A new model formulation of the SiO2–Al2O3–B2O3–MgO–CaO–Na2O–F glass-ceramics

ARTICLE IN PRESS

Biomaterials 26 (2005) 2255–2264 www.elsevier.com/locate/biomaterials

A new model formulation of the SiO2–Al2O3–B2O3–MgO–CaO–Na2O–F glass-ceramics Simeon Agathopoulosa,, Dilshat U. Tulyaganova, Patricia Vale´riob, Jose´ M.F. Ferreiraa a

Department of Ceramics and Glass Engineering, University of Aveiro, CICECO, 3810-193 Aveiro, Portugal Department of Biophysics and Physiology, Federal University of Minas Gerais, ICB-UFMG, 31270-901 Belo Horizonte, Brazil

b

Received 16 December 2003; accepted 12 July 2004 Available online 15 September 2004

Abstract Mono-phase glass-ceramics of akermanite were successfully produced from a Ca-mica and wollastonite via low-temperature sintering and crystallization. Doping with P2O5 considerably improves sintering behaviour since P2O5 increases the stability of glass against crystallization at the temperature of sintering onset. The resulting glass-ceramics feature good in vitro acceptance from osteoblasts, and moderate bioactivity due to the enrichment of the glassy phase with Ca and Si. The good quality of the white colour at the surface and throughout the bulk, the matching of microhardness with tooth enamel, and the possibility to coat other biomaterials such as ZrO2, Ti or hydroxyapatite make these materials promising for medical applications. r 2004 Elsevier Ltd. All rights reserved. Keywords: Akermanite; Glass; Glass-ceramics; Properties; Biocompatibility; In vitro

1. Introduction Glass-ceramics have been broadly used as dental restorative and osteoconductive materials due to their biocompatibility and potential bioactivity [1–4]. For instance, apatite-wollastonite (A/W) has been used in orthopaedic prostheses [5–7], while glass-ceramics based on potassium mica, such as Dicors, have been applied in dental restorations due to their good physical, chemical and mechanical properties [8,9]. In this work, we present a new model system, which comprises components of boron-containing calcium mica (CaMg3Si2Al0.5B1.5O10F2) and wollastonite (CaSiO3). The most important challenge in this work was to determine the optimum ratio between mica and wollastonite, which would eventually result in a single crystalline phase. In such systems, several phases of complex structure are usually formed. However, a multi-phase Corresponding author. Tel.: +351-234-370242; fax: +351-234425300. E-mail address: [email protected] (S. Agathopoulos).

0142-9612/$ - see front matter r 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.biomaterials.2004.07.030

material is generally not desirable because the likely mismatch of thermal expansion coefficients between the different crystalline phases can cause weakening of material. To our knowledge, there are no similar studies reporting single-phase glass-ceramics tested or used in biomedical applications. In this investigation, akermanite (Ca2MgSi2O7) was the optimum demanded phase. Akermanite and ghelenite are the end members of melilite, which belongs to the group of sorosilicates [10], having a general formula of X2YZ2O7, where X is Ca or Na, Y is Al or Mg, and Z is Si or Si and Al. X is a large 8-coordinated site, while Y and Z are tetrahedral ones. Akermanite and ghelenite form solid solutions with a general formula of 2CaO
(1x) MgO
xAl2O3
(2x)SiO2. Melilites usually occur in igneous and metamorphic rocks and meteorites, but also in the slag of blast furnaces [10,11]. These facts unequivocally evidence their high stability. Akermanite exhibits relatively high refractory properties, melting point at 1454 1C, density of 2.944 g/cm3, and hardness of 56 in Moose scale.

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The mechanism, whereby akermanite would be produced from mica and wollastonite, is assumed to be as follows. Micas consist of two sheets made of tetrahedral units, sandwiching one sheet made of octahedral structural units. The tetrahedral structural units of SiO4 exhibit high structural stability, due to 4 the four hybridized sp3 orbitals of Si4+, which can give the same configuration of these structural units in both the glass melt and the crystals of the glass-ceramic. In the octahedral layer, where there is coexistence of 4s and 2p bonds, the weaker s-bonds are primarily ruptured over heating, facilitating diffusion of cations to escape the octahedral layer [12]. Hence, the Mg2+ ions, accommodated in the octahedral units of Ca-mica, diffuse easily towards the units of wollastonite during heating according to the following simplified chemical equation: 3MgO þ 7ðCaO
SiO2 Þ ! 3ð2CaO
MgO
2SiO2 Þ þ CaO
SiO2 :

(1)

In this work, to synthesize akermanite, the optimal molar ratio between mica and wollastonite was taken as 1:7. Thus, the basic formula, hereafter denoted as MW, was (3MgO
0.25Al2O3
0.75B2O3
2SiO2
CaF2)
7(CaO
SiO2)
Na2O. The ratio 1:7 was selected because it seems that it better approaches the structural and bioactivity requirements of akermanite glass-ceramics. In particular, apart the stoichiometry of Eq. (1), i.e. 1:6, the excess of wollastonite aimed to react with the other components of mica and result in formation of glassy phase. The presence of B and F in the glassy phase was expected to favour densification of the resulting glassceramics, while Ca and Si would potentially enhance bioactivity [13]. A modified composition, denoted as MWP, was derived from MW by doping with P2O5, having a formula of (3MgO
0.25Al2O3
0.75B2O3
2SiO2
CaF2)
7(CaO
SiO2)
Na2O
0.3P2O5.

2. Materials and experimental procedure 2.1. Production of glasses and glass-ceramics and characterization techniques Powders of reagent grade of SiO2, Al2O3, CaCO3, H3BO3, 4MgCO3
Mg(OH)2
5H2O, Na2CO3, CaF2 and NH4H2PO4 were used. The chemical compositions of MW and MWP glasses are presented in Table 1. Homogeneous mixtures of 100 g batches obtained by ball milling were preheated at 1000 1C for 1 h for decarbonization and then melted in Pt crucibles at 1400 1C for 1 h, in air. Glass frit was obtained by quenching of the melt in cold water. The frit was dried and then milled. The resulting powders had specific

Table 1 Chemical compositions of the investigated glasses Glass

SiO2

Al2O3 B2O3 CaO

MW (wt%) (mol%) MWP (wt%) (mol%)

42.51 40.91 41.13 40.36

2.00 1.14 1.94 1.12

4.10 3.41 3.97 3.46

MgO P2O5 Na2O CaF2

30.86 9.51 — 31.82 13.64 — 29.86 9.20 3.24 31.40 13.45 1.35

4.87 4.54 4.72 4.44

6.14 4.54 5.94 4.44

surface area of 0.719 m2/g for MW and 0.354 m2/g for MWP, measured by BET technique (Micromeritics, Gemini II 2370, USA). Glasses in bulk form were also obtained by casting, followed by appropriate annealing to reduce internal stresses. Powder of glass frit was granulated by mixing in a 5 vol% polyvinyl alcohol solution (PVA) (wt% proportion frit/PVA=97.5/2.5). Parallelepiped bars (4  5  50 mm3) were prepared by uniaxial pressing (200 MPa). After debinding at 450 1C for 2 h, the bars of glass-powder compacts were sintered at several temperatures between 650 and 920 1C for 1 h at a slow heating rate (2–31/min) to avoid deformation. Production of coatings (thickness 100 mm) was also attempted as follows. Glass frit, granulated with PVA, was applied on pellets of several polished substrates (i.e. zirconia, Ti and hydroxyapatite) and crystallized similarly to the glass compacts at 800 1C for 1 h. The resulting materials were characterized by the following techniques: Differential thermal analysis (DTA, Labsys Setaram TG-DTA/DSC, France, heating rate 51/min, in air). Dilatometry (Bahr Thermo Analyse DIL 801 L, Germany, heating rate 31/min). X-ray diffraction (XRD, Rigaku Geigerflex D/Mac, C Series, Cu Ka radiation, Japan). Microstructure observations by scanning electron microscopy (SEM, Hitachi S-4100, Japan, 25 kV acceleration voltage) under secondary electron and back scattering modes, equipped with energy dispersive spectroscopy apparatus (EDS, minimum spot diameter 1 mm). The Archimedes method was used to measure the apparent density of the glass and glass-ceramic blocks (immersion in ethylenoglycol). Vickers microhardness was estimated from ten indentations for each sample (Shimadzu microhardness tester type M, Japan, load of 9.8 N). Weibull statistics were employed to evaluate the mechanical reliability of materials (i.e. three-point flexural strength of parallelepiped bars, 3  4  40 mm3, Shimadzu Autograph AG 25 TA, 0.5 mm/min displacement). Water absorption was measured according to the ISO-standard 10545-3, 1995 (i.e. weight gain of dried bulk samples after immersion into boiling water for 2 h, cooling for 3 h and sweeping of their surface with a wet towel). The linear shrinkage during sintering was also calculated from the dimensions of the green and the resulting sintered samples.

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2.2. Immersion in SBF The dissolution rates and the possible bioactivity of the glasses and the glass-ceramics were investigated by immersion of either powders (pulverized from sintered bulk samples) or bulk materials in simulated body fluid (SBF), at 37 1C. SBF solution has an ionic concentration of Na+ 142.0, K+ 5.0, Ca2+ 2.5, Mg2+ 1.5, Cl 147.8, 2 2 HCO 3 4.2, HPO4 1.0, SO4 0.5 (in mM), buffered at pH=7.25 by tris-hydroxymethyl-amminomethane (Tris, 50 mM) and hydrochloric acid [14]. Sampling took place at 1, 3, 6, 10, 24 (1 d), 72 (3 d), 168 (7 d), 336 (14 d), and 744 (31 d) h (d: days). The obtained results were independent, which means that each sample was individually treated without interfering with others. An amount of glass powder corresponding to 1 m2 surface area was immersed in 10 ml of filtered (through sterilized filters, cameo 25 AS-MSI, pore size 0.22 mm) SBF solution and immediately sealed into sterilized plastic flasks, which were stored afterwards at 37 1C (70.11). After each experiment, the powder was separated from the liquids by filtering. Similar processing was followed for the bulk materials. In this case, the response of material in SBF was tested at both fracture and polished surfaces (mirror finishing, i.e. polishing with 1 mm diamond paste, and ultrasonic agitation). The analysis of the liquids comprised pH measurements and determination of the concentration of Ca2+, Mg2+, B3+, P5+, Si4+, Na+, and Al3+ by inductively coupled plasma-optical emission spectroscopy (ICPOES, Jobin Yvon, JY 70 plus, France). The dried samples were examined by SEM/EDS. 2.3. In vitro tests with osteoblasts For the in vitro tests with osteoblasts the following materials were used. Penicillin, streptomycin, fetal bovine serum, Dulbecco’s phosphate buffered saline, trypsin-EDTA, [3(4,5 dimethylthiazol-2yl) 2,5 diphenyltetrazoliumbromide] MTT, BCIP-NBT kit: Gibco (Burlington, Ont., Canada). Crude bacterial collagenase: Boehringer (Biberach, Germany). RPMI Cell culture medium: Sigma (St. Louis, USA), SIRCOL kit: Biocolor (Newtonabbey, N. Ireland) T25 culture flasks and 24 well plates: Nunc products (Naperville, USA). Osteoblasts were isolated from the calvaria of 1–5 days old neonatal Wistar rats [15]. The calvaria were dissected and freed from soft tissue, cut into small pieces and rinsed in phosphate-buffered saline without calcium and magnesium. The calvaria pieces were incubated with 1% trypsin-EDTA for 5 min, followed by four sequential digestions with 2% collagenase at 37 1C for 45 min each. The supernatant of the first collagenase digestion, which contains a high proportion of periosteal fibroblasts, were discarded. The other digestions produced a suspension of cells with high proportion of osteoblasts.

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After centrifugation at 1000 g for 5 min, each pellet was re-suspended in 5 ml of RPMI medium supplemented with 10% FBS, 1% antibiotic-antimycotic. The cells were seeded into 25 ml tissue culture flasks, and led to grow in a controlled 5% CO2, 95% humidified incubator at 37 1C. After confluence, the cells were used for experiments on passage 2. To test the stimulation of osteoblasts with a medium containing glass powder, osteoblasts were platted 1  105 cell density, in 24 well plates. After 2 h, the medium was changed to a medium containing glass powder (i.e. MW or MWP; grain size of powdero35 mm). The concentration used was 0.005 g powder/ 1 ml of culture medium. After 72 h of incubation, osteoblasts’ morphology, viability/proliferation and secretion capability (i.e. collagen, alkaline phosphatase) were tested. The cellular proliferation and viability tests were done as follows: After 72 h of incubation in the presence of each glass powder, osteoblasts’ viability was evaluated by MTT assay, based on the reduction of tetrazolium salt to formazan crystals by dehydrogenase present in living cells mitochondria. The formazan salt formation is directly related to the amount of dehydrogenase, providing an indirect measurement of cell proliferation [16]. In total, 60 ml of MTT (5 mg/ml) were added to each well. Two hours later, the cell morphology was analysed by inverted optical microscopy and formazan salts were solubilized with SDS 10% HCl. After incubation for 18 h, optical density was measured at 595 nm [17]. The alkaline phosphatase production was evaluated by BCIP-NBT assay. This assay is based on a chromagenic reaction initiated by the cleavage of the phosphate group of BCIP by alkaline phosphatase present in the cells. This reaction produces a proton, which reduces NBT to an insoluble purple precipitate. In brief, the supernatant of each well was removed and the cell layer was rinsed twice with PBS. Then, 200 ml of BCIP-NBT solution, prepared according to manufacturer protocol, were added to each well. After 2 h of incubation, the cells were observed by optical microscopy and the insoluble purple precipitates were solubilized with 210 ml of SDS 10% HCl and incubated for 18 h. The optical density was measured at 595 nm [18]. The osteoblast collagen production was analysed by SIRCOL assay in the cultures’ supernatant, following the manufacturer instructions. This method is based on the selective binding property of the syrius-red dye to the [Gly–X–Y] tri-peptide end sequence of mammalian collagen [19]. The collagen present in the supernatant, precipitated by syrius-red, was solubilized and measured by an optical density analysis at 595 nm. The collagen concentration was calculated basing on a linear regression from previously known concentrations of type I collagen and their optical density measurement.

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The results are presented as Means7SD (the numbers of the experiments are mentioned in the legend of the figure). The statistical significance was measured by ANOVA and Bonferroni’s post-test.

3. Results and discussion 3.1. Properties The thermal analysis of the glasses, summarized in Fig. 1, provided the necessary information for setting up the optimum crystallization schedule. For both MW and MWP glasses, dilatometry measurements showed that the transition temperature (Tg) was at 625 1C and the softening point at 670 1C. Tg was detected by DTA at slightly higher temperature (640 1C). The DTA showed a single exothermic peak for crystallization with maximum at 780 1C for MW and 797 1C for MWP. The decomposition of PVA was also detected at 400 1C (endothermic peak). Accordingly, after debinding (450 1C, 2 h), sintering of bars of glass-powder compacts was carried out starting at 650 1C, i.e. slightly lower than the dilatometric softening point. Then, sintering of these samples was conducted at seven different temperatures, i.e. 700, 750, 800, 825, 850, 900 and 920 1C. The influence of sintering temperature on density and linear shrinkage as well as the mechanical properties of the resulting glass-ceramics are presented Table 2 and Fig. 2, respectively. Several results indicate that the modified MWP composition features superior sintering behaviour than MW. In particular, sintering starts at

1.0 625

MW

0.5

MWP 797

∆l/lo (%)

670

lower temperatures. Densification is always higher and safely occurs until 800 1C, reaching a maximum of density (2.89 g/cm3) and shrinkage (13.8%) at 750 1C. Linear shrinkage values were remarkably constant over the investigated temperature range. At 650 1C, powdered frit of MWP achieved higher density than that of bulk glasses, whose density was measured as 2.80 g/cm3 for both MW and MWP. The dramatic decrease of mechanical reliability of MW glass-ceramics sintered at temperatures X800 1C is characteristically depicted in Fig. 2a, which is in accordance with the values of Table

Table 2 Density and linear shrinkage of compacts made of glass powders and sintered at different temperatures Sample

Sintering temperature (1C) 650

MW MWP

2.51 2.83

MW MWP

4.4 13.6

MWP

825

Density (g/cm ) 2.82 2.74 2.77 2.46 2.88 2.89 2.86 2.78 Linear shrinkage (%) 12.7 12.2 12.6 13.6 13.8 13.6

0

200

400

600

2.34 2.67

2.13 2.50 10.8 13.5

2.04 2.09 10.0 13.3

-2

4.0

4.2

y=16.105x-74.039 R2=0.9892

4.4

4.6

4.8

ln (T)

(a)

2

700ºC 750ºC 800ºC

0

y=7.3746x-33.124 2 R =0.9885

-2

y=26.679x-125.94 2 R =0.9779

-4

641

3.9

-3 0

920

y=6.1094x-28.158 R2=0.9892

y=8.0658x- 38.338 2 R =0.9884

396

900

700ºC 750ºC 800ºC

-1 MW

850

-4

ln{ln[1/(1-Pi)]}

Heat flow (mV)

1

800

y=15.347x-69.953 R2=0.9903

0.0

Exo Endo

750

2

780

3

700

3

ln{ln[1/(1-Pi)]}

2258

800

1000

Temperature (ºC) Fig. 1. Thermal analysis (dilatometry top curves and DTA bottom curves) of MW and MWP glasses.

(b)

4.1

4.3

4.5

4.7

4.9

ln (T)

Fig. 2. Weibull statistics of flexural strength (3 point bending) of glassceramics obtained from (a) MW and (b) MWP glasses at different crystallization temperatures (T: bending strength, P: fracture probability).

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2. On the other hand, MWP glass-ceramics sintered at 800 1C anticipate the highest mechanical reliability (Fig. 2b). From these curves, the maximum average flexural strength is calculated at 95 MPa for MW sintered at 750 1C and 110 MPa for MWP sintered at 800 1C. Consequently, the sintering temperature ranges between 700 and 800 1C for the MW composition, having a span of about 1001. The modified MWP composition features a broader sintering span of about 1501, since remarkable consistency in the values of the properties were recorded for the materials sintered between 650 and 800 1C. For glass-ceramics sintered at 800 1C (1 h) both materials have homogeneous white colour throughout the entire bulk and the surface. Vickers microhardness was 4.01 (70.02) GPa for MW and 4.43 (70.66) GPa for MWP. Water absorption was 0.26% for MW and 0.16% for MWP. The linear thermal expansion coefficient (CTE) calculated from the expansion curves (not shown) between 70 and 500 1C was 9.4  106 K1 for MW and 9.9  106 K1 for MWP. Fig. 3 shows that at 650 1C both materials are still amorphous. Crystallization occurs at X700 1C, resulting in formation of a single phase, akerminite. A singlephase formation was correctly anticipated from the single exothermic peak of DTA (Fig. 1). It should be mentioned that in this type of systems, heating rate often affects the temperature of crystallization onset (i.e. lower heating rate favours evolution of crystalline phases at lower temperatures) [20]. Comparison of the intensity of the XRD peaks (Fig. 3) indicates that MWP generally featured lower crystallinity than MW. This finding agrees fairly well with the microstructure of glass-ceramics sintered at 800 1C, for 1 h (Fig. 4). The dense and highly crystallized structure of MW glass-ceramics comprises a network of interlocking elongated and prismatic crystals of

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akermanite embedded in a glassy phase. The MWP features similar microstructure, but glassy phase is more evident. In the case of MW, the dramatic decrease of density, shrinkage and mechanical properties recorded at temperatures4800 1C (Table 2, Fig. 2a), would be attributed to secondary porosity, probably developed after completion of the crystallization process [21]. Doping with P2O5 evidently suppressed such detrimental effects on the properties of the MWP glass-ceramics. Consequently, P2O5 improved the sintering ability of the glass compacts because it shifted crystallization towards higher temperatures and broadened the temperature range of densification. Owing to the fact that coarsening usually reduces sintering ability [22,23], the effect of P2O5 is more sounded since the MWP powder comprised more coarsened particles than MW. Evidently, P2O5 should considerably influence the diffusion process of the components. Most probably, phosphorous oxide should increase the stability of the glass against crystallization at the temperature of sintering onset. Earlier studies have reported that small addition of several oxides, such as P2O5, to certain glasses can impart excellent sintering ability [24,25]. Hence, for both the investigated compositions, sintering should have started at temperatures close to

MW

800 ºC 700 ºC 650 ºC

MWP

800 ºC 700 ºC 650 ºC 10

20

30

40

50

60

2θ Fig. 3. Akermanite exclusively forms at X700 1C (JCPDS card of akermanite-synthetic No. 87-0052). Higher crystallinity is observed in the MW than in the MWP composition.

Fig. 4. Characteristic microstructures of (a) MW and (b) MWP glassceramics sintered at 800 1C for 1 h observed at fracture surfaces of bulk samples after immersion in SBF for 31 days. The insets show clearly that the surface of both materials was homogeneously covered with submicron particles enriched in Ca-P (verified by EDS).

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the glass transition point. A liquid phase was formed and wetted the surface of the powder grains. Heating at temperatures close to the softening point causes lowering of viscosity of the liquid and densification advances via viscous coalescence. In several glassceramic systems, crystallization process has been attributed to have an inhibition effect on densification occurring via viscous flow [26–28]. 3.2. Dissolution and mineralization capability in SBF Fig. 5 outlines the change of ionic concentrations in the liquid over immersion time of the powders of both glasses in SBF at 37 1C. Evidently, alkaline reaction was observed in all cases. Dissolution spontaneously took place, even after 1 h. MW dissolves slightly faster than MWP. This finding further supports the aforementioned speculation whereby phosphorous oxide can increase the stability of the glassy phase, not only with regard to the delay of crystallization at the temperature of sintering onset but apparently also towards SBF. However, the plots feature similar characteristics for both glasses. In particular, the concentration of Ca2+ was continuously increased after 1 month of immersion. Phosphorous concentration was rapidly decreased, indicating the active role of P on the transformation of

9.0 8.6 8.2 7.8 7.4

2.0

25 1.5 20 15

1.0

10 0.5 5 0 1

Ca

Ca, Mg, B (mM) pH

9.8 9.4

0.0 1000

10 100 Immersion time (h)

(a) Mg

B

pH

P

30

Si

Na

Al 2.0

25 1.5 20

9.0 8.6 8.2

P,Si, Al, Na/100 (mM)

9.4

30

15

1.0

10 0.5 5

P,Si, Al, Na/100 (mM)

Ca, Mg, B (mM) pH

9.8

7.8 7.4

(b)

0 1

10 100 Immersion time (h)

0.0 1000

Fig. 5. Evolution of ionic concentrations in the liquid over immersion time of: (a) MW and (b) MWP glass powders in SBF at 37 1C.

glasses’ surface during immersion in SBF. The concentration of Si reached a plateau, descending afterwards. There are no considerable changes of the dissolution of Mg2+ and B3+ over immersion time, while there is no clear evidence about the possible Na+ exchange between SBF and the investigated materials. Finally, the concentration of Al3+ was always very low, being at the limit of the sensitivity of the ICP apparatus. Comparison of the curves of Fig. 5 with earlier reports indicates that the evolution of ionic concentrations over immersion time strongly resembles the published reports of oxides containing CaO–SiO2 [29–34], rather than the calcium phosphates, such as apatites, TCP, etc. [35,36]. The reactivity of the investigated glasses should have followed the mechanism proposed by Kokubo: (i) Ca2+ ions are exchanged with H+ resulting in a pH increase, (ii) a silica gel layer forms and provides favourable nucleation sites for apatite, and (iii) apatite layer forms and thickens on the silica gel layer. In the light of this mechanism, the investigated glasses showed evidences of mild biomineralization capability. Fig. 6a shows that submicron precipitates, enriched in Ca and P (found by EDS), were uniformly formed on the surface of MWP glass after 7 days of immersion, but not earlier (i.e. 3 days or shorter). Their tiny size hampered their complete characterization by EDS (e.g. Ca/P ratio), XRD and FTIR (attempts with all these techniques gave ambiguous results and thus omitted). Extensive cracking is also obvious at the surface of the glass. In literature, there are studies reporting similar intensive corrosion of very well-known bioactive glasses [37]. Chemical analysis by EDS at steep angles showed that the glass surface was rich in P and Ca, comparing to the bulk (fracture surface) of (as cast) glass (Fig. 6b). This finding agrees fairly well with the conclusion with regard to the active participation of P ions in surface transformation process of glass during immersion in SBF. The increase of Al should be likely apparent, due to the leaching of the other elements during the first stage (i) of the aforementioned mineralization mechanism, since SBF does not contain Al ions. Nevertheless, prolonged immersion in SBF only widened surface corrosion but did not result in formation of apatite layer. Similar submicron Ca-P-rich precipitates (verified by EDS) were formed all over the surface of both MW and MWP glass-ceramics after 1 month of immersion (Fig. 4). In this case, the remaining glassy phase should have favoured the formation the Ca-P precipitates while the crystalline structure maintained the structural integrity of materials and suppressed corrosion [38]. However, the surface of the glasses appears to enable suitable modifications resulting in a layer, which could act as nucleation centre of Ca-P precipitates, according to the aforementioned mineralization mechanism. For

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(a)

P

O

C

Al

Ca Si

Na Mg

Surface

Cl K

Ca

Bulk 0 (b)

1

2

3

4

5

2261

comparing to other elements. This conclusion is in accordance with the low Al3+ concentration of Fig. 5. We have also obtained similar evidences in our earlier study in the system fluorapatite-anorthite, where Al was not detected in the SBF solution after 108 days of immersion of glass powders in SBF [43]. Therefore, the strong adherence of Al in the glass structure should cause intrinsic structural constraints with regards to the freedom of the glass structure for undergoing the necessary transformations anticipated in stage (ii) at the maximum extent. In the literature, Hench and Andersson have conjectured that alumina inhibits some of the stages of apatite formation retarding therefore the mineralization of osteoids into bone [44]. Kokubo experimentally determined the Al2O3-concentration threshold between bioactivity and bioinertness regimes in Al2O3-containing wollastonite glasses, which is in the range of 1.5–1.7 mol%, while glasses with 5% Al2O3 were entirely inert [45]. The investigated glasses had, however, lower amount of Al2O3 content (Table 1). Therefore, the apparent enrichment of the surface of the glass in Al after the stage (i) of leaching should have locally exceeded this concentration threshold.

Energy (keV)

Fig. 6. (a) Formation of Ca-P-rich submicron precipitates and corrosion evidences at the surface of MWP- glass bulk sample after 7 days of immersion in SBF at 37 1C. (b) Comparison of the elemental analysis of the surface shown in (a) and the bulk (fracture surface) of the virgin (i.e. as cast) glass, obtained by EDS analysis done at very speed angles.

instance, there was an incubation time of 7 days for the formation of the first calcium-phosphate precipitates [29,39]. Kokubo has mentioned 10 days of incubation time for the 20Na2O–80SiO2 glass [2]. Moreover, the fact that the pH value approaches the acidity constant of silicic acid, pKSi(OH)2=9.6 (Fig. 5a) [40], probably supports the existence of stage (ii), whereby silica is hydrolysed to silicic acid. Silicates, which also greatly dissolve from the glasses (Fig. 5) may be actively involved in this process [4,32,40,41]. In earlier studies, dissolution capability and bioactivity performance have been directly related to network connectivity of glasses [3,4,42]. In brief, glasses comprising networks with dense cross-linking feature more rigid structures, which undergo less dissolution, while flexible structures with lower network connectivity can promote dissolution and then mineralization. This work proposes that not only the structure of the initial glass but also the resulting structure of the surface of the glass, derived after the leaching of stage (i) of the mineralization mechanism should be enough flexible to favour apatite formation. Under this perspective, the experimental results of Fig. 6b clearly show that Al has low leaching capability,

3.3. In vitro tests with osteoblasts Under light microscopy, osteoblasts showed no evidence of morphological changes after 24, 48 or 72 h of incubation in the presence of MW and MWP glasses, when compared to control. The viability/proliferation of osteoblasts was not altered in the presence of both glass powders, when compared to control (Fig. 7a). The production of alkaline phosphatase was slightly higher in the presence of the glasses (Fig. 7b). The osteoblasts collagen secretion increased about 20% in the presence of MW and MWP, when compared to control (Fig. 7c). These results are in accordance with earlier findings [46], likely attributed to the silicon content of the glasses. It has been reported that release of silicilic acid enhances production of collagen type I [47] as well as that bioactive glasses containing wollastonite components stimulates in vitro biomineralization [48,49]. Accordingly, the results of Fig. 7c can support fairly well the discussion of Section 3.2, with regards to the evidences of the occurrence of the stages of biomineralization mechanism in the investigated glasses. 3.4. Potential applications, coatings The results showed that both MW and MWP glassceramics exhibited few properties that may be suitable for biomedical applications. These glass-ceramics have homogenous colour and high whiteness, good matching of microhardness with enamel [50], formation of a single (and not multi) phase material during crystallization,

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Optical density

0.3

0.2

0.1

0.0 (a)

MW

MWP

Control

MW

MWP

Control

shows typical microstructures of the interfaces developed between MWP glass-ceramic and the three substrates after 1 h heating at 800 1C. In the case of zirconia, the good matching of the CTEs between TZP and MWP apparently resulted in a continuous interface, free of cracks or gaps (Fig. 8a). Therefore, good affinity between the contacting phases might be also assumed [52]. In the case of coatings on Ti, a reaction zone of TiO2 (by EDS) was clearly visible in back-scattering mode (Fig. 8b). Strong chemical reaction can be suggested since the interface is not planar but the reaction zone was seemingly diffused towards both

0. 5

Optical density

0. 4 0. 3 0. 2 0. 1 0. 0 (b)

Collagen secretion(mg/ml)

75

* 50

25

0 (c)

MW

MWP

Control

Fig. 7. (a) Viability/proliferation. (b) Alkaline phosphatase production. (c) Collagen secretion. 1  105 osteoblasts were plated in the presence of MW and MWP powder. After 72 h of incubation, viability and proliferation were evaluated by MTT assay (a), alkaline phosphatase production by NBT-BCIP assay (b), and the collagen present in the supernatant was measured by SIRCOL method (c). In the latter case (c), the supernatant of all samples, from each condition, were pooled and measured. In (a) and (b) results represent Mean7SD of duplicates from five different experiments (Po0:05) and (c) represent Mean7SD of three different measurements (Po0:05).

and good matching of CTE (9–10  106 K1) with other materials used in biomedicine, such as ZrO2 and Ti [51], and therefore may be appropriate for dental applications. These advantages motivated a preliminary testing of producing coatings using MWP glass-ceramic on zirconia (TZP), titanium and hydroxyapatite. Fig. 8

Fig. 8. Microstructure of interfaces developed between MWP glassceramic and (a) zirconia (TZP), (b) Ti, and (c) hydroxyapatite (HA) after crystallization at 800 1C, 1 h. The grooves shown in (a) and (b) were purposefully introduced by polishing with a 6 mm diamond paste to give a better perspective of the interface and allow the comparison between the hardness of the contacting phases.

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the ceramic and the metal phase. Although continuous and strong interface without gaps was observed, the possible brittleness of the reaction zone in conjunction with the possible mismatch of the thermal and the elastic properties between the three contacting phases might cause failure of the Ti/MWP joint, especially under fatigue conditions. The MWP coating was ideally joined with hydroxyapatite. Evidently, the two phases diffused one to another and no interface can be observed (Fig. 8c). Therefore, production of porous materials, which will function as scaffolds for tissue ingrowth, by combining MWP glass-ceramics and other bioactive ceramics into a composite material structure is possible [53].

4. Conclusions Glass-ceramics of akermanite from a Ca-mica and wollastonite were successfully designed and produced. Accordingly, the liquid phase formed during sintering of glass-compacts facilitated densification due to the presence of B and F and favoured moderate bioactivity due to the presence of Ca and Si. Doping with P2O5 improved the sintering behaviour of the compacts because P2O5 increases the stability of glass against crystallization at the temperature of sintering onset and broadens the range of sintering temperatures. Crystallization occurred at X700 1C and resulted in only one crystalline phase, akermanite, which maintained material’s structural integrity towards body fluids. With regards to the use in biomedical applications, the most important properties of the resulting glassceramics are the good in vitro acceptance from osteoblasts, the good quality of the white colour at the surface and throughout the bulk, the matching of microhardness with tooth enamel, and the possibility to coat other biomaterials, such as Zirconia, Ti or hydroxyapatite.

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