Effects of silicate and carbonate substitution on the properties of hydroxyapatite prepared by aqueous co-precipitation method

Materials and Design 87 (2015) 788–796

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Effects of silicate and carbonate substitution on the properties of hydroxyapatite prepared by aqueous co-precipitation method L.T. Bang a, S. Ramesh a,⁎, J. Purbolaksono a, Y.C. Ching a, B.D. Long a, Hari Chandran b, S. Ramesh c, R. Othman d,e a

Center for Advanced Manufacturing & Material Processing, Department of Mechanical Engineering, University of Malaya, Kuala Lumpur 50603, Malaysia Division of Neurosurgery, Faculty of Medicine, University of Malaya, Kuala Lumpur 50603, Malaysia Centre for Ionics University of Malaya, Department of Physics, Faculty of Science, University of Malaya, 50603 Kuala Lumpur, Malaysia d School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia, 14300 Nibong Tebal, Penang, Malaysia e Faculty of Manufacturing Engineering, Universiti Teknikal Malaysia Melaka, 76100 Durian Tunggal, Malacca, Malaysia b c

a r t i c l e

i n f o

Article history: Received 13 February 2015 Received in revised form 27 July 2015 Accepted 16 August 2015 Available online 25 August 2015 Keywords: Hydroxyapatite Ionic-substituted hydroxyapatite Characterization In vitro evaluation Cell response

a b s t r a c t This research was conducted to investigate the effects of silicate and carbonate substitutions in hydroxyapatite (HA) on phase retention, physical properties, and in vitro biological response. A wet chemical method was employed to synthesize silicate-substituted HA (Si–HA) and carbonate-substituted HA (CHA). It was shown that the presence of silicate and carbonate ions in the HA lattice increases the specific surface area and Ca/P molar ratio. In addition, the substitution also resulted in a reduction in the crystallinity of the HA powder and promoted better bioactivity of the sintered body. This was confirmed through in vitro cell culture experiment which revealed that the osteoblast cells adhered well on the surface of the ceramic and that ionic-substituted HA exhibited better cell performance than the control HA. © 2015 Elsevier Ltd. All rights reserved.

1. Introduction Bioceramics made of calcium phosphates (CaPs) have been widely used as artificial bone substitute because of their excellent biocompatibility, bone bonding ability, and chemical composition similarity with natural bone [1,2]. Approximately 70 wt.% of bone is made of hydroxyapatite (HA: Ca10(PO4)6(OH)2), and therefore, HA has been intensively studied as a bone substitute material [3,4]. However, stoichiometric HA also does not dissolve over time but instead remain as a permanent fixture which can be susceptible to long-term failure [5]. Therefore, stoichiometric HA would not be suitable as bone substitute material which requires replacement by natural bone and hard tissues [3]. In addition, native bone differs from stoichiometric HA incorporated with other ions, such as carbonate, silicate, magnesium, and zinc. Reports confirm that the addition of ions that is present in native hard tissue into HA structure can lead to advantageous effects on biomaterial properties, such as structure order, solubility and dissolution rate. Therefore, a key target of biomaterial research is the preparation of a synthetic HA bone-substitute ceramic that mimics the chemical composition of hard tissue. Ionic substitution into HA (or CaPs) has been widely studied recently [6–9]. Bone apatite contains approximately 3 wt.% to 8 wt.% carbonate ion, and hence, bone apatite has been referred to as carbonate apatite (CHA) ⁎ Corresponding author. E-mail address: [email protected] (S. Ramesh).

http://dx.doi.org/10.1016/j.matdes.2015.08.069 0264-1275/© 2015 Elsevier Ltd. All rights reserved.

[5]. The incorporation of carbonate ions is known to pose a considerable impact on the crystal lattice of apatite structure and on the mineralization process [10]. Introduction of carbonate ions in HA promotes dissolution and could enhance the osteointegration rate [11–13]. In addition, CHA can be resorbed by osteoclasts and be replaced with new bone [14]. Therefore, carbonate-substituted HA or CHA is a prospective candidate for bone-substitute material. Apart from carbonate, silicon is a trace element found in bone with a specific metabolic role connected to bone growth, especially during the initial formation stages [15,16]. Silicate substitution at a small concentration simulated the activation of seven families of genes in osteoblast [17], thus increasing osteoblast proliferation, differentiation, and bone extracellular matrix production. Although no study has linked the improved biological performance of Si-substituted CaPs to Si release, and no evidence of silicon released rate from Si–HA in vivo is reported, the presence of Si in the HA structure is widely accepted to cause chemical or topographical changes that would eventually lead to a change in biological responses [18]. In addition, Si substitution promotes ion release that is essential for biological process [19]. The solubility of ionic substituted HA was observed to be higher than that of pure HA. This difference can be attributed to an increase in structural order, which is due to the presence of foreign ions [20–22] in the HA structure. However, only a few studies have investigated the ion release in synthetic fluids [22]. Although research on silicate or carbonate-substituted HA have been reported in the literatures [7,15, 23], the actual effect of the ion substitutions have not been well

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understood. In addition, a systematic comparison of the physiochemical and CO2− properties of SiO4− 4 3 substituted HA has not been carried out. Also, the interaction between biological cells and ionic substitution and CO2− (SiO4− 4 3 ) of HA synthesized using similar precipitation conditions have not been studied. Therefore, the purpose of the present work 2− is to investigate the effect of ion substitutions of SiO4− 4 and CO3 into the HA structure on the phase retention, physiochemical properties, and early in vitro biological response. The paper reports the preparation and characterization of CHA and silicate-substituted HA (Si–HA) by a precipitation method. In this work, the prepared powders were subjected to heat-treatment at various temperatures (800 and 1200 °C). The powder analysis was carried out by means of powder X-ray diffraction (XRD), transmission electron microscope (TEM), fourier transform infrared (FTIR), X-ray fluorescence (XRF), inductively coupled plasma atomic emission spectroscopy (ICP/AES), specific surface area by means of the BET (Brunauer,Emmet and Teller) while the morphology of heat-treated compact was observed under scanning electron microscope (SEM). In addition, the biocompatibility nature of these materials in vitro with MC3T3-E1 osteoblast-like cell was also evaluated. 2. Experimental procedures 2.1. Powder synthesis Preparative conditions of alkalinity pH = 9.4 ± 0.1 and reaction temperature of 40 °C ± 1 °C were used to synthesize pure HA and HA 2− substituted with SiO4− 4 (1.6 wt.% Si) and CO3 (8 wt.% to 10 wt.% carbonate content) by aqueous precipitation. The reaction was allowed to proceed in a closed system of a reaction flask placed in a heating mantle and under overhead stirring at 400 rpm. Pure HA powder, Si–HA powder, and CHA powder were synthesized using calcium hydroxide [Ca(OH)2, 96% purity, Fluka], phosphoric acid (H3PO4, Merck), silicon tetraacetate [Si(COOCH3)4, 98% purity, Sigma-Aldrich], and CO2 gas as Ca, P, Si, and carbonate sources. The Ca/P molar ratio of the precursors was based on 1.67 for the synthesis of the HA, CHA and Si-HA powders. To obtain the highest degree of carbonation and a B-type CHA (carbonate ion replace for phosphate ion) precipitation with respect to A-type (carbonate ion replace for hydroxyl ion), the CO2 gas flow was set at 1 bubble/2 s as outlet flux [12]. The PO34 −solution was added drop-wise (~ 1.25 mL/min) to Ca2 + solution under continuous stirring to obtain pure HA. For the synthesis of Si–HA, a preset amount of Si(COOCH3)4 was first dissolved in the Ca2 +solution prior to the addition of PO34 − solution. To form CHA, the PO34 − solution was added to Ca2 + solution, while CO2 gas (40 mL/min) was passed through the reaction flask. The pH was maintained at 9.4 ± 0.1 by a pH meter by adding small amount of ammonium hydroxide (NH4OH 29%, J.T. Baker, USA). The final solution was continuously stirred for 2 h at 40 °C and left to settle for 1 day. The precipitate was then separated by filtering, washing with deionized water, and drying at 70 °C for 24 h. Finally, the precipitate was ground into a fine powder prior to uniaxial compacting at 60 MPa into disc samples (11 mm diameter × 3 mm thickness). The HA and Si–HA green disks were sintered in air at 1200 °C for 2 h, whereas the CHA compact was sintered at 800 °C for 2 h in CO2 gas (40 mL/min) to prevent carbonate loss. The heating and cooling rates were set at 5 °C/min. 2.2. In-vitro experiment 2.2.1. Simulated body fluid (SBF) test In vitro experiments were performed on substrates of the heattreated pure HA, Si–HA, and CHA samples (25 ml SBF, sample size ~10 mm × 2.8 mm). Ion release evaluation was performed in triplicate by immersing the compacts in a SBF solution at 36.5 °C for 1, 3 and 7 days. The SBF solution was prepared according to the procedure described by Kokubo et al. [25]. After the predetermined soaking time,

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the samples were removed, the ion concentrations of liquid media were analyzed by ICP, and pH was measured by a pH meter. 2.2.2. Cell culture test: cell morphology and proliferation Prior to cell culture, all heat-treated samples were sterilized by heating in a vacuum oven at 120 °C for 3 h. MC3T3-E1 osteoblast-like cells (Riken Cell Bank, Tokyo, Japan) were cultured in L-glutamine containing alpha-minimum essential media (α-MEM, GIBCO/Invitrogen, Grand Island, NY, USA) supplemented with 10 vol.% fetal bovine serum (FBS, Invitrogen, Carlsbad, CA, USA), and 1 vol.% penicillin (10,000 units) and streptomycin (10 mg/mL). Cells were maintained at 37 °C under 5% CO2 in a humidified atmosphere. The medium was changed every 2 days. At confluence, adherent cells were passaged and harvested using 0.25% trypsin–EDTA (Trypsin; GIBCO, Invitrogen™, NY. USA). Cells at passage 2 were used for culture on the sample surface with a cell density of 2 × 104 cell/mL. Each experiment was performed in quadruple (n = 4) for each sample group. Cell morphology and attachment ability were observed after 4 h, 1 day and 3 days of culture duration. Cell proliferation was evaluated after 2, 4, 6 and 8 days of culture using Alamar Blue dye test (AlamarBlue™; BioSource International Inc., Camarillo, CA, USA). At the end of each time point, the culture medium was replaced with culture medium containing 10% Alamar Blue reagent. Specimens with cells were then further incubated for 4 h. The resulting 200 μL solution was obtained from all wells, placed in a 96-well clear bottom plate, and the fluorescence was measured on a plate reader (Tecan Infinite M200, Austria GmbH, Grӧdig, Austria) using an excitation and emission wavelength of 520 and 590 nm, respectively. 2.2.3. Osteoblastic differentiation: alkaline phosphatase activity Alkaline phosphatase (ALP) activity was measured by LabAssay™ ALP (Wako, Japan) using p-nitrophenylphosphate as the substrate following instruction of the manufacturer. Cells were seeded onto specimens at an initial density of 4 × 104 cells/mL. After 7, 14 and 21 days, cells attached on the specimens were rinsed twice with PBS and lysed with cell lysis buffer M (Wako, Japan) for 30 min. After centrifugation, 20 μL cell lysate supernatant was incubated with the assay buffer at 37 °C for 15 min and the absorbance was spectrophotometrically measured at 405 nm. Total protein content was measured using Bio-Rad Protein Assay (Pierce Chemical Co., Illinois, USA). 5 μL of the cell lysate supernatant was added to 100 μL assay buffer and shake for 1 min, and then spectrophotometrically measured at 595 nm. ALP activity was normalized by total protein content and expressed as unit/μg protein. 2.2.4. Mineralization assay Cells were seeded at an initial density of 4 × 104 cells/mL. After 14 days of culture, specimens were washed twice with PBS and the cells were fixed with 99.5% methanol for 30 min at room temperature. After rinsing with distilled water, specimens were stained by Calcified Nodule Staining Kit (Cosmo Bio Co., Ltd) solution composed of Alizarin Red S and Chromogenic Substrate for 5 min. After removing excess dye by substrate-containing buffer, the images of stained specimens were observed using an optical microscope (LG-PS2, Olympus, Tokyo, Japan). For the calcium quantification, after 14 and 21 days of culture, specimens were washed twice with PBS and ultrapure water. 1 ml of 0.1 mol/L HCl was subsequently added and the plate was subjected to vibration for 12 h at room temperature. The suspension was centrifuged at 12,000 rpm for 5 min. Calcium content of the supernatant was measured using a calcium C-test Wako kit (Wako, Japan) at an absorbance wavelength of 610 nm. 2.3. Material characterization The phases present in the synthesized powders were examined by X-ray diffraction (XRD; D5000 Siemens) using CuKα radiation (λ =

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1.54060 Å) operating at 40 kV and 20 mA over the 2θ range 10–60° at a step size of 0.03 and counting time of 4 s per step. The (002) XRD peak was selected for determining the crystallite size using the Scherrer's equation (Eq. (1)) since it is the strongest peak without any overlapping with the HA, Si–HA, and CHA peaks. d¼

Kλ β cos θ

ð1Þ

where d = crystallite size, K = the shape factor equal to 0.9, λ = wavelength of Cu Kα radiation equal to 1.5406 Å, θ = half of the diffraction angle and β = full width at half maximum (FWHM). The diffraction peak at about 25.9° (2θ) corresponding to the (002) Miller plane family of the apatite and the diffraction peak at about 32.1o (2θ) corresponding to the (300) Miller plane family, were chosen to calculate the distance between crystal planes along to the crystallographic axis c and a, respectively according to Eq. (2) 1 2

d

3

¼

2

4 h þ hk þ l 3 a2

!

was heat-treated at 550 °C for 5 h in air to eliminate the carbonate adsorbed on the material surface. The Si content substituted in Si–HA and the ion concentration of Ca and Si in the SBF medium after immersion were determined by inductive coupled plasma (ICP) spectrometer (ICP/AES, ARL-3410). The detection limits for the ICP analysis are 10–100 ppt for Ca, 0.1–1 ppb for P and Si. In this experiment, the solution was diluted in nitric acid for 500 times. The as-synthesized powder morphology was observed by transmission electron microscopy (TEM, Philips CM12). The size and morphology of heat-treated samples, as well as cell morphologies and viability of MC3T3-E1, were observed by scanning electron microscope (SEM: S-3400 N, Hitachi High-Technologies Co., Tokyo, Japan). For cell morphology observation, dehydration of cells was performed by sequential immersion of the specimen in serially diluted ethanol solutions (50, 60, 70, 80, 90, and 100 vol.%) for 30 min each. Finally, the specimens were dried by hexamethyldisilazane evaporation and sputter-coated with gold prior to SEM examination.

2

þ

l c2

ð2Þ

3. Results and discussion 3.1. Material characteristics

where d is the distance between the crystal planes obtained from XRD results at the (002) and (300) Miller plane family of HA, Si–HA and CHA, respectively. X-ray fluorescence spectrometry (XRF; Rigaku RIX-300 wavelength dispersive) was used to obtain the Ca/P ratio of the as-prepared powders. The specific surface area was measured by nitrogen adsorption method and Brunauer–Emmet–Teller (BET) using a surface area analyzer (Flow Sorb 2300-Micromeritics). For the sample preparation, about 1 g powder sample after being sieved through 300 μm mess was placed in a glass tube and outgassed before measurement. The chemical absorption characteristics of heat-treated powders were characterized by Fourier transform infrared spectroscopy (FTIR; Perkin-Elmer FT-IR 2000, FTIR spectrometer). For this purpose, 1 wt.% of the powder was mixed and ground with 99% of potassium bromide (KBr, 0.2 g). The spectrum was taken in the range of 400–4000 cm−1 with resolution 4 cm−1 and 64 iteration scans. The carbonate content of the CHA powder was determined using an elemental analyzer (CHN test; Perkin Elmer series 2, 2400 CHNS/O). Prior to CHN analysis, the CHA powder

The size and morphology of the as-synthesized powders are shown in Fig. 1. All precipitated powders were observed to be composed of nanosized, needle-like particles in the range of approximately 8.7 nm to 20.7 nm in width and 40.2 nm to 81.7 nm in length. These results indicate that nanosize particles were successfully obtained. No discernible difference in the morphology of HA and ionic-substituted HA was observed. Experimental conditions such as pH and temperature is known to possibly influence the resultant phase and particle morphology [26]. In this work, similar particle morphology observed is attributed to the identical synthesis conditions employed in this study, in which the pH (9.4) and temperature (40 °C) were maintained. Table 1 shows the physical and chemical properties of the as-synthesized powders. The presence of silicate and carbonate decreased the crystallite size of HA. The measured Ca/P molar ratios of the as-synthesized Si–HA and CHA powders were found to be higher than the stoichiometric value of 1.67, which indicates 4− 3− groups in HA that the substitution of CO2− 3 and SiO4 ions for the PO4

Fig. 1. TEM micrographs of as-prepared powders: (a) HA, (b) Si–HA, and (c) CHA.

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Table 1 Physical properties of as-prepared HA, Si–HA and CHA powders. Sample

HA Si–HA CHA

Ca/P ratio

1.69 ± 0.02 1.85 ± 0.03 2.08 ± 0.01

Lattice parameters (Å) a ±0.0005

c ±0.0005

9.424 9.470 9.422

6.880 6.926 6.879

Crystallite size (nm)

Specific surface area (m2/g)

26.36 ± 0.003 22.81 ± 0.004 23.12 ± 0.004

76.96 ± 0.2 85.58 ± 0.3 79.87 ± 0.3

had taken place. These substitutions reduced the amount of the PO3− 4 group, thus leading to an increase in the Ca/P ratios [13,24]. A higher specific surface area was obtained in Si–HA and CHA samples compared with HA, suggesting that ionic substitution effectively increases the surface area, which would enhance biological activity in vivo. Lattice parameters (a and c) were calculated for a hexagonal structure as listed in Table 1. It was found that there is an increased in the length of the a-axis and c-axis in the as-prepared Si–HA, whereas the a-axis value in the CHA powders decreased if compared with pure HA. With reference to silicate substituted groups, the average length of Si\\O and P\\O bonds are 1.66 and 1.57 Å, respectively [27]. Silicate substitution resulted in an increase in the c-axis of HA unit cells, and, in certain observations, a small decrease in the a-axis [24,28]. The subby SiO4− is assumed to contribute to the increase in stitution of PO3− 4 4 the lattice parameters of the Si–HA powder [23]. However, for the CHA powder, the a-axis has been reported to increase when carbonate was substituted for the hydroxyl group (A-type CHA), but decreased when carbonate replaced the phosphate group (B-type CHA) in the apatite structure [29]. Therefore, the decrease in the lattice parameter of the a-axis in this research work was due to the substitution of carbonate for the phosphate group (B-type CHA). B-type CHA is known to enhance solubility and collagen deposition in vivo compared with HA or A-type CHA [20]. The XRD patterns of the heat-treated HA, Si–HA, and CHA are shown in Fig. 2. It was found that HA and Si–HA retained the HA phase (JCPDS No# 9-432) with no secondary phases present after heat treatment at 1200 °C, whereas the CHA sample was stable at 800 °C. In fact, the CHA sample decomposed to CaCO3 upon heat treatment at 900 °C in the CO2 gas atmosphere (data not shown). Therefore, evaluation was performed for CHA after heat treatment at 800 °C.

Fig. 2. XRD patterns of sintered ceramics: (a) HA at 1200 °C, (b) Si–HA at 1200 °C, and (c) CHA at 800 °C. All peaks corresponded to the hydroxyapatite phase (JCPDS No# 9-432).

Fig. 3. FTIR spectra of heat-treated powders: (a) HA at 1200 °C, (b) Si–HA at 1200 °C and (c) CHA at 800 °C.

According to Marchat et al. [30], the theoretical maximum limit of incorporation of Si into a hexagonal apatitic structure is y b 1.5 for the formula Ca10(PO4)6–y(SiO4)y(OH)2–y(VOH)y. However, this limit is a function of temperature and atmosphere during heat treatment. As a result, the Si1.0HA (y = 1.0) and Si1.25HA (y = 1.25) powders showed an apatitic structure after calcination at 400 °C for 2 h that were unstable at 1000 °C for 15 h in comparison with Si–HA with a silicon content ranging between 0 ≤ y ≤ 0.75. In the present work, the silicon content that was selected for substitution in the HA lattice was 1.6 wt.%, where the ySi in the aforementioned formula was approximately 0.57. Thus, the thermal stability of the HA phase in the Si–HA sample at 1200 °C in this work can be explained by the substitution of silicon, which is within the prescribed limit [31]. The FTIR spectra of heat-treated samples are shown in Fig. 3. Typical bands of HA (Fig. 3a) were observed for the phosphate groups at approximately 960 to 1100 and 560 cm−1 to 604 cm−1, along with the bands corresponding to hydroxyl groups at approximately 3570 bands at around 890, 752, and and 630 cm−1. The presence of SiO4− 4 500 cm−1 in Si–HA together with the phosphate and hydroxyl groups (Fig. 3b) indicates silicate that is substituted in the HA lattice structure [24]. The Si content measured by ICP in the as-synthesized Si–HA sample was approximately 1.57 ± 0.05 wt.%, which is comparable with the starting value of 1.6 wt.%. As indicated in the FTIR result, Si\\O\\Si band (at approximately 798 cm−1 to 800 cm−1) in amorphous silica [30], which may exist on the material surface, was not detected in the Si–HA sample. This result suggested that silicon had been totally substituted into the HA lattice. Typical bands of carbonate groups in CHA were observed at approximately 1410, 1455 and 875 cm−1 (Fig. 3c). These results indicated that CHA obtained in this work is a B-type CHA, in which PO3− 4 is substituted [32]. The carbonate content of CHA sample measured by CHN by CO2− 3 analysis is approximately 8.0 ± 0.2 wt.%, which is close to the typical amount of carbonate in human bone [14]. The presence of OH− groups in all HA samples suggests that the dehydroxyalation of HA, Si–HA, and CHA did not occur at the evaluated temperatures [33]. The most notable effect of silicate and carbonate substitution on the FTIR spectra of Si–HA and CHA is the change in PO3− 4 and OH− bands. The phosphate bands (v2, v3, and v4) and hydroxyl bands [at 3570 (vS) and 630 cm−1 (vL)] are less intense in comparison with those of the HA sample. The incorporation of silicon (or silicate) into the phosphorous (or phosphate) sites in the lattice is believed to result in a change in the bonding and symmetry of the phosphate group [23,34], and induce OH− loss to compensate for the extra negative charge of the silicate groups [24,35]. The slight decrease in OH− group

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observed at 3570 cm−1 in CHA compared with that of HA favored the formation of carbonate group in B-type CHA. In addition, the substituwith PO3− was confirmed by lattice parameter measuretion of CO2− 3 4 ment (Table 1) for B-type CHA.

3.1.1. Wavenumber (cm−1) The morphology of heat-treated samples as observed by SEM is shown in Fig. 4. The sintered HA and Si–HA at 1200 °C exhibited dense compacts (Fig. 4a, b), and the grain sizes were approximately 0.59 ± 0.03 μm for HA and 0.37 ± 0.03 μm for Si–HA. However, the SEM image of CHA (Fig. 4c) heat-treated at 800 °C showed a porous structure with individual grains having sizes below 100 nm. These results are in agreement with previous reports [24,36] indicating that Si–HA composition presents significantly smaller grain size than the corresponding HA ceramic when sintered at the same temperature. The decrease in grain size of Si–HA is consistent with the specific surface area result given in Table 1.

Fig. 4. SEM images of sintered samples: (a) HA at 1200 °C, (b) Si–HA at 1200 °C and (c) CHA at 800 °C.

3.2. In vitro analysis The pH value and ion concentrations of Ca and Si of SBF after immersion of HA, Si–HA, and CHA for 1, 3, and 7 days are summarized in Fig. 5. The value given is an average of three specimen measurements (in ppm). The estimated error in ICP analyses for the measurement of each point in Fig. 5 is less than 5%. When HA and ionic-substituted HA samples were immersed in SBF, fluctuation in pH was recorded, and the values were generally below the standard pH of SBF prepared in this work. The lower pH recorded is due to the increase in H+ concentration, which would enhance the dissolution of the surface layer [37]. According to Fig. 5, the calcium concentration from Si–HA and CHA samples notably increased after 1 and 3 days immersion, respectively, while silicon concentration increased continuously with immersion duration. This finding was due to the ion release of Ca and Si from the ionic-substituted HA to SBF.

Fig. 5. pH and ion concentration of SBF solution after immersion of samples.

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Fig. 6. SEM images showing the attachment of osteoblast cell on the surface of HA, Si–HA, and CHA samples.

Calcium release for CHA was higher than that of Si–HA, and Ca was continuously released from Si–HA after 3 days. The release of Ca and Si from the ionic-substituted HA indicated the solubility of these materials [20]. In fact, various defects such as dislocation and grain boundary are introduced into crystal lattice of Si–HA and CHA, and dissolution begins at these defects. In addition, the surface area of the sample is another important factor that controls the dissolution of ionic-substituted HA. Higher surface area would promote ion release and subsequently increase solubility [38]. As shown in Table 1, the higher surface area of Si–HA and CHA in comparison with HA could be a possible reason for this enhancement. The ion concentration of Ca and Si (in Si–HA) reached a maximum point after a certain immersion duration, i.e., Ca and Si ions in Si–HA reached a maximum at 3 days, and then decreased with further increase in soaking duration. The increase or decrease of ion concentrations in the SBF is as a consequence of dissolution and precipitation processes when a bioactive material is immersed in SBF for certain duration. At the beginning of immersion, the increase of Ca and Si ion concentrations is attributed to the dissolution of the samples. When the ion concentration in the SBF reaches supersaturated with respect to apatite, the precipitation occurs thus leading to the formation of amorphous calcium phosphate or apatite layer deposited on the sample surface and as such reduces the ion concentration. As shown in Fig. 5, the decrease in ion concentration and pH for all samples could be an indicator of the occurrence of bone-like apatite layer precipitation during soaking in SBF [39]. This observation indicated that Si–HA and CHA possess higher bioactivity and solubility in comparison with those of HA, as depicted in Fig. 5. The morphological feature and attachment ability of MC3T3-E1 osteoblast-like cells on the surface of HA, Si–HA, and CHA after 4 h, 1 day and 3 days of culture duration are shown in Fig. 6. These SEM images show the presence of osteoblast cells attached on these surfaces thus confirming the biocompatibility nature of the materials. An increase in the density of osteoblast cells, as well as a change in the typical morphology from elongated (4 h) to well-flattened polygonal shape (after day 1 to day 3) were observed. Cells attached, spread,

and grew on all samples with filopodia extending from the cell edges. After day 1 and day 3 of incubation, the cells have spread and are in physical contact with each other to form a dense cell layer, covering the surface of samples to induce extracellular matrix development [40]. Interestingly, cells attached on CHA samples exhibited more flattened appearance with abundant filopodia extension after 4 h in comparison with the others. The proliferation of MC3T3-E1 osteoblast-like cells on HA, Si–HA, and CHA samples is shown in Fig. 7. The data is presented as the average of the four samples ± standard deviation. A statistically significant difference (p b 0.05) exists in cell numbers over culture time. Cells were found to proliferate, and their number increased over time, representing an active proliferation period. The comparison of proliferation of samples at the end of each time point revealed that cells proliferated well

Fig. 7. MC3T3-E1 cell proliferation on HA, Si–HA, and CHA samples.

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on both Si–HA and CHA specimens within the culture period. The significant increase in osteoblast cells is attributed to several factors, in which the metabolic activity is more likely to be caused by topography [41]. In this work, the high specific surface area, smaller particle size, and the smaller crystallite size (Table 1) may explain the higher proliferation 4− of Si–HA and CHA samples. The substitution of CO2− 3 and SiO4 ions resulted in a significantly smaller particle size. Thus, the increase in grain boundary at the surface, which is due to the smaller grain size, would be beneficial for cell proliferation [42]. Although CHA possesses a preferred surface for cell to proliferate, the proliferation activity of CHA was considerably lower than that of HA after 2 days. This finding could be due to the loose compaction of CHA particles, shown in Fig. 4c, which could inhibit cell growth at the beginning of culture. However, the proliferation of CHA reached a comparable value and surpassed those of HA and Si–HA samples after 4 days and 6 days, respectively. Besides from surface morphology and topography, material chemistry exerts an effect on the interactions between cells and materials [43]. A number of studies had demonstrated that the direct interaction of cell with the material surface may be responsible for osteoinduction– osteoconduction behavior through the dissolution and release of specific ions from the material implanted in a biological environment. These ions act as a guide signal to attract local cells to facilitate osteoblast proliferation [13,44]. For instance, silicon substitution in α-TCP cement enhanced cell proliferation because of the Si released into the culture medium [16]. In addition, the in vitro studies on Si–HA reported that an increase of silicon content stimulated osteoblast proliferation [21, 45,46]. On the other hand, CHA showed better cellular response than HA because of the higher level of Ca2+ ion extension and, consequently, the change in membrane potential for ionic exchange [47]. Based on these studies, an increase in ion release leads to additional enhancement for cell proliferation. In the present study, the release of Si4 + in Si–HA and the higher Ca2 + concentration of Si–HA and CHA samples in comparison with HA indicate a better dissolution by releasing ions (Fig. 5). As a result, Si–HA and CHA samples exhibited a greater degree of chemical surface interaction and, thus, a better interaction with the cells. However, from day 6 to day 8 of culture, cells on HA and Si–HA samples continued to proliferate, and then proliferation was comparable for all samples after day 8. This finding is in general agreement with the results obtained in previous study [48], which showed that both HA and CHA presented insignificant cell proliferation after a certain period, i.e., after 7 days. However, a decrease in cell proliferation of CHA sample is observed at day 8 of culture. This can be explained that for the bone formation, after initial adhesion to the surface of an implant, osteoblast actively proliferates. At the end of proliferation, the extracellular matrix development and maturation begin. Osteoblasts start to differentiate to

calcium depositing cells. At this stage, alkaline phosphatase is increased as well as osteoblasts synthesize and deposit osteocalcin along with other matrix proteins. Therefore, cell growth, in fact, reaches a peak before the cells begin to differentiate and finally mineralize into bone. As such, a decline in cell proliferation normally signals that cells are in the differentiation stage [49]. The ALP activities of cells on HA, Si–HA and CHA substrates are shown Fig. 8. ALP is widely used as an early osteoblast differentiation marker, and its increased activity is associated with osteoblastic differentiation. The ALP activities of the cells grown on all samples increase continuously with an increase in culture time of 7 to 14 days. This result suggests that these cells are able to carry out osteogenic functions and differentiate on all the samples [50]. Overall, the enzyme activity of the samples follow the order CHA N Si–HA N HA. The higher enzyme activity indicates a higher number of cells differentiated into the osteoblast phenotype. Cells on the CHA sample started to differentiate at day 7 of culture. However, after 21 days of culture, all the samples showed a tendency of an increased

Fig. 8. ALP activity of cell cultured on HA, Si–HA and CHA samples.

Fig. 9. Bone nodule-like formation on (a) HA, (b) Si–HA, (c) CHA samples.

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Acknowledgments This study was supported by the AUN/SEED-Net project under the Japan International Cooperation Agency (JICA), HIR grant number H-16001-00-D000033 and PPP grant number RP011D-13AET.

References

Fig. 10. Calcium quantification detected on HA, Si–HA and CHA samples after 21 days of culture.

in ALP. This is attributed to the late stage of culture as the osteoblast cells advancing to the maturation and production of mineralized matrix [51]. Fig. 9 shows the Alizarin red S staining results of different samples after 14 days of culture. As can be seen, the areas of mineralized nodules detected on Si–HA and CHA surfaces were larger than that on the HA surfaces. Therefore, calcium deposition on Si–HA and CHA samples was supposed to be earlier and more efficient than that of HA sample. However, there was not much difference observed on the bone nodule obtained from the Si–HA and CHA samples. The calcium concentration of the sample after 14 days of culture is shown in Fig. 10. The result is in agreement with the bone nodule observation shown in Fig. 9. This indicates that there is more calcium present on the Si–HA and CHA samples when compared to the pure HA sample. Based on the cell study results, it can be deduce that the substitution of carbonate and silicate in HA results in better biological performance of MC3T3-E1 osteoblast cell culture while the carbonate substitution showed the best cell response. Adhesion, proliferation, differentiation and mineralization processes of these ionic-substituted HA ceramics are good indicators of the cell response that could be expected when a biomaterial is used in vivo [15].

4. Conclusions Si–HA and CHA were successfully prepared by a precipitation method and characterized by XRD, TEM, FTIR, XRF, ICP/AES, and CHN analyses. The physico-chemical properties of the synthetic materials were investigated to determine the role of different ionic substitutions in the HA lattice. XRF and FTIR results confirmed the presence of the substituted species of silicate and carbonate in the lattice. Substitution causes lattice disorder and decreases crystallite size, as well as suppresses grain coarsening of HA after heat treatment at 1200 °C. SiO4− 4 and CO2− ion substitution increases ion release from the substituted 3 HA, resulting in better dissolution of CHA and Si–HA compared with pure HA. Meanwhile, silicate substitution induced better solubility than carbonate substitution. Based on cellular tests using MC3T3-E1 osteoblast-like cells, it was found that the substitution of SiO44 − and ions enhances cell proliferation, differentiation as well as bone CO2− 3 mineralization. However, carbonate substitution has been observed to exhibit an overall better biological response than silicate substitution. This study has demonstrated that the properties of HA could be tailored through the combination of carbonate and silicate co-substitution to achieve excellent biocompatibility for orthopedic applications.

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