Processing, physico-chemical characterisation and in vitro evaluation of silicon containing β-tricalcium phosphate ceramics

Materials Science and Engineering C 31 (2011) 531–539

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Materials Science and Engineering C j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / m s e c

Processing, physico-chemical characterisation and in vitro evaluation of silicon containing β-tricalcium phosphate ceramics N. Douard a, R. Detsch b, R. Chotard-Ghodsnia a, C. Damia a, U. Deisinger c, E. Champion a,⁎ a b c

Université de Limoges, CNRS, ENSCI, SPCTS, UMR 6638, 123 Avenue Albert Thomas, 87060 Limoges Cedex, France BioCer Entwicklungs-GmbH, Ludwig-Thoma-Straße 36c, 95447 Bayreuth, Germany Friedrich-Baur-Research Institute for Biomaterials, University of Bayreuth, Universitätsstr. 30, 95440 Bayreuth, Germany

a r t i c l e

i n f o

Article history: Received 17 June 2010 Received in revised form 9 August 2010 Accepted 12 November 2010 Available online 18 November 2010 Keywords: Tricalcium phosphate Silicon Powder synthesis Sintering In vitro evaluation

a b s t r a c t For bone grafting applications, the elaboration of silicon containing beta-tricalcium phosphate (β-TCP) was studied. The synthesis was performed using a wet precipitation method according to the hypothetical theoretical formula Ca3 − x(PO4)2 − 2x(SiO4)x. Two silicon loaded materials (0.46 wt.% and 0.95 wt.%) were investigated and compared to a pure β-TCP. The maturation time of the synthesis required in order to obtain β-TCP decreased with the amount of silicon. Only restrictive synthesis conditions allow preparing silicon containing β-TCP with controlled composition. To obtain dense ceramics, the sintering behaviour of the powders was evaluated. The addition of silicon slowed the densification process and decreased the grain size of the dense ceramics. Rietveld refinement may indicate a partial incorporation of silicon in the β-TCP lattice. X-ray photoelectron spectroscopy and transmission electron microscopy analyses revealed that the remaining silicon formed amorphous clusters of silicon rich phase. The in vitro biological behaviour was investigated with MC3T3-E1 osteoblast-like cells. After the addition of silicon, the ceramics remained cytocompatible, highlighting the high potential of silicon containing β-TCP as optimised bone graft material. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Bone mineral is a biological calcium phosphate material which contains various ionic substitutions [1]. Therefore, calcium phosphate materials are widely used for bone replacement in surgery due to their chemical compositions close to the bone mineral. These bioceramics favour bone reconstruction, thanks to high properties of resorbability for β-tricalcium phosphate (β-TCP — Ca3(PO4)2) and good osteoconductivity for hydroxyapatite (HA — Ca10(PO4)6(OH)2) [2]. The bone remodelling process involves the coupled action of osteoblasts (boneforming cells) and osteoclasts (bone-resorbing cells) [3] and silicon (Si) seems to play a role during this process. Indeed, Carlisle suggested that Si is implied in the first stage of mineralisation [4]. Xynos et al. stated that ionic products of bioactive glass dissolution, rich in Si element, increased the proliferation of osteoblast cells, up to 155% of the control [5]. Thus, the doping of calcium phosphates with Si would be a potential method to improve their bioactivity. The synthesis and characterisation of pure Si-substituted hydroxyapatite

⁎ Corresponding author. Tel.: + 33 555457460; fax: + 33 555457586. E-mail address: [email protected] (E. Champion). 0928-4931/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2010.11.008

(Si-HA), where silicate substitutes for phosphate, was well investigated and conducted to the formula Ca10(PO4)6 − x(SiO4)x(OH)2 − x [6,7]. In vivo [8] and in vitro [9] studies revealed an enhancement of the biological behaviour of HA after the silicon doping. In order to prepare calcium phosphate bioceramics with both higher resorption rates than HA and increased bioactivity, the study of silicon doped βtricalcium phosphate appears of interest. Ghaith et al. prepared a silica containing β-TCP material through laser irradiation of a silica sol spin coated at the surface of dense β-TCP pellets [10]. Droplets of an amorphous phase containing silica remained embedded in the β-TCP. Bandyopadhyay et al. also reported the preparation, via a solid state method, of a nontoxic β-TCP with silica as dopant [11]. Wei et al. investigated the (Si, Zn) co-doping of TCP which led to β-TCP or α-TCP or a mixture of the two phases depending on the amount of dopants [12,13]. Except these studies, mainly the doping of α-TCP with silicon has been investigated [14–16]. On this basis, the aim of the present work was to investigate the possible incorporation of silicon into the β-TCP lattice. The study is focused on the powder synthesis via a wet precipitation method and further sintering. Particular attention was given to the determination of the chemical composition of ceramics (nature and location of the chemical phases within the ceramic material). Then, the influence of silicon on the in vitro cytocompatibility of the materials was investigated using a murine pre-osteoblastic cell line.

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2. Materials and methods 2.1. Ceramic processing To respect the charge balance mechanism, silicon containing βtricalcium phosphate, was prepared according to the following hypothetical equation: 2 +

ð3−xÞCa

3−

+ ð2−2xÞPO4

4−

+ x SiO4 → Cað3−xÞ ðPO4 Þð2−2xÞ ðSiO4 Þx : ð1Þ

Though purely hypothetical, this equation was chosen on the theory that tetrahedral silicate (SiO4) could substitute for phosphate (PO4) in the β-TCP lattice, as it is the case in HA [6,7] and suggested by Wei et al. [13]. Two silicon loadings, x = 0.05 and 0.10, were investigated in this study which correspond respectively to 0.46 wt.% and 0.95 wt.% of Si. The corresponding materials are noted Si0.46–βTCP and Si0.95–βTCP. As a reference, a pure β-TCP (x = 0) was also prepared. The synthesis of Si containing β-TCP was performed using the classical route of β-TCP synthesis by a wet precipitation method [17]. In this method, β-TCP cannot be synthesised directly. The compound that precipitate is a calcium deficient apatite Ca10 − y(HPO4)y(PO4)6 − y(OH)2 − y, with 0 ≤ y ≤ 1, (CDHA) which leads, when y = 1, to β-TCP after a thermal treatment above 750 °C. CDHA powders containing silicon were prepared by the addition, at a control rate of 50 mL min− 1, of 1.5 ± 0.1 L of an aqueous solution of diammonium phosphate (NH4)2HPO4 (Aldrich, France) and silicon tetra-acetate Si(OOCCH3)4 (Aldrich, France), into a reactor containing 2.0 ± 0.1 L of a calcium nitrate tetra-hydrate solution Ca(NO3)2·4H2O (Aldrich, France). Molar amount of starting materials and expected molar ratio of the samples are given in Table 1. The reactor was placed under dynamic flow of argon in order to prevent any presence of CO2 which could dissolve in water and result in the precipitation of carbonated apatite. The pH of the solution was maintained at a constant value of 7.0 ± 0.1 by the addition of an ammonium hydroxide solution at 3 mol L− 1 using a pH controller and dosing pump system (Hanna Instruments). The temperature was controlled and regulated at 30.0 ± 1.0 °C. The suspension was continuously stirred and refluxed. As the maturation time could influence the final powder composition [17], different maturation steps, for each silicon loading, were examined. When no silicon was added (x = 0, reference powder) during the precipitation at pH 7 and T = 30 °C, a maturation time of 10 h was applied [17]. For each sampling, the precipitate was centrifuged (500 g, 5 min) and washed with distilled water to remove synthesis residuals. The resulted cake was dried at 100 °C for 24 h and then characterised. For the shaping process, powders were first heat-treated at 750 °C for 30 min [17] and then pressed in a cylindrical die under a compressive stress of 125 MPa. The resulting pellets were sintered in a Super Kanthal furnace in air atmosphere. Sintering parameters (time and temperature) were determined, for each composition, according to the results of the dilatometric analysis described in the next subsection. 2.2. Characterisation methods To assess the phase composition, X-ray diffraction (XRD) and Fourier transformed infra-red (FTIR) analysis were performed either

on as synthesised and heat-treated powders. XRD patterns were collected using Cu Kα radiation on a θ/2θ diffractometer (model D 5000, Siemens, Germany). Phase identification was achieved by comparing the diffraction patterns with the International Center for Diffraction Data (ICDD) Powder Diffraction Files (PDF). FTIR was carried out using KBr technique (Spectrum One, Perkin-Elmer, USA). The Ca/P molar ratio of powders was determined after calcination at 1000 °C for 15 h from integrated intensity ratios of characteristic diffraction peaks according to an ISO standard procedure when Ca/ P≥1.500 [18] and according to Destainville et al. when Ca/P b 1.500 [17]. After calcination, the CDHA precipitates are single-phase β-TCP for Ca/ P=1.500, biphasic mixture of HA and β-TCP for 1.500b Ca/Pb 1.667 or biphasic mixture of β-calcium pyrophosphate (β-CPP — Ca2P2O7) and βTCP for 1.000b Ca/Pb 1.500 whose proportions relate to the Ca/P ratio. Cell parameters of the 1000 °C heat-treated powders were determined by Rietveld refinement of the XRD data with the FullProf Suite software [19]. Data sets were collected over the range 10–180° with a step size of 0.015° and a count time of 5 s. Specific surface area (S) of the powders was measured by the Brunauer, Emmett and Teller method (8 points, Analyzer Micromeritics ASAP2010, USA) after degassing under vacuum at 250 °C. The linear shrinkage was registered by dilatometry in air up to 1300 °C (SETARAM TMA 92, France). The 750 °C heat-treated powders were pressed in a cylindrical die prior to the dilatometric experiment. The relative density of sintered samples was measured by the Archimedean method in water. It was quoted as a percentage of the theoretical density of β-TCP. Polished sintered ceramics were characterised by means of XRD. In order to reveal the microstructure of the ceramics, polished samples were thermally etched for 6 min at 20 °C below the sintering temperature. The outer surface of the thermal etched ceramics was then examined using scanning electron microscopy (SEM, Philips XL30, The Netherlands) and Scion software was used to determine the grain size distribution. The grain size (G) was the equivalent diameter of a disc having the same area as the grain. The surface of the polished samples was analysed by X-ray photoelectron spectroscopy (XPS) with an Axis Ultra DLD spectrometer (Kratos Analytical Ltd., UK) using a monochromatic source Al Kα (1486.6 eV). Peak positions were referenced to the C1s line of the residual carbon set at 285.0 eV. After background correction with the Shirley method, atomic compositions of the samples were determined using the Scofield coefficients [20]. The samples were also examined by transmission electron microscopy (TEM) with a JEOL 2010 microscope (Japan). Chemical analyses were performed using energy dispersion spectroscopy (EDS, Princeton Gamma-Tech, USA) and selected area electron diffraction (SAED) patterns were recorded to investigate the structure of the ceramics. 2.3. In vitro cytocompatibility study All ceramic samples were polished and their average roughness was determined from 1 mm length roughness profiles of 10 surfaces (SJ-201P, Mitutoyo, USA). Before carrying out the cell culture assay, the polished samples were cleaned and sterilised at 134 °C in an autoclave (Systec, Germany). An osteoblast-like cell line, derived from mouse calvaria and denoted MC3T3-E1 (DSMZ, Germany), was used. The cells were cultured at 37 °C in a humidified atmosphere of 95% air and 5%CO2, in alpha-modified

Table 1 Quantities of reagents, expected Si content (x), expected and calculated, from XPS analysis, molar ratio of the sintered ceramics. Sample

nCa (mol)

nP (mol)

nSi (mol)

x (mol)

Ca/P Expected

Calculated

Expected

Calculated

β-TCP Si0.46–βTCP Si0.95–βTCP

1.280 1.259 1.237

0.853 0.811 0.768

0 0.021 0.043

0 0.05 0.10

1.50 1.55 1.61

1.51 ± 0.07 1.36 ± 0.07 1.43 ± 0.07

1.50 1.51 1.53

Not measured 1.31 ± 0.07 1.34 ± 0.07

Ca/(P + Si)

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essential medium (α-MEM, Gibco, Germany) containing 10 vol.% fetal bovine serum (FBS, Sigma-Aldrich, Germany) and 1 vol.% penicillin/ streptomycin (Sigma-Aldrich, Germany). Cells were grown to confluence in 75 cm2 culture flasks (Nunc, Denmark), harvested using Trypsin/EDTA (Sigma, Germany), counted by a coulter counter (Beckman, Germany) and diluted to a concentration of 100 000 cells/mL cell culture medium. After sterilisation, the β-TCP, Si0.46–βTCP and Si0.95–βTCP discs (diameter 14.8 mm), designed to geometrically fit the respective culture well plate, were placed into a 24-well plate (Greiner, Germany). Afterwards, 1 mL cell suspension was plated either onto the upper surface of the materials or in the absence of material (plastic of multiwell plate). After 48 h of incubation onto the materials, cell number, cell viability, cell membrane permeability and cell morphology were determined according to standards [21]. The proliferation of the MC3T3-E1 cells was analysed by means of the cell counting. The adherent cells were first detached from the surface by Trypsin/EDTA and then counted by coulter counter. The cell viability was analysed as the mitochondrial dehydrogenase activity [22] and measured by the WST-1 test (Roche, Germany), with a WST-1 working concentration of 15 μl mL− 1. In this colorimetric assay, the UV/Vis absorbance at 450 nm is proportional to the amount of dehydrogenase activity in the cell grown on the sample. The cell membrane permeability was quantified using Resazurin assay (Sigma-Aldrich, Germany) [23]. The results of this analysis were normalised to the determined cell number and indicated the degree of cytotoxicity caused by the material on the cells. Each experiment was normalised to the reference (plastic of multiwell plate: REF = 100%) and repeated with four samples. Differences between the mean values were evaluated by a student-t test. For SEM-characterisation, cells on samples were fixed in 3 vol.% paraformaldehyde, 3 vol.% glutaraldehyde (Sigma-Aldrich, Germany) and 0.2 M sodiumcacodylate (Sigma-Aldrich, Germany). After a dehydration through incubation with a series of graded acetone, the samples were critical point dried with CO2 (Parr Instrument Company Moline, USA) and sputtered with gold (Cressington, UK). The cell morphology was analysed by SEM (Quanta 200, FEI, The Netherlands).

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Fig. 1. XRD patterns of powder with x = 0.05 (a) raw and (b) 750 °C heat-treated powders for 70 min of maturation and 1000 °C heat-treated powder for (c) 40 min, (d) 70 min and (e) 100 min of maturation.

3.1. Synthesis and characterisation of the powders

(PDF 09-0346) was present as a secondary phase. Thus, at the beginning of the synthesis, the Ca/P molar ratio of the precipitates was lower than 1.500. The increase of the Ca/P molar ratio from 40 min to 70 min of maturation indicated a decrease in the amount of β-CPP in the powders. For 70 min of maturation, only the β-TCP phase was detected, meaning that the Ca/P molar ratio reached the value of 1.500. For longer maturation times, the β-TCP phase was still observed but a diffraction peak attributed to the HA phase (PDF 09-0432) was highlighted. From 80 min of maturation, the Ca/P molar ratio of the precipitate was higher than 1.500 and HA appeared in growing content. For x = 0.10, the same changes of the powder composition were observed (Fig. 3), but shifted to shorter maturation times. The synthesis of a powder constituted of β-TCP was here achieved after only 30 min of maturation.

Fig. 1a presents XRD pattern of the raw powder obtained for x = 0.05 and after 70 min of maturation time. The broad peak observed matched those of crystalline apatite (PDF 09-0432). The synthesised powder was composed of poorly crystalline apatite. The FTIR spectra of the same raw powder exhibited the characteristic absorption bands of CDHA (Fig. 2a). Absorption bands for phosphate groups appeared at 1120–1020 cm− 1 (ν3 PO4), at 960 cm− 1 (ν1 PO4), at 600–560 cm− 1 (ν4 PO4) and at 480 cm− 1 (ν2 PO4). The bands at 3540 and 630 cm− 1 are attributed to the hydroxide groups (νs OH and νl OH respectively) and those at 1140 and 875 cm− 1 to hydrogenophosphate groups (HPO4). The broad band between 3700 and 3000 cm− 1, as well as the one at 1640 cm− 1, corresponded to residual water adsorbed at the particle surface. The ammonium (NH4) and nitrate (NO3) groups resulting from synthesis residuals, appear at 3200, 1380 and 830 cm− 1. Regardless of the amount of silicon or the maturation time, the XRD patterns and the FTIR spectra of the raw powders were similar to those presented in Figs. 1a and 2a respectively. To complete the description of the powders, complementary characterisations after a thermal treatment at 1000 °C for 15 h, as described in Section 2.2, were required. This treatment allows the decomposition of the raw CDHA. The XRD patterns of the 1000 °C heat-treated powders for x = 0.05 (Fig. 1c–e) exhibited different calcium phosphate crystalline phases with the maturation time and as a consequence different global values of Ca/P molar ratio (Fig. 3). The powder was first mainly composed of β-TCP (PDF 09-0169) but β-CPP

Fig. 2. FTIR spectra of (a) raw and (b) 750 °C heat-treated powder with x = 0.05 for 70 min of maturation and 1000 °C heat-treated powders with (c) x = 0, (d) x = 0.05 and (e) x = 0.10 for respectively, 10 h, 70 min and 30 min of maturation time.

3. Results

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Fig. 3. Evolution of the Ca/P molar ratio of the powders calcined at 1000 °C for 15 h with the maturation time and the amount x of silicon.

Table 2 Lattice parameters and cell volume calculated by Rietveld refinements on powders calcined at 1000 °C for 15 h. Sample

a (Å)

c (Å)

Volume (Å3)

β-TCP Si0.46–βTCP Si0.95–βTCP

10.4316 (5) 10.4247 (2) 10.4259 (1)

37.4595 (4) 37.4390 (5) 37.4255 (2)

3530.10 (1) 3523.50 (2) 3523.06 (2)

To fulfil characterisation of the calcined powders, complementary analyses by means of FTIR were performed. Indeed, by means of XRD, the presence of HA as a secondary phase in β-TCP can be detected down to 0.5 wt.% [24] whereas the presence of β-CPP could not be detected from XRD below 4 wt.% [17]. The FTIR spectra of the 1000 °C heat-treated powders containing β-TCP (i.e. after 10 h of maturation for x = 0, after 70 min of maturation for x = 0.05 and after 30 min of maturation for x = 0.10) showed only the characteristic bands of phosphate groups (Fig. 2c–e). The absence of pyrophosphate absorption bands confirmed the absence of β-CPP in the powders [25]. After the thermal treatment at 1000 °C, the OH and HPO4 absorption bands, which were present in the raw powders (Fig. 2a), disappeared. This is in agreement with the dehydration of CDHA into β-TCP that occurs around 750 °C [17,25]. No other peak, which would reveal the presence of Si, was observed. From now on, in this paper, the notation Si0.46–βTCP and Si0.95–βTCP will correspond to powders constituted of β-TCP, synthesised after a maturation time of 70 min and 30 min, respectively. The refined lattice parameters of the powders treated at 1000 °C for 15 h are listed in Table 2. The calculated values of a, c and the unit cell volume for β-TCP were slightly greater than those given in the PDF 09-0169. With growing amount of silicon, the c-axis decreased.

Fig. 4. (a) Linear shrinkage and (b) derivative plot of the linear shrinkage vs. the temperature.

When the silicon content increased from 0 wt.% to 0.46 wt.%, the a-axis and the unit cell volume decreased whereas for the higher amount of silicon they remained constant. 3.2. Processing of ceramic materials An initial thermal treatment at 750 °C for 30 min was performed in order to transform CDHA into β-TCP. The XRD profile (Fig. 1b) of Si0.46–βTCP after this thermal treatment exhibited only the β-TCP phase. On its FTIR spectra (Fig. 2b), only the absorption bands relative to phosphate groups appeared. Same observations were made for the β-TCP and the Si0.95–βTCP powders (data not shown). The decomposition of CDHA into β-TCP has well occurred. This thermal treatment at 750 °C for 30 min is also required to reduce the specific surface area of the powders. Indeed, the specific surface area of the raw powders (about 80 m2 g− 1, Table 3) was too high to allow the pressing of pellets without delamination. The resulting specific surface area authorised pressing samples up to a compaction ratio around 55% of the theoretical density of β-TCP (Table 3). The linear shrinkage of pressed samples and the derivative plot vs. the temperature are presented in Fig. 4. The shrinkage, i.e. the beginning of the sintering, started at 800 °C for β-TCP. The shrinkage temperature increased with increasing silicon content, to reach 900 °C for Si0.95–βTCP. The maximum rate of densification shifted from 1015 °C for β-TCP up to 1100 °C for Si0.95–βTCP. The addition of Si slowed the densification process. After the shrinkage step, an expansion occurred around 1200 °C. This came from the allotropic transformation β → α of TCP which induced material expansion due to the difference of bulk density between the two polymorphs of TCP [26]. Thus, in order to obtain dense materials, the pellets must be sintered at a temperature lower than 1200 °C to prevent this transformation. In this study, the

Table 3 Specific surface area of raw powders (Sraw) and after heat treatment at 750 °C for 30 min (S750 °C), compaction and densification ratio of the ceramics, average grain size (Ga) and average roughness (Ra) of sintered pellets used for biological tests. Sample

Sraw (m2 g− 1)

S750 °C (m2 g− 1)

Compaction ratio (%) 750 ° C — 30 min

Densification ratio (%) 1100 °C — 2 h

Ga (μm)

Ra (μm)

β-TCP Si0.46–βTCP Si0.95–βTCP

89.0 ± 0.3 81.3 ± 0.2 76.5 ± 0.2

11.5 ± 0.1 11.6 ± 0.1 21.1 ± 0.2

52.7 ± 0.2 56.5 ± 0.2 50.9 ± 0.2

99.7 ± 0.1 98.1 ± 0.1 98.0 ± 0.1

1.6 ± 0.8 0.8 ± 0.4 0.5 ± 0.2

0.08 ± 0.01 0.04 ± 0.02 0.06 ± 0.02

N. Douard et al. / Materials Science and Engineering C 31 (2011) 531–539

Fig. 5. XRD patterns of the polished sintered at 1100 °C for 2 h samples.

ceramics were sintered at 1100 °C for 2 h. Under these conditions and regardless of the composition, the relative density of the sintered materials was higher than 97% of the theoretical density of β-TCP (Table 3). XRD data of the polished sintered ceramics did not highlight any difference with the pattern of β-TCP (Fig. 5). This means that no secondary crystalline phase was formed during the sintering. SEM observations of thermal etched samples (Fig. 6) showed a dense microstructure. With the addition of silicon, and the increase of its content, the grain size decreased, as illustrated with the grain size distributions in Fig. 7. The average grain size (Ga) decreased from 1.6 μm for β-TCP to 0.5 μm for Si0.95–βTCP (Table 3). The addition of silicon in β-TCP slowed the grain growth. No significant secondary peak appeared in the distributions, the grain size distributions were monomodal (Fig. 7b). This indicates that during the sintering, grain growth remained normal whatever the composition might be. Fig. 8 shows high resolution XPS spectrum of Si 2p acquired at the surface of polished β-TCP, Si0.46–βTCP and Si0.95–βTCP. As expected, no signal was detected for β-TCP. The Si 2p peaks were detected at 103.6 ± 0.2 eV which would correspond to a Si–O bond [9]. The intensity of the signal increased with the amount of silicon. From these results, calculated molar ratios of the samples were exposed in Table 1. Even if the Ca/P and Ca/(P + Si) molar ratios of the silicon containing β-TCP remained lower than those expected, their values increased with the silicon content. TEM micrographs of Si0.95–βTCP ceramic showed a homogeneous microstructure (Fig. 9a). The grain boundaries were clearly visible, no

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Fig. 7. Grain size distribution of the samples sintered at 1100 °C for 2 h: (a) cumulative frequency; (b) frequency.

phenomenon of silicon segregation was detected. A typical SAED pattern obtained on these crystals is shown in Fig. 10a. A crystalline structure, in good agreement with the PDF card of β-TCP, was revealed. On other areas of the sample (Fig. 9b), the presence of clusters of about 500 nm of diameter (indicated by the white arrows) was revealed. Such regions were observed very locally in the material but with a higher frequency in Si0.95–βTCP than in Si0.46–βTCP. Nothing comparable was observed in β-TCP (data not shown). The typical SAED pattern of this region (Fig. 10b) highlighted an amorphous phase. EDS analysis (Fig. 11) revealed the presence of calcium and phosphorus in the crystallised domain of Si0.95–βTCP whereas the amorphous phase contained mainly silicon and oxygen. 3.3. In vitro cytocompatibility study Before cell seeding, the surface roughness (Ra) of the materials was similar (Table 3). This was necessary to prevent an eventual topography effect on the cell behaviour. Indeed, in vitro cellular response is influenced by physical and chemical characteristics of materials surface [27]. A same surface topography was required to study the sole effect of the addition of silicon in β-TCP on the biological response. During the culture period, no dramatic change of pH of the culture medium was observed and cell culture medium was not influenced by the presence of silicon in the β-TCP. The results of the

Fig. 6. SEM observations of the outer surface of sintered materials: (a) β-TCP, (b) Si0.46–βTCP and (c) Si0.95–βTCP.

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Fig. 10. SAED patterns of Si0.95–βTCP: (a) crystallised phase and (b) amorphous phase.

4. Discussion

Fig. 8. High resolution XPS spectrum of Si 2p peak acquired at the surface of polished βTCP, Si0.46–βTCP and Si0.95–βTCP samples.

cytocompatibility assay are presented in Fig. 12. After 48 h of culture, the cell number was always lower on calcium phosphate materials than on the reference (Fig. 12a). However, the cell number increased with the amount of silicon and significant differences were observed between the two silicon containing materials and pure β-TCP. The mitochondrial activity is downregulated for β-TCP and upregulated for Si0.46–βTCP and Si0.95–βTCP as compared to the reference, suggesting an alteration in cell viability. Interestingly, cell viability slightly increases with the amount of silicon (Fig. 12a), but no significant differences were highlighted between the materials. The cell membrane permeability on all the ceramics was higher than on the reference (Fig. 12b). There was no significant difference in the cytotoxicity between Si0.95–βTCP and the reference, but significantly higher values were calculated for β-TCP and Si0.46–βTCP. Note that a higher value means a higher membrane permeability, and therefore, a higher toxicity. Cell morphology was investigated by SEM (Fig. 13). After 48 h of culture, MC3T3-E1 cells were widely spread on the samples (Fig. 13a, c and e). A dense and nearly confluent monolayer was observed on the ceramics. In all cases, cells have a typical flattened shape (Fig. 13b, d and f). A highly structured cell membrane indicated a good metabolic activity on all samples. The cells exhibited a star-like morphology (numerous cytoplasmic extensions) which shows that MC3T3-E1 cells cultured for 48 h have well adhered and grown.

Fig. 9. TEM micrographs of Si0.95–βTCP: (a) crystallised phase and (b) amorphous phase (arrows indicate the amorphous phase).

The elaboration of silicon containing β-TCP was investigated by synthesising a CDHA with silicon, using a wet precipitation method, and further thermal treatment. The changes in the powder composition during the synthesis have been studied through the evolution of the Ca/P molar ratio. For the calculation of Ca/P ratios, the content of silicon was ignored. Indeed, SAED and EDS analyses (Figs. 10 and 11) revealed that silicon contributed to an amorphous phase only constituted of silicon and oxygen. Assuming that this amorphous phase is SiO2, its content in the powder is about 2 wt.%. In view of this low amount, intensity ratios of the diffraction peaks of the crystalline phases should not change. For all amounts of silicon, the Ca/P molar ratio of CDHA increased with the maturation time (Fig. 3). According to Destainville et al., the synthesis of CDHA happened in two steps: formation of a precipitate, followed by its hydrolysis [17]. During this second step, calcium and hydroxide ions are incorporated into the precipitate inducing an increase of the Ca/P molar ratio. To get β-TCP, a precise maturation time, which decreased with the silicon content (Fig. 3), was required. When no silicon is added during the precipitation, in the same pH and temperature conditions, a much larger maturation time of at least 5 h is needed [17]. This behaviour is explained by an increase of the kinetic of the hydrolysis reaction due to the increase of the Ca/P molar ratio of the reagents with the addition of silicon (Table 1). Finally, the maturation time influences greatly the crystalline phase composition which implies that strict synthesis conditions are required to produce β-TCP after thermal treatment.

Fig. 11. EDS analyses of Si0.95–βTCP: (a) crystallised phase and (b) amorphous phase.

N. Douard et al. / Materials Science and Engineering C 31 (2011) 531–539

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Fig. 12. Cell number, mitochondrial activity (a) and cell membrane permeability (b) of MC3T3-E1 cells on β-TCP, Si0.46–βTCP, Si0.95–βTCP and reference (REF) after 48 h of incubation (normalised to REF = 100%) (*p b 0.05 vs. REF and **p b 0.005 vs. β-TCP).

Fig. 13. SEM images of MC3T3-E1 cells cultured 48 h on (a and b) β-TCP, (c and d) Si0.46–βTCP and (e and f) Si0.95–βTCP.

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Nurse et al. have examined high temperature phase equilibria in the system 2CaO,SiO2–3CaO,P2O5 [28]. The hypothetical theoretical powder composition of our study is Ca(3 − x)(PO4)(2 − 2x)(SiO4)x, which can be rewritten x(2CaO,SiO2),(1 − x)(3CaO,P2O5), with x the amount of silicon previously described. For x = 0.05, the weight ratio of 3CaO, P2O5 in the mixture is 97.2%. According to the phase diagram established by Nurse et al., such a composition should be composed of a solid solution made of silicocarnotite Ca5(PO4)2SiO4 and α-Ca3(PO4)2 below a temperature of 1230 °C. XRD data of the calcined powder showed no traces of either silicocarnotite or α-TCP (Fig. 1). This observation suggests that silicon might not be incorporated into a calcium phosphate structure to form the silicocarnotite phase. Moreover, the same study of Nurse et al. highlighted that the addition of 2CaO,SiO2 from ~1 wt.% (x= 0.02) up to ~20 wt.% (x= 0.30) stabilised the α-TCP phase with the silicocarnotite as second phase for more than 4 wt.% of 2CaO,SiO2, whereas the β-TCP was stabilised below 1 wt.% of 2CaO,SiO2. Several studies also reported that Si addition stabilised the α-TCP polymorph [12,15,16]. On the hypothesis that Si4+ would substitute for P5+ in αTCP, Sayer et al. proposed two mechanisms of charge compensations: formation of oxygen vacancies or calcium excess which respectively leads to Ca3(P1 − xSixO4 − x/2)2 or Ca3 + x(P1 − xSixO4)2 formula [29]. According to ab initio calculations, the second possibility leads to a more stable α-TCP structure [30]. Wei et al. explained the stabilisation of αTCP by Si from crystallographic considerations [12]. Tetrahedral PO3− 4 ions in α-TCP occupy a 180 Å3 volume compared to 168 Å3 in β-TCP. Since Si4+ ionic radius (0.26 Å) is larger than the P5+ one (0.17 Å), its incorporation will favour the α-TCP structure, which offers a larger volume for substitution. All these results showed the difficulties to synthesise Si doped β-TCP in which Si4+ would substitute for P5+, an αTCP being preferentially formed in the presence of silicon. Nevertheless, XRD results of our study did not show any presence of α-TCP. Even if silicon was added during the precipitation reaction, only the β-TCP appeared after high temperature thermal treatment. Moreover, cell parameters refinement (Table 2) highlighted modifications of the a-axis and c-axis after the addition of silicon. Changes of the lattice constants may indicate an incorporation of silicon into the β-TCP. Regardless of the amount of silicon, the a-axis was constant suggesting that the maximum of substitution was reached for, or below, 0.46 wt.% of silicon. The presence of an amorphous phase in Si0.46–βTCP suggests that the maximum substitution amount would be reached for a content value lower than 0.46 wt.%. Wei et al., who studied by means of neutron powder diffraction a (Si, Zn) co-doped β-TCP, have shown that Si substitutes 3.9% of total P sites and is located on P(1) sites [13]. Thus, a substitution of Si into the β-TCP structure may be possible. This study was completed by nuclear magnetic resonance (NMR) analysis of the material which revealed that part of the Si was located in amorphous clusters [31]. XPS, SAED and EDS analyses (Figs. 8, 10 and 11) of our study also showed that Si contributes to an amorphous phase out of the β-TCP. Under the synthesis conditions used in our work, the silicon tetra acetate could undergo hydrolysis that can be followed by a polycondensation process [32]: SiðOOCCH3 Þ4 þ 2H2 O→SiO2 þ 4CH3 COOH:

ð2Þ

The silicon tetra acetate, in this way, produces an amorphous phase in a way similar to that encountered in the synthesis of sol–gel silicate-based glasses. Such results have been previously suggested in the case of hydroxyapatite doped with high amounts of silicon [7,33]. Such an amorphous phase was also observed in the system HA-TCP doped with silicon [29]. Finally, the synthesised materials consist of crystalline β-TCP, probably partially substituted with silicon, coexisting with amorphous clusters of silicon rich phase. The introduction of silicon into β-TCP slowed the densification process (Fig. 4) and the grain growth (Fig. 6 and Table 3). A decrease of 57% of the average grain size was observed after the addition of 0.95 wt.% of Si. Bandyopadhyay et al. reported a decreased of 16% in the

average grain size of a sintered β-TCP containing 1 wt.% SiO2 [11]. The presence of an amorphous phase in the material could explain such sintering behaviour. Indeed, no glassy phase was observed by TEM in the grain boundaries, the amorphous phase is randomly located in the material in the form of small spherical grains (Fig. 9). Thus, the material resembles to a biphasic compound that can be described as a ceramic composite where the minor phase (i.e. amorphous silica) constitutes inert inclusions in a ceramic matrix (i.e. β-TCP). It is well known that in such ceramic systems, the inclusions slow down the sintering of the ceramic matrix [34]. The in vitro cell culture experiments, with osteoblast-like cells over a culture period of 48 h, demonstrated a high biocompatibility of the TCP materials (Fig. 12a). Cell number was significantly lower on β-TCP than on the silicon containing β-TCP. The addition of silicon induced an increase of the proliferation of osteoblast-like cells over the period of 48 h. Cell proliferation on silicon containing β-TCP has also been reported in another study [11]. The mitochondrial activity indicated that cells were active and viable in contact with β-TCP even after the addition of silicon. On the contrary, β-TCP and Si0.46–βTCP appeared more cytotoxic than the reference (Fig. 12b). If the ceramics were cytotoxic for MC3T3-E1, cells would tend to reduce their adherence to the material surface. As SEM images showed dense and confluent monolayer of MC3T3-E1 cells on all the ceramics (Fig. 13a, c and e), it could be stated that cells were not affected by the addition of silicon. With respect to the cell number and mitochondrial activity results, the differences observed for the cell membrane permeability were marginal. As previously observed [35], the TCP materials remained nontoxic even after silicon addition. Since the biocompatibility of β-TCP is not significantly improved with the addition of silicon, this does not mean that silicon does not influence the whole biological behaviour of β-TCP. The biological studies that aimed at increasing the bioactivity of TCP with ionic substitutions are, by a majority, focused on the improvement of α-TCP with silicon. Camiré et al. examined the in vitro and in vivo bioactivity of a cement made of Si-substituted α-TCP [14]. They showed that addition of 1 wt.% of Si significantly improved osteoblastic activity and bone integration. Another study conducted by Langstaff et al. reported the in vivo behaviour of implants made of HA and silicated α-TCP [36]. Results showed the rapid integration of the implant in the host bone tissue accompanied with the formation of a new mineralised bone matrix. A rapid phenomenon of biomimetic precipitation was observed at the surface of a multiphase material containing HA, β-TCP and silicated αTCP probably explained by the presence of Si at the grain boundaries of the material [37]. However, a recent review points out the fact that the beneficial effect of silicon on the enhanced biological properties of calcium phosphate materials has not been demonstrated [38]. More, in our opinion the chemical composition and structure of materials are often not properly defined (unverified presence, location and nature of glassy phase, or of secondary crystalline phases) which may influence greatly the biological behaviour. So, the effect of silicon is still a subject of discussion and well defined materials are essential to state on it, which was the main goal of the present study. From these results, a wider biological investigation will be carried on to better understand the effect of silicon in β-TCP ceramics. In any case, since silicon did not influence the biocompatibility of β-TCP, its presence in the material is potentially very interesting for further biofunctionalisation of the ceramic surface. Indeed, the provided Si–O–Si bounds should facilitate the silanisation of the material for subsequent covalent attachment of biomolecules [39], giving opportunity to develop a new generation of calcium phosphate implants with improved biological properties. 5. Conclusion A β-TCP containing different amounts of silicon has been synthesised via a wet precipitation method. The maturation time

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required to obtain a β-TCP phase is very precise and decreased with the increase of silicon content. The sintering behaviour of Si containing β-TCP was also influenced by the presence of silicon. Dense materials were obtained, and their grain size decreased with the addition of silicon. The structure analysis revealed that silicon was probably partially incorporated into β-TCP but mainly formed amorphous silica clusters leading to small inclusions in the β-TCP matrix. In vitro cell culture experiment on the surface of the dense materials highlighted proliferation of osteoblast-like cells. However, no major difference was observed with the silicon content. The cytocompatibility of β-TCP was not modified with the addition of silicon, so silicon containing β-TCP ceramics were biocompatible. Complementary studies (dissolution and resorption of materials, release of silicon…) are in progress to characterise more precisely the biological response of silicon containing β-TCP. Acknowledgments The authors would like to thank Etienne Laborde for the synthesis of the powders and the XPS experiments and Professor Gilles Trolliard, for the TEM analysis, both from the SPCTS-UMR CNRS 6638. The biological study has been performed in collaboration with Susanne Schäfer and Matthias Schumacher from the Friedrich-BaurResearch Institute for Biomaterials directed by Prof. Günter Ziegler, University of Bayreuth and Franziska Uhl from the BioCer Entwicklungs-GmbH (Bayreuth). The authors acknowledge their help during the experiments and interpretation of the results. References [1] J.C. Elliot, Structure and Chemistry of the Apatites and Other Calcium Orthophosphates, Elsevier Science, Amsterdam, 1994. [2] M. Jarcho, Calcium phosphate ceramics as hard tissue prosthetics, Clin. Orthop. 157 (1981) 259. [3] E.F. Eriksen, G.Z. Eghbali-Fatourechi, S. Khosla, Remodeling and vascular spaces in bone, Bone Miner. Res. 22 (2007) 1. [4] E.M. Carlisle, Silicon: a possible factor in bone calcification, Science 167 (1970) 279. [5] I.D. Xynos, A.J. Edgar, L.D.K. Buttery, L.L. Hench, J.M. Polak, Ionic products of bioactive glass dissolution increase proliferation of human osteoblasts and induce insulin-like growth factor II mRNA expression and protein synthesis, Bioch. Biophys. Res. Commun. 276 (2000) 461. [6] I.R. Gibson, S.M. Best, W. Bonfield, Chemical characterization of silicon-substituted hydroxyapatite, J. Biomed. Mater. Res. 44 (1999) 422. [7] M. Palard, E. Champion, S. Foucaud, Synthesis of silicated hydroxyapatite Ca10 (PO4)6 − x(SiO4)x(OH)2 − x, J. Solid State Chem. 181 (2008) 1950. [8] A.E. Porter, N. Patel, J.N. Skepper, S.M. Best, W. Bonfield, Effect of sintered silicatesubstituted hydroxyapatite on remodelling processes at the bone–implant interface, Biomaterials 25 (2004) 3303. [9] F. Balas, J. Perez-Pariente, M. Vallet-Regi, In vitro bioactivity of silicon-substituted hydroxyapatites, J. Biomed. Mater. Res. 66A (2003) 364. [10] E.S. Ghaith, T. Kasuga, M. Nogami, Preparation of β-tricalcium phosphate containing silica by CO2-laser-irradiation, Key Eng. Mater. 309–311 (II) (2006) 779. [11] A. Bandyopadhyay, S. Bernard, W. Xue, S. Bose, Calcium phosphate-based resorbable ceramics: influence of MgO, ZnO, and SiO2 dopants, J. Am. Ceram. Soc. 89 (2006) 2675. [12] X. Wei, M. Akinc, Resorption rate tunable bioceramic: Si&Zn-modified tricalcium phosphate, Ceram. Eng. Sci. Proc. 26 (2005) 129.

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