J Mater Sci (2014) 49:4590–4594 DOI 10.1007/s10853-014-8159-6
Bioactivity potential of calcium alumino-silicate glasses and glass–ceramics containing nitrogen and fluorine A. R. Hanifi • C. M. Crowley • M. J. Pomeroy Stuart Hampshire
•
Received: 21 November 2013 / Accepted: 11 March 2014 / Published online: 27 March 2014 Ó Springer Science+Business Media New York 2014
Abstract Calcium alumino-silicate glasses of general composition (in eq.%) 28Ca:57Si:15Al:[100 - (x ? y)]O:x N:yF (x = 0 or 20 and y = 0, 3 or 5) and their glass–ceramic counterparts were immersed in simulated body fluid (SBF) at 37 ± 0.5 °C for 28 days to assess their potential bioactivity. The glasses showed no Ca release or surface calcium phosphate deposition due to their high network connectivities ([2.55). The glass–ceramics all showed potential bioactivity, as the SBF became enriched in Ca and calcium phosphate deposits formed on their surfaces. This was a result of Ca release from crystalline phases (predominantly wollastonite in the case of CaSiAlOF glass–ceramics and gehlenite in the case of CaSiAlONF glass–ceramics). No aluminium was leached from the glass–ceramics into the SBF, due to its pH always exceeding 7.0.
Introduction Since the initial development of BioglassÒ 45S5, numerous glasses and glass–ceramics have been identified that are capable of bonding to bone tissue but clinical application is presently restricted to BioglassÒ [1, 2], CeravitalÒ [3], apatite-wollastonite glass–ceramic (AW-GC) [4] and Bio-
A. R. Hanifi C. M. Crowley M. J. Pomeroy S. Hampshire (&) Materials and Surface Science Institute, University of Limerick, Limerick, Ireland e-mail:
[email protected] Present Address: A. R. Hanifi Department of Chemical and Materials Engineering, University of Alberta, Edmonton, AB, Canada
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veritÒ [5]. These materials are osteoconductive and bond to bone tissue following the deposition of apatite at the material–tissue interface [2, 4–6] by leaching of calcium from the surface of the glass or glass–ceramic. Formation of surface silanols occurs and they polymerise to form a silica hydrogel layer [1, 6, 7]. Trigonal siloxane rings act as nucleation sites for formation of an amorphous calcium phosphate layer, which crystallises to form a biologically active, carbonate-containing hydroxyapatite layer that is chemically similar to the mineral phase of bone [1, 2, 7]. For this to occur the glass (or glass–ceramic) must be soluble, that is, it should have low network connectivity (NC) so that ion release is facilitated. Consequently, the potential of a material to be considered as a bone substitute has frequently been correlated with its chemical composition, or in the case of a glass, NC [7]. In SiAlON glasses, the incorporation of nitrogen in place of some oxygen atoms in the alumino-silicate glass network increases NC [8, 9]. The increase in NC of CaSiAlON glasses, with increasing N content, also results in increases in mechanical properties. Addition of monovalent cations or anions, e.g. fluorine, is known to reduce NC of glasses. In CaSiAlON glasses, the addition of fluorine independently decreases NC but has no significant effect on molar volume, compactness, Young’s modulus or microhardness [10, 11]. Crystallisation of CaSiAlONF glasses [12] results in glass–ceramics with mechanical properties comparable with AW-GC or BioveritIÒ [4, 5]. Some CaSiAlON glass–ceramics have been reported to be cytocompatible in vitro [13], and addition of nitrogen to bioactive glasses has improved mechanical properties [14] but, to date, bioactivity of CaSiAlON(-F) glasses and glass–ceramics has not been examined. Thus, the aim of this paper is to report on an initial study of potential bioactivity of these materials.
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Table 1 Nitrogen (x) and fluorine (y) contents of glasses of composition (in eq.% of elements): 28Ca:57Si:15Al: [100 - (x ? y)]O:xN:yF and corresponding weight percentages of starting compounds with network connectivity (NC) values Glass Type
(eq.%) x
(wt%) y
CaO
SiO2
Al2O3
CaF2
Si3N4
NC
Oxide
0
0
41.4
45.2
13.4
0
0
2.81
Oxyfluoride
0
3
36.3
44.4
13.2
6.1
0
2.65
Oxyfluoride Oxynitride
0 20
5 0
33.1 42.9
43.9 30.4
13.1 13.9
10 0
0 12.8
2.55 3.56
Oxyfluoronitride
20
3
37.6
29.8
13.7
6.3
12.6
3.35
Oxyfluoronitride
20
5
34.2
29.5
13.5
10.4
12.4
3.20
Materials and methods The generic compositions of the glasses and glass–ceramics were (in equivalent percent, defined in refs [8, 9]): 28Ca:57Si:15Al: (100-x-y)O, xN:yF. Values of x and y are shown in Table 1 together with the weight percentages of constituent powders. Glass batches were prepared by melting, casting and annealing as previously described [10, 11]. Table 1 also shows NC = 2 ? CLD, where CLD is cross-link density calculated according to [11]. The glass–ceramics were prepared by heat treating the glasses under 0.1 MPa N2 in a horizontal tube furnace [12]. X-ray powder diffraction was performed on sub-45 lm samples (Phillips X’Pert diffractometer) using CuKa radiation ˚ ). The proportions of crystalline phases in the (k = 1.54056 A glass–ceramics were semi-quantitatively assessed by comparison of relative intensities of the strongest reflections [12]. Glass and glass–ceramic samples, 15 9 10 9 3 mm3, were immersed at 37 ± 0.5 °C for 28 days in simulated body fluid (SBF) prepared using the method of Kokubo and Takadama [15]. The concentrations of Ca and Al in the SBF prior to and after immersion for 28 days were determined by Atomic Absorption Spectroscopy (Varian Spectra AA220). The concentration of Al in the SBF as prepared was 0.34 ppm and that of Ca was 111.5 ppm. The trace of Al in SBF is due to the use of distilled water for preparation rather than ‘ultrapure’ water (as per e.g. ISO3696) and the use of reagents which were only 99.5 % purity. pH of the SBF was measured before and after immersion. Following immersion, samples were washed in deionised water before drying in a desiccator. Glass and glass–ceramic sample surfaces were gold-coated (Edwards Sputter SI50B) before examination by scanning electron microscopy (SEM, JEOL JSM-840) combined with energy dispersive spectroscopic (EDS) analysis.
Results Glasses were successfully prepared and were completely X-ray amorphous. Weight loss during glass melting was
Table 2 Phase assemblages of glass–ceramics 28Ca:57Si:15Al:[100 - (x ? y)]O:xN:yF glasses
derived
from
x (eq.%)
y (eq.%)
Phase assemblage
0
0
pseudo-wollastonite [ gehlenite anorthite
0
3
wollastonite [ anorthite [ cuspidine
0
5
cuspidine anorthite [ wollastonite
20
0
gehlenite [ [ [ quartz
20
3
gehlenite [ [ [ cuspidine
20
5
gehlenite [ [ [ fluorite [ cuspidine
Pseudo-wollastonite (CaSiO3), wollastonite (CaSiO3), gehlenite (Ca2Al2SiO7), anorthite (CaAl2Si2O8), cuspidine (Ca4Si2O7F2), quartz (SiO2), fluorite (CaF2)
less than 1 %, indicating that N or F losses, as SiO/N2 or SiF4, respectively, were insignificant. This was confirmed by EDS analysis of the glasses containing 20 eq.% N and 5 eq.% F. Heat treatment of the glasses resulted in surface nucleation and crystallisation to form the phases identified in Table 2. Residual glass was also present in all materials. Table 2 shows that the oxide glass primarily crystallised to pseudowollastonite, gehlenite and anorthite, while addition of fluorine to the oxide glass resulted in a change in crystalline phase assemblage to wollastonite, cuspidine and anorthite, with the relative proportion of wollastonite to cuspidine decreasing as the fluorine content increased. The CaSiAlON and CaSiAlONF glasses, however, primarily crystallised to gehlenite but the relative proportion of gehlenite decreased with increasing fluorine content as more cuspidine, and then cuspidine and fluorite formed as devitrification products. Visual and SEM examination of surfaces of the glasses and glass–ceramics prior to immersion in SBF showed no evidence of voids or cracks. Figure 1 shows data for Ca enrichment of the SBF due to solution of the glasses and glass–ceramics after immersion for the 28 day period. It is clear that the Ca enrichment ratio, [CaSBF]28/[CaSBF]0, is *1 for all the glasses, i.e. no enrichment of SBF by Ca occurs, whereas for the glass–ceramics, the ratios are from 1.2 to 2.65. Clearly, devitrification enhances calcium
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Fig. 1 Effect of F content on change in Ca content of SBF after immersion of glasses and glass–ceramics ([CaSBF]28) with respect to initial Ca content of SBF ([CaSBF]0)
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indicate an increase in surface coverage when 5 eq.% F substitutes for O. In contrast, substitution of 20 eq.% N for O causes separate spheroidal deposits of the Ca–P containing phase (Fig. 3c). Additional substitution of 5 eq.% F for O again increases surface coverage (Fig. 3d), and thus it is possible to modify the potential bioactivity of CaSiAlONF glass–ceramics by manipulation of composition and crystallisation. As indicated above, XRD confirmed only hydroxyapatite, and no other reflections were observed corresponding to the underlying crystalline phases of the glass–ceramics. However, as shown in Table 2, for the glass–ceramics without nitrogen, wollastonite is the major phase (except with high F content where Ca enrichment is less after immersion). For the glass–ceramics containing nitrogen, gehlenite is the major phase. Figure 4 shows that as the amount of wollastonite or gehlenite in the glass–ceramics increase, as indicated by their X-ray diffraction intensity ratios, the Ca enrichment of the SBF increases in a reasonably linear manner. In addition, potential bioactivity decreases as fluorine content of the glass–ceramics increases.
Discussion
Fig. 2 Effect of F content on change in Al content of SBF after exposure of glasses and glass–ceramics [AlSBF]28 with respect to initial Al content of SBF [AlSBF]0
release from the materials. The corresponding data for aluminium, [AlSBF]28/[AlSBF]0, is given in Fig. 2 and clearly shows that Al content of the SBF appears to decrease over the 28 day period, and thus no enrichment of SBF with Al occurred for either glasses or glass–ceramics. SEM examination of the glass surfaces after immersion in SBF showed no surface deposition of Ca or P, as might be expected if the glasses had potential bioactivity. In direct contrast, the glass–ceramics all exhibited surface coverings containing Ca/P in a ratio, analysed by EDS in the SEM, which is consistent with a carbonated hydroxyapatite (1.69). Thin film XRD confirmed that the deposits were crystalline hydroxyapatite. Figure 3 shows the effect of fluorine, nitrogen and nitrogen ? fluorine substitution in the parent glasses on surface coverage of glass–ceramics by the hydroxyapatite phase. Comparison of Fig. 3a, b
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It is clear from above that none of the CaSiAlO(N)(F) glasses readily hydrolyse to facilitate hydroxyapatite formation. This most probably reflects the high NC values for the glasses which lie in the range 2.55–3.56 (Table 1). The lack of hydrolysation, Ca release and Ca/P deposition following immersion of the glasses in SBF correlates with previous findings for glasses with high Si contents and high NC [7] where it was proposed that hydrolysis of bioactive glasses was the result of SiO4 tetrahedra being associated with a high number of non-bridging oxygens, and thus low NC values (\2.4). When the glasses are crystallised to their glass–ceramic counterparts, it can be shown that the network connectivities of the residual glasses increase further from the high values for the oxide, oxyfluoride and oxynitride glasses. For the oxyfluoronitride materials, there is a slight decrease in NC but they remain [3.0. This observation would strongly suggest that it is the crystalline phases which are hydrolysable and soluble to enable Ca leaching from the glass–ceramics followed by subsequent deposition as hydroxyapatite confirmed by XRD. There is evidence from Raman spectroscopy [16] that the structure of gehlenite glass is more ‘polymerized’ than that of crystalline gehlenite. Thus, the crystalline form would be expected to be more susceptible to hydrolysation. This conclusion is further endorsed by the fact that Ca release levels increase with increasing X-ray intensity fractions of wollastonite and gehlenite (see Fig. 4). Based on the findings and
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Fig. 3 Surface topographies of glass–ceramics showing development of Ca–P surface deposit a 0 eq.% N:0 eq.% F, b 0 eq.% N:5 eq.% F, c 20 eq.% N:0 eq.% F, d 20 eq.% N:5 eq.% F
Fig. 4 Effect of pseudo-wollastonite ? wollastonite (PW ? W) or gehlenite (G) content of glass–ceramics on Ca enrichment of SBF after 28 days
arguments presented, it may be concluded that glass– ceramics derived from these CaSiAlONF glasses have potential bioactivity.
When gehlenite acts as the crystalline source for Ca, in the case of the oxynitride and oxyfluoronitride materials, it might be expected that in addition to Ca enrichment, the SBF would show Al enrichment too. As this is not observed, it is thought that this is because pH values of the SBF following glass–ceramic exposure always exceeded 7.0. Fluoro-alumino-silicate glasses do not generally release Al unless the pH is acidic [7]. Thus, the :Al–O– Si: linkage is only hydrolysed under acidic conditions (i.e. pH B 7.0), and so Al would be expected to remain bound to the surface layers of the material, either in the high NC residual glass or within a surface hydrolysed layer, where the :Al–O–Si: linkage is maintained. The glass materials investigated here only show potential bioactivity if devitrified, and the nature and amount of crystallisation products can be controlled by fluorine and nitrogen substitution levels, as well as by heat treatment. Accordingly, the family of materials reported here can be carefully manipulated such that their phase assemblage can be altered to vary their potential bioactivity. This suggests that further evaluation of the bioactivity of the glass– ceramics using other tests would be useful.
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Conclusions For the CaSiAlONF glasses studied here, the following can be concluded: (1)
(2)
(3)
(4)
Oxide, oxyfluoride, oxynitride and oxyfluoronitride glasses in the CaSiAlO(N)(F) system do not release Ca or Al into SBF, and thus are not soluble. This is due to high glass network connectivities which exceed 2.55. Crystallised glass–ceramic derivatives facilitate calcium phosphate deposition on material surfaces due to significant calcium release into the SBF from the crystalline phase(s) formed. For nitrogen containing glass–ceramics which devitrify predominantly to gehlenite, the neutral to basic nature of the SBF during exposure precludes Al release from gehlenite. The potential bioactivity of the CaSiAlONF glasses studied here can be controlled by tailoring of composition and crystallisation.
Acknowledgements The authors wish to thank Science Foundation Ireland for financial support for this research under grant no. 04/BR/ C0163.
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