivity: it was pure silica gel with mesopore and micropores inside [24, 25]. Therefore, it is worthwhile to take a deeper look into the gels. Nakanishi elegantly employed phase separation phenomenon to fabricate mesoporous silica gels [24]. The systems consist of Si(OR)₄ (R = CH₃ (TMOS) and C₂H₅ (TEOS)); water as the solvent and hydrolyzer; organic additives such as formamide, polyethylene glycol (PEG), and poly(acrylic acid) (PAA); and nitric acid as the catalyst. As hydrolysis of Si–OR groups (Eq. 5.7a) and condensation polymerization of the resultant silanol groups (Si–OH) (Eq. 5.7b) proceed, silicate oligomers are produced, together with water and ROH as the by-products that are involved into the solvent.

Miscibility among those silicate oligomers, solvent (H₂O + ROH (= a reaction product), and organic additives controls the phase separation into two phases: silicate-rich and solvent-additive phases. As those reactions proceed, the composition of the solvent continuously changes, and the miscibility is also changed, accordingly. In conventional cases, the phase separation is caused by the decrease in miscibility due to reduced temperature, and it is the physical cooling. In the Nakanishi’s case [24], the separation takes place under chemical cooling where the interactions among the species change with the growth in the molecular size of the silica oligomers and the solvent composition. Nakanishi interpreted those changes in a thermodynamic way. Regardless of the cooling principle, the change in the free energy of mixing ∆G (Eq. 5.8) controls the thermodynamic stability of the system.

Here, ϕ and P stand for the volume fraction and degree of polymerization of the separating phases 1 and 2, respectively. Parameter χ represents the interactions between those phases and contributes either to destabilize or to stabilize the system. The first two terms represent the entropy contribution and the third the enthalpy one. Equation 5.8 indicates that the degree of polymerization of either phase destabilizes the system. Thus, proper selection of the starting materials in the precursor solution leads to appropriate phase morphology as well as the volume fraction of the silicate phase. When the system is metastable, ∂ΔG/∂ϕ ² > 0, certain level of density fluctuation is necessary prior to the separation, as a classic model of nucleation and growth postulates, where the droplet-type separation takes place to yield particle aggregates or gels with isolated pores. In contrast, when the system becomes unstable on the polymerization, ∂ΔG/∂ϕ ² < 0, any small fluctuation in density triggers the phase separation known as spinodal decomposition, leading to entangled microstructures. As the rate of polymerization is highly dependent on the temperature, and affinity among the atomic groups and molecular in the systems, as well, the temperature control and the period of reaction are critical to obtain the gels with desired microstructures.

The hydrated silica layers on glass or glass-ceramics, aforementioned, have presumably the same as or similar to the Nakanishi’s mesoporous silica gel in atomic structure. Cho et al. examined the apatite formation on several mesoporous silica gels derived under different additives (polyethylene glycol: PEG and polyacrylic acid: PAA) [25] as well as the formation on those heated up to 1,000 °C [21]. Figure 5.3 shows scanning electron micrographs of the fracture surfaces of the PEG-derived mesoporous silica gels, calcined at 400 and 1,000 °C for 2 h. Entangled pores used to be filled with the solvent and PEG [25]: both have a very similar macroporous microstructure. Strangely enough, the gel from the PEG-containing precursor solution was able to induce the apatite deposition in SBF within 14 days but not the gel from the PAA-containing solution [25]. The PAA gel had a finer entangled microstructure than the PEG gel, while the former exhibited a type I N2 adsorption–desorption isotherms and the latter showed a type IV profile. No other obvious difference was noticed. Among the structural defects, those known as D1 and D2 defects, i.e., four- and three-Si–O-membered cyclic groups, respectively, were detected in both PAA- and PEG-involved gels by Raman spectroscopy, whereas the former defects remained even after soaking both gels in SBF but the latter disappeared on soaking. It was also found that dissolution of Si(IV) from both gels was at the same level, as well as both had almost the same amount of Si–OH groups regardless of being soaked in SBF or not. At this moment, therefore, no definite reasoning or correlation to the structure has been given for the apatite deposition on the mesoporous silica gels depending on the polymers in the precursor solutions. In contrast, the calcining temperature effects superficially seem more straightforward: according to the analyses by Fourier transform infrared spectroscopy and scanning electron microscopy, the gels calcined at 900 °C or above would not deposit any apatite in SBF within 14 days, but the gels calcined at 400 °C or above up to 800 °C deposited apatite. Since dissolution of Si(IV) in 14 days from the gels calcined above 800 °C was half as much as that from the gels calcined at 400–800 °C, such release of Si(IV) species is common among the sol-gel-derived silica and silicate systems. When they are dead-burned, the surface silanol groups are fully condensed. Again, it is the hydrated silica layer that is active to induce apatite deposition in SBF or plasma.

A simple question of knowing the mesoporous silica gels are bioactive is whether those gels are structurally very similar to the hydrated silica layers produced on those bioactive silicate glass or glass-ceramics. No known attempts were made to compare both in terms of their structures, partly because the silica layers are mostly too thin to deserve analyses and partly because the bioactivity of the silica layers is dependent on the calcium ions in the glasses and glass-ceramics.

The mother glass of Cerabone A-W® involves F as CaF₂ and is subjected to crystallization or devitrification procedure before solid and pore-free bone substitutes are produced. Calliano and Lopez conducted a similar study on crystallization and sintering of F-containing lime silicophosphate glasses [103]. While F can be involved in silicate glasses in two ways, the free ions (ionic F) are coordinated by the corresponding modifier ions, and the other (covalent F) is covalently bonded to Si and P (Si–F or P–F) [104, 109]. The former ones were predominant in the calcium silicate glasses [104]. When F remains ionic in glass, the network connectivity is determined only by the ratio CaO/SiO₂. That is, the connectivity is the same as that for the F-free glass. In contrast, the formation of Si–F bonds yields a degraded glass network, comparing with the F-free glass. This may lead to better susceptibility of hydrolysis when in contact with SBF or plasma. Moreover, the ionic F species, or the free fluoride ions, will be incorporated in the apatite lattice to form fluorapatite, which is chemically more stable than ordinary hydroxyapatite. An earlier work by Bauer et al. reported that an increase in the F content in the glass promoted the CaF₂ precipitation but suppressed the precipitation of apatite, in comparison to the F-free glass [105]. Bauer et al. [110] suggested that the phosphate content determined the CaF₂ and CaCO₃ (calcite) formations in the F-containing and F-free glasses, respectively. Lynch et al., in contrast, reported a seemingly contradicting result that a series of fluoride-containing glasses were so active to deposit apatite within 6 h. In consequence, it is surely accepted that we have varied factors to control, such as glass compositions (content of F, P₂O₅, SiO₂, CaO, or other components), the period of soaking the samples in solution, sample size, surface area to solution volume ratio, and solution itself (SBF or other buffer solution), to name a few. Therefore, as Coccihi et al. postulated [108], multivariate data analysis is necessary when one should elucidate reliable conclusions out of individual experiments.

Bioglass® 45S5 glass is liable to devitrification on shaping. Moreover, according to a differential thermal analysis to derive Johnson-Mehl-Avrami exponent conducted by Massera et al. [111], 45S5 glass strangely changed the crystallization mode from surface crystallization to bulk one when the grain size increases. Some efforts thus have been paid to fabricate “longer” glass, that is, glasses with better workability or a wider working temperature range [112]. Glass-ceramics involving magnetic phases can be applied to hyperthermia, where the heat evolved from the magnetic phases susceptive to microwave irradiation. A few groups attempted to provide such magnetic glass-ceramics with bioactivity [82–84]. For example, Ebisawa et al. [82] were the first who developed not only magnetic but also bioactive glass-ceramics involving ferrites. Here, the bioactivity was provided by the wallastonite phase precipitated in the course of crystallization. For the hyperthermia cancer therapy, various types of magnetic particles were proposed like hollow and mesoporous silica spheres [113, 114], cobalt ferrite (core)–apatite (shell), and magnetite–apatite composite nanoparticles [115, 116], while Maehara et al. [117] preferred magnesium ferrite (MgFe₂O₄) to the other ferrites of Cu, Fe (magnetite), Mn, Ni, Sr, or Co. Incidentally, Mother Nature gives us bacterial magnetic particles (BMPs) [118, 119], which are mostly 50–100 nm in size and nowadays are considered for gene therapy or other clinical treatments [120]. It is a key disadvantage in the hyperthermia with magnetized materials that microwave can only penetrate into shallow regions under the skin. In order to overcome this issue, some applicators like antenna or needles have been employed to raise the temperature at deeper region, but a good care is necessary for avoiding burning the skin.

In 2001, Xynos et al. reported that the ionic species dissolved from 45S5 Bioglass® should stimulate expression of the genes relevant to human osteoblasts [121]. Then, a movement was emerged [88, 122, 123] to develop third-generation biomaterials that should stimulate the body to heal itself, i.e., cell- and gene-activating materials, in contrast to the previous generation materials that would simply be active only to replace the damaged tissues, to fill up the defects, or to protect from the invasion of cells or tissues harmful for swift recovery. A tendency is emerged to drive researches to develop biodegradable materials from the molecular biology aspect. The materials derived from the sol-gel procedure are advantageous in degradation compared to those via the high-temperature processes such as melt-quenching, sintering, and dead-calcining, since they commonly have many –OH groups due to incomplete calcination. Table 5.5 lists some examples of sol-gel-derived glasses. Unfortunately, against expectations, the sol-gel amorphous materials, like 58S [87], had not yielded statistically significant results in terms of osteoblastic marker gene expression. Note here that although most people praise apparent homogeneity of the sol-gel-derived amorphous materials, the truth is that the atomic distribution in the sol-gel-derived bodies is heterogeneous as far as the multicomponent ones are concerned. This would be justified by the consideration that on the derivation of the sol-gel materials, the alkoxides and inorganic salts are different in the rate of hydrolysis or condensation. In the classic sol-gel science, a few techniques for homogeneous gels were attempted such as prehydrolysis of the components, use of chelating agents, and application of ultrasound [124, 125].

Classic model of bioactivity is based on the principle that the material surface should give a metastable hydrated gel layer to serve nucleation sites of apatite deposition [5, 6, 126]. As has been discussed in Sect. 5.2.2, the Missouri group [63, 64, 123, 127, 128] has proposed a method to convert glass into apatite, i.e., to obtain apatite hollow spheres and solid spheres, dependent on the glass composition, at the expense of glass itself. When dissolvable glass containing calcium ions is in contact with a phosphate solution of pH > 8, the chemical state of the glass–solution interface described in Fig. 5.4 will be established. At the beginning, the hydrolysis proceeds at the glass surface to leave a stable hydrated gel. When the surface layer is chemically unstable or dissolved, the calcium ions in the glass are supplied almost freely and ready to precipitate apatite particles. The borate glass derived by substituting B₂O₃ for SiO₂ in 45S5 glass is highly susceptible to the dissolution attack by the alkaline medium, according to the process similar to that described by Eqs. 5.3a and 5.3b, to form a less dense hydrated layer, in which the calcium ions can migrate in all directions to meet the phosphate ones from the solution. That is, calcium borate glass particles are fully converted to polycrystalline hydroxyapatite with random orientation [64] while the borosilicate glass particles are partially converted to yield apatite core–shell glass particles or hollow apatite particles depending on the compositions [123, 127, 128].

In contrast, when a piece of flat calcium silicate glass with moderate dissolvability is in contact with similar phosphate solution, the metastable hydrated layer is sustained as the surface hydrolysis and dissolution of cations proceed. In this situation, the Ca ion supply in the direction parallel to the surface is impractical, being different from the particle case. As a result, established is the Ca²⁺ ion concentration gradient within the interface region of the glass, the normal of which is perpendicular to the interface. If a large number of reaction sites are available to nucleate the crystallites, the nuclei, earlier produced or comparatively larger ones, will have a freedom to grow. Then nanometer-scale apatite rods will be formed in array perpendicular to the mother glass surface. Indeed, Hayakawa et al. [129] found for a series of soda-lime glasses with lower silica contents, xNa₂O–(45−x)CaO–(54 or 55)SiO₂, x = 0, 10, 20 (mol%). Among a few glass compositions, a glass-coded G36 (containing 36 % CaO) yielded the most elegant self-assembled apatite nano-rod arrays that grew with an alignment in the direction of the c-axis of the apatite lattice when it was kept in contact with 0.01 M Na₂HPO₄ at 80 °C for 1 h. The degree of alignment is as high as 0.85, in terms of the measure proposed by Lotgering [130], which was almost equivalent to that of the dental enamel. They could not, unfortunately, detect the hydrated silica layer at the glass–apatite interface. Yet, longer soaking in the phosphate solution fully converted the glass substrates to a block of apatite nano-rod array. Interestingly, when soaked in SBF (Kokubo’s solution), glass G36 gave hemispherical agglomerates of apatite crystallites for which the c-plane grew to give the same petallike morphology as the conventional bioactive silicate glasses or glass-ceramics [13, 18–22].

In addition to the silicate, borate, and borosilicate glasses and glass-ceramics, some phosphate glass-ceramics in the systems CaO–P₂O₅