Bioglass-assisted preparation of network sodium titanate bioceramics

Abstract

The objective of this work is to prepare new network sodium titanates ceramics (N-STC) via a common ceramics fabrication method and to evaluate the influence of a bioglass additive and sintering temperature on the mechanical properties and biological activity of N-STC. Samples with different concentrations of bioglass were prepared at 900–1050 °C. The crystal structure and morphology of the materials were characterized by powder X-ray diffractometry (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). The porosity was investigated by a conventional liquid substitution method. Hardness and fracture toughness were determined by using the Vickers indentation method. As a preliminary biological evaluation, vitro simulated body fluid (SBF) tests were carried out using the Kokubo method. It was found that the bioglass additive and the sintering temperatures significantly affect the structure and performance of N-STC. The best performance was obtained by sintering the network sodium titanates (ST) together with 15 wt% bioglass at 950 °C. The porosity of the material is 82.75% and bending strength is 40.2 MPa. These N-STC samples with low cytotoxicity were immersed in SBF for 7 days, and covered by calcium titanate (CT) together with hydroxyapatite (HAP), indicating the potential application of this material in tissue engineering.

---

1. Introduction

Sodium titanate (ST) is of great interest for potential applications such as photocatalysts,¹ an insertion material for lithium ions,² humidity materials,³ gas sensitive materials,⁴ and ion exchange materials,⁵ etc. In recent years, sodium titanate ceramics (STC) have been investigated as a new type of bioceramics. Becker et al.⁶ obtained STC by a sol–gel method. This material exhibited excellent bioactivity when exposed to simulated body fluid (SBF). After being treated in NaOH, the sintered STC could be used as a ceramic implant with bone-bonding ability. Various technologies, such as the solid-state reaction method, wet chemical method, and laser irradiation have been used to fabricate STC. Holzinger et al.⁷ carried out solid-state reaction to fabricate ST from stoichiometric amounts of Na₂CO₃ and TiO₂. A template-free direct laser irradiation technology was also used to fabricate micropatterned film with nano-porous ST.⁸ It has been reported that nanocrystalline bioceramics with high purity and appropriate structural characteristics have better functions in biomedical applications. The sol–gel method has been used successfully for synthesizing ST powder. Baliteau et al.⁹ directly prepared STs mixtures (Na₂Ti₃O₇/Na₂Ti₆O₁₃) via a sol–gel reaction and investigated electrical properties of each ceramic composite. The hydrothermal method is powerful for obtaining nanoparticles, nanotubes, nanobelts, and nanowires, allowing large-scale production of the materials only by one single step. Through a simple hydrothermal method, Wang¹⁰ and Aziz¹¹ fabricated ST with various morphologies for the application of a new electrode for high energy density supercapacitors. Zhang et al.¹² also synthesized ST nanorods via a hydrothermal reaction in alkaline solution.

It is common sense that the liquid-phase sintering route by using glasses as sintering additives is an effective way to improve the mechanical properties of bioceramics.^13–16 The bioglass additives can promote sintering process via additional diffusion mechanisms such as dissolution/precipitation, capillary forces, and particle rearrangement owing to formation of liquid-phase at low sintering temperatures. The porous bioceramic with high mechanical properties plays an important role in determining the degree and rate of bone ingrowth. The interconnectivity, porosity, and pore size are three key parameters.¹⁷ To favor cellular and vascular penetration which ensures bone ingrowth inside pores, the porous architecture of material must be well controlled, and particularly the pores must interconnects closely. To obtain a ceramic body with high porosity, various methods have been used. Adding the pore-forming combusting organic material (e.g., polyethylene glycol (PEG), rosin^18,19) during sintering process can efficiently generate free spaces and voids inside the resulting body. Other compounds can also be used as pore-forming agents. Fang et al.²⁰ pointed out that CaSO₄ can decompose into CaO with smaller size in air above 1220 °C and then the pores formed. Lee et al.²¹ fabricated hydroxyapatite (HAP) bioceramics with the porosity of 75% and compressive strength of 0.94 MPa. The camphene-based freeze casting method produced three-dimensionally interconnected pore channels by removing the frozen vehicle network via sublimation. Recently, Huang et al.²² obtained CaSiO₃ bioceramics with interconnected lamellar pores and macro-channels of pores by an ice/fiber-templated method.

Up to now, to meet with the growing demands of bioactive materials, bioceramics were extensively investigated for the reconstruction and repair of diseased or damaged part of human body. Li et al.²³ studied the influence of fluoride additions on biological and mechanical properties of glass-ceramics and found that it suffers from low bending strength of 26.74 Mpa. β-TCP bioceramics were also widely used as bone replacements in the fields of oral and plastic surgery. However, it is well known that the low mechanical behavior of β-TCP bioceramics is the main limitation in load-bearing applications.²⁴ In order to promote the progress in the implant materials, for which they must possess high mechanical properties escaping the breakage in collision, high porosity allowing 3-dementional proliferation for cells, and good biological activity determining the degree and rate of tissues ingrowth, we should pay more attention on the synthetic methods obtaining well-behaved materials. Herein, we design the synthetic project to prepare ST bioceramics with high porosity, excellent interconnection, good biological activity, and high mechanical properties. Effects of bioglass additive and sintering temperature on properties of N-STC were investigated here. And this study could also provide a more comprehensive understanding for the mechanism of the porous materials formation and help to determine the optimum bioceramics for use as implant materials that must have enhanced mechanical performance and better biochemical properties. This simple and efficient approach can be considered to be a facile, controllable, and generic way for preparing advanced bioceramics.

2. Experimental procedure

2.1. Materials

All chemicals used in this work are of analytic grade and without further purification. Titania P-25 (TiO₂, ca. 80% anatase, and 20% rutile) was purchased from China National Medicines Co., Ltd. Calcium carbonate (CaCO₃), calcium phosphate dibasic dehydrate (CaHPO₄·2H₂O), magnesia (MgO), boron trioxide (B₂O₃), silicate oxide (SiO₂), aluminum oxide (Al₂O₃), sodium carbonate (Na₂CO₃), and sodium hydroxide (NaOH) are all purchased from Sinopharm Chemical Reagent Co., Ltd.

2.2. Preparation of Na₂Ti₃O₇ nanobelts and bioglass powders

Na₂Ti₃O₇ nanobelts were prepared via a hydrothermal process in a concentrated NaOH aqueous solution. 0.2 g of P-25 powder was dispersed in 25 mL 10 M NaOH aqueous solution. After being sonicated for 15 min, the solution was transferred into a 25 mL Teflon autoclave, heated at 180 °C for 72 h, and cooled to room temperature in air. The obtained powder was washed repeatedly with deionized water until the pH value of washing water reached 7. Subsequently, the samples were dried at 120 °C for 12 h in air.

The bioglass was prepared by the melt-quenching method. The chemical composition of the glass is 7.0 SiO₂-43.4 B₂O₃-3.7 P₂O₅-3.2 Al₂O₃-35.0 CaO-7.7 Na₂O (wt%). The raw materials were extensively mixed, placed in a platinum crucible, and heated to 1150 °C for 2 h. Subsequently the melt was rapidly quenched in water. The obtained sample was dried, milled with dehydrated alcohol, and sieved to obtain particles with size less than 64 μm.

2.3. Fabrication of N-STC

The bioglass powders were mixed with ST nanobelts by planetary high-energy ball for 20 min in alcohol using zirconia balls as medium. The green body of N-STC with the dimension of Φ: 30 mm × 10 mm was formed by slurry suction filter molding process. After formation, the green body was dried in the oven at 110 °C for 24 h. The obtained samples were sintered at the temperatures between 900 and 1050 °C for 1 h with a heating rate of 4 °C min⁻¹.

2.4. Tests for bioactivity in vitro

The bioactivity of the obtained N-STC samples was evaluated by examining the calcium titanate (CT) and HAP formation on the surface of samples in SBF. The SBF was prepared according to the method by proposed Kokubo.²⁵

0.1 g N-STC samples were accurately weighed and soaked in the SBF of 100 mL, and then were put into constant temperature control rotary table with speed of 170 rpm at 37 ± 0.5 °C. The samples in SBF were filtered for a certain time, and the ion concentration such as Ca, P, and pH value of the filtrates were determined. The samples were refreshed by acetone and then dried at room temperature in the air.

2.5. Mechanical properties

The bending strength was measured by three-point bending method on squareness specimens in RGD-5 type electronic tensile machine. This test was repeated for eight times, and the test results were the means of the values, not including the minimum and maximum. The strengths of samples were calculated by using the below formula:

S = 3PL/2bd²

where P is the fracture load (kg), L is the span length (mm), b is the width of the samples (mm) and d is the thickness of the samples (mm).

2.6. Characterization

Scanning electron microscopy (SEM) observation of samples was performed on a field-emission scanning electron microscope (SEM; Shimadzu EPMA-8705QHII, Japan) with energy dispersive X-ray (EDX) analysis. The open porosity was determined by the Archimedian method. The crystallographic structure and composition were characterized by X-ray powder diffraction (XRD) (Bruker D8) with Cu-Kα radiation (λ = 0.15406 nm). High resolution transmission electron microscope (HRTEM) images were obtained through a JEOL JEM 2100 microscope. The bending strength of the sintered samples was measured using a mechanical testing machine (Shimadzu AG-5kN, Japan). FT-IR analysis was done by using a Varian FTS 1000 spectrometer scimitar series (Varian Inc, Co., Australia). The samples were prepared by pelletized with potassium bromide. Cytotoxicity assay was performed on the total internal reflection fluorescence microscopy Olympus IX81. All measurements were performed at room temperature.

3. Results and discussion

3.1. Phase compositions

The XRD patterns of the bioglass, Na₂Ti₃O₇ nanobelts, and the N-STC samples sintered at 900 °C, 950 °C, 1000 °C, and 1050 °C, respectively, are shown in Fig. 1. The results reveal that bioglass powders show a typical diffraction pattern of amorphous characteristics. The powder prepared via an alkaline hydrothermal is ST nanobelts (Na₂Ti₃O₇),²⁶ The lattice parameters are a = 9.128 Å, b = 3.803 Å, c = 8.562 Å, and β = 101.60°. For sintered N-STC samples, when sintering temperature was raised to 900 °C from room temperature, the compound was changed into Na₂Ti₆O₁₃ whose lattice parameters are a = 15.120 Å, b = 3.738 Å, and c = 9.160 Å. No significant difference is observed except the intensity of diffraction peak for samples sintered at 950 °C, indicating that no new crystalline phase was formed and the degree of crystallinity of N-STC sintered at 950 °C is better than that sintered at 900 °C. However, when the sintering temperature was further increased, several peaks corresponding to rutile TiO₂ (PDF: 65-1119) appeared, the lattice parameters are a = 4.588 Å, b = 4.588 Å, and c = 2.967 Å, which results from the decomposition of Na₂Ti₆O₁₃ to Na₂O and rutile TiO₂. There is a little of Na₂Ti₆O₁₃ in the sample that sintered at 1050 °C. The decomposition can be described with the following equations:

2Na₂Ti₃O₇ → Na₂Ti₆O₁₃ + Na₂O↑

Na₂Ti₆O₁₃ → 6TiO₂ + Na₂O↑

Fig. 1 XRD patterns of the bioglass, Na₂Ti₃O₇ nanobelts, and N-STC samples sintered at different temperatures. (a) N-STC-1050 °C, (b) N-STC-1000 °C, (c) N-STC-950 °C, (d) N-STC-900 °C (e) Na₂Ti₃O₇ nanobelts, (f) bioglass (●) Na₂Ti₃O₇ (★) Na₂Ti₆O₁₃ (▼) TiO₂.

Therefore, the optimized sintering temperature of 950 °C for N-STC is adopted in this work.

3.2. Microstructure analysis

Fig. 2 shows the microstructure of Na₂Ti₃O₇ nanobelts, ST nanorods and N-STC sintered at 950 °C. It is clear that Na₂Ti₃O₇ nanobelts and calcined Na₂Ti₃O₇ nanobelts at 950 °C possess distinctive characteristics in shape. The SEM image of Na₂Ti₃O₇ nanobelts is shown in Fig. 2a, revealing that the samples possess uniform belt-like microstructures. The as-prepared nanobelts are 100–200 nm in width, 20–50 nm in thickness. The insert graph is high resolution TEM image for Na₂Ti₃O₇ nanobelts. Fig. 2b shows the morphology of ST nanorods which are sintered at 950 °C. As we can see, Na₂Ti₃O₇ nanobelts have become nanorods after sintered at 950 °C, and the nanorods are shorter in length and larger in diameter than that of Na₂Ti₃O₇ nanobelts. Fig. 2c shows the morphology of N-STC obtained by mixing Na₂Ti₃O₇ nanobelts with bioglass powders and consequently sintering at 950 °C. It is clear that N-STC has network structure and possesses high porosity. Fig. 2d is a larger view of Fig. 2c. It shows that the nanorods of N-STC connect with each other so that the network structure is formed with evenly distributed interconnected pores.

Fig. 2 SEM images of as-prepared Na₂Ti₃O₇ nanobelts (a), ST nanorods calcined at 950 °C (b), the cross section of N-STC samples (c), and the high magnification figure of N-STC samples (d).

3.3. Mechanical properties

In order to optimize the best structure and bioactivity, we investigated the effects of bioglass concentrations and the sintering temperature on the fabricating process. Table 1 summarizes the porosity and bending strength of the samples with different concentrations of bioglass at different sintering temperature. The porosity of the samples fabricated at 950 °C with 5 wt%, 15 wt%, 30 wt%, 40 wt% bioglass are 90.21%, 82.75%, 75.23% and 40.3%, respectively. The distinct porosity of the samples due to different bioglass concentration could result in quite different mechanical properties. Obviously, the bending strength of the samples sintered at 950 °C increases with increasing the amount of bioglass, namely, 34.8 MPa for 5 wt%, 40.2 MPa for 15 wt%, and 51.0 MPa for 30 wt% bioglass, respectively. However, if the content of bioglass increased to 40 wt%, the bending strength decreased to 37.8 MPa. That is to say, the porosity progressively diminishes and the mechanical strength progressively improves as the bioglass content increases. At the same time, the mechanical strength of the sample sintered at 950 °C for 1 h is 65% higher than that of the samples sintered at 900 °C.

Table 1 The porosity and mechanical strength of the samples

Bioglass content (wt%)	Sintering temperature (°C)	Porosity (%)	Bending strength (MPa)
40	1000	32.5	38.2
30	1000	65.56	61.4
15	1000	72.78	56.2
5	1000	83.95	45.3
40	950	40.3	37.8
30	950	75.23	51.0
15	950	82.75	40.2
5	950	90.21	34.8
10	900	86.23	24.4

---

The porosity and bending strength are controlled by changing the fabrication process, such as molding process, bioglass content, and sintering temperature, etc. To address this problem, slurry suction filter molding process was adopted, which could enhance the uniform distribution of bioglass and nanobelts. Fig. 3 shows the SEM images of the N-STC samples sintered at 950 °C with different content of bioglass. Fig. 3a and b shows the SEM graphs of N-STC samples with 15% bioglass content. As we can see, there are a lot of bioglass bead in ST rods combination position including point combination, line combination, and plane combination, which helps to promote the formation of network structure, like metal welding. That is to say, the connection of ST rods with each other is facilitated by the presence of bioglass liquid phase. This may be due to the mechanical bite force and the liquid phase surface tension during the suction filter molding process. However, with increasing the amount of bioglass addition, a clear decreasing trend was observed in mesh size and porosity, even destroying the network structure. The excess of bioglass leads to rods orientation arrangement (Fig. 3c and d). The formation of continuous and strong network between the ST rods could be the main reason accounting for the increased mechanical strength. Therefore, appropriate bioglass content is a key factor.

Fig. 3 The SEM graphs of the N-STC samples sintered at 950 °C. (a) Bioglass content of 15% (b) bioglass content of 15% at high magnification (c) bioglass content of 30% (d) bioglass content of 40%.

Sintering temperature is one of the key factors to control the network structure. As sintering temperature was increased from 950 °C to 1000 °C, the ST green body contracts and the mechanical strength increases for all resulting samples. The green body contraction of N-STC differs from those of common ceramics, which are mainly due to the capillary force and the particle rearrangement. Fig. 4 shows the principle of the formation of N-STC. ST changed from belts to hexagonal rims and then to cylinder and the bioglass melting to liquid phase with the increase of sintering temperature. At the same time, ST rods tend to orientation arrangement with the action of bioglass liquid, which results in the forming of small mesh and improvement of the mechanical strength. With increasing the sintering temperature, more bioglass bead melted in point combination, line combination, and plane combination, welding the junctions and connecting the ST rods into a network. The circular structure can be found in all rods junction at high temperature, which indicated that this network structure is very strong and has bigger surface area than other structures.

Fig. 4 Schematic of the formation of N-STC (a) Na₂Ti₃O₇ nanobelts (b) N-STC sintered at 900 °C (c) N-STC sintered at 950 °C (d) N-STC sintered at 1000 °C.

3.4. Evaluation of the bioactivity in vitro

Ion concentration of SBF and blood plasma are listed in Table 2. SBF reproduces the internal environment of physiological fluids containing the same ionic composition and therefore SBF-based in vitro studies simulate the inorganic reactions taking place once the material is implanted into the body.^27,28

Table 2 Ion concentration of SBF and human blood plasma

	Ion concentration (mg L^−1)
	Na^+	K^+	Mg^2+	Ca^2+	Cl^−	HCO3^2−	HPO4^2−	SO4^2−
SBF	142.0	5.0	1.5	2.5	147.9	4.2	1.0	0.5
Blood plasma	142.0	5.0	1.5	2.5	103.0	4.2	1.0	0.5

---

Fig. 5a shows the evolution of pH values of the fluid after ST nanobelts, N-STC, and ST porous ceramics were immersed in the SBF. The pH of SBF is 7.4. However, an obvious variation of pH values of the samples is observed after the ceramics were immersed into SBF. The pH values of the samples increases with immersing time. The pH value increases very quickly at the beginning and then reaches a plateau after certain immersing time. After being immersed in SBF for 80 h, the pH values of the samples are 8.5, 8.2, and 8.0 for ST fibers, N-STC, and ST porous ceramics, respectively. The concentration of Ca and P in SBF vs. time curves is shown in Fig. 5b. As we can see, the concentrations of Ca and P decrease with immersing time. Moreover, decrease of ionic concentrations is more pronounced for N-STC samples due to the higher degree of bioactivity. In conclusion, this phenomenon indicates that 3D porous network can increase biological activity than that of ST porous ceramics. Fig. 5c and d shows the SEM images of the surfaces of N-STC with the content of bioglass of 15% before and after soaked in SBF for 7 days. It is clearly that some particles are formed after immersed in SBF for 7 days. And these particles with size ranging from 80 to 300 nm are homogeneously dispersed on the on the surface of N-STC.

Fig. 5 The evolution of pH (a) and Ca²⁺, P⁵⁺ (b) in SBF and SEM micrographs of N-STC before (c) and after (d) immersion in SBF for 7 days.

Fig. 6 shows the XRD patterns of N-STC with the content of bioglass of 15% before and after immersed in SBF for 7 days. Before soaking, the XRD patterns of the sample corresponded to the phases Na₂Ti₆O₁₃. After soaking for 7 days, the XRD patterns of sample show the diffraction peaks at 2θ values of 33.25, 47.62, 59.43, and 69.13° corresponding to the pure CaTiO₃ with perovskite phase (JCPDS: 42-0423, CT). Meanwhile, the weak but clear peaks of HAP are found at 2θ values of 31.77, 32.19, 39.81, 46.71, and 54.44° corresponding to the presence of HAP. So we conclude that the CT and HAP are formed on the surface of N-STC.

Fig. 6 The XRD patterns of N-STC before (b) and after (a) immersion in SBF for 7 days. ● Na₂Ti₆O₁₃, ▼ perovskite CaTiO₃, ★ Ca₁₀(PO₄)₆(OH)₂.

There are three points for explaining the reason why the CT and HAP can be easily formed on the surface of N-STC. Firstly, the Ti–OH functional group can be formed on the surface of N-STC. When N-STC immersed in SBF, the N-STC exchanges the Na⁺ ions from the surface with the H₃O⁺ ions in the fluid to form Ti–OH groups on its surface. And consequently, these Ti–OH groups incorporated with the calcium ions from the fluid to form CT. The process promoting the nucleation of CT can be described as followed reaction equations:⁶

In SBF: Ti–O–Na + H₂O → Ti–OH + NaOH

Ti–O–Na + H₃O⁺ → Ti–OH + Na⁺ + H₂O

Ti(OH)₄ + Ca²⁺ + 2OH⁻ → CaTiO₃ + 3H₂O

The CT can significantly speed up the formation of bone calcium phosphate,²⁹ which has an enhanced effect on the proliferation and differentiation of osteoblasts under isolated conditions and improved alkaline phosphatase activity and relative protein content. It is more conducive to osteoblasts adhesion and proliferation without toxic effects on osteoblasts.³⁰

After a long soaking period, Na⁺ leads to the increase of the OH⁻ concentration and consequently an alkaline environment. This alkaline environment can also promote the nucleation of HAP. Further, the CT can provide a positive surface charge that interacts with the negatively charged phosphate ions in the fluid to form HAP on the surface of CT. HAP can directly bond to the bone and is regarded as the most stable material for hard tissue replacement implants.³¹ The process can be described as the following equation:

10Ca²⁺ + 6PO₄³⁻ + 2OH⁻ → Ca₁₀(PO₄)₆(OH)₂

To further elucidate the composition in the surface layer, the samples are characterized by EDX and FT-IR analysis. Fig. 7 shows the EDX spectroscopy determining the elemental distribution and composition. We selected a small area of a single ST rod in N-STC immersed in SBF for 7 days as the analysis region as shown in Fig. 7(a). As displayed in Fig. 7(b)–(e), Ti, P, O, and Ca elements in the testing region with even distribution are detected in the surface layer. All elements are homogeneously distributed in the covering layer. Fig. 8 shows the FT-IR spectra of N-STC samples before and after immersed in SBF for 7 days. As observed, the absorption bands below 700 cm⁻¹ are assigned to Ti–O of ST and CT. New absorption band 972 cm⁻¹ can be recognized at 972 cm⁻¹ after immersed, which was caused by the symmetric stretching vibration of PO₄³⁻ group and indicated the formation of HAP. The HAP also generated the characteristic –OH band at 3385 cm⁻¹.

Fig. 7 The SEM of N-STC with the content of bioglass of 15% immersed in SBF for 7 days (a) and corresponding EDS maps. (b) Ti K map, (c) O K map, (d) P K map, (e) Ca K map.

Fig. 8 FT-IR spectra of N-STC samples before and after immersed in SBF for 7 days.

Secondly, the surface of ST is negatively charged. The ST nanobelts were fabricated by strong alkaline hydrothermal method, which resulted in a large amount of hydroxyl groups on the surface of nanobelts. The hydrogen ions lose easily due to the strong polarization effect of titanium ions, which decreases the number of hydroxyl groups on surfaces:

–Ti–O–H ⇔ –Ti–O⁻ + H⁺

Fig. 9 shows the evolution of Zeta potential by varying pH values. Isoelectric points were 4.2 and 6.3 for the ST nanobelts and ST rods, respectively. As the pH value is beyond 6.5, ST (Na₂Ti₃O₇ and Na₂Ti₆O₁₃) is negatively charged and the Zeta potential is inversely proportional to pH value. When pH is higher than 8, the Zeta potential is nearly unchanged with a minimum of −71.6 mV for Na₂Ti₃O₇ and −63.07 mV for Na₂Ti₆O₁₃, respectively. The ST particles are negatively charged due to hydroxyl groups on the surface of the particles under the lower pH value solution. However, ST particles have negatively charged because of losing hydrogen ions from surface of particles under high pH value solution. A chemical reaction occurs on the surface of particles, which can be described by the following equation.

–Ti–OH + H⁺ ⇔ –Ti–OH⁺₂

–Ti–OH + OH⁻ ⇔ –Ti–O⁻ + H₂O

Fig. 9 Zeta potential of before and after calcinating ST nanobelts.

The Zeta potential vs. pH curves move rightward after ST nanobelts were sintered, as shown in Fig. 9; the isoelectric pH value changes from 4.2 in ST nanobelts to 6.3 in calcined ST nanobelts, which results in the formation of negatively-charged surface of ST nanobelts and rods under neutral/alkaline environment. Because CT has positively charged surface, CT and ST could attract each other through electrostatic interaction, and PO₄³⁻ is preferentially adsorbed from solution to the surface of CT and followed by the adsorption of Ca²⁺. Then the nucleation of HAP occurred as the adsorption of ions reached a certain level. CT becomes an intermediate between N-STC and HAP.

Thirdly, the surface of CT is rough and N-STC has porous structure. The pores in N-STC improve the growth of bone tissue, which results in biocompatibility between the human-body and artificial materials. The rough surface and porous structure lead to the increase of surface areas, better performance in solubility, and accumulation of ions around the surface. The pores also retard the diffusion of ions. As a result, nucleation of CT preferentially occurred in this region. More Ca²⁺, TiO₃²⁻, and PO₄³⁻ ions are required for growth of the nuclei, and these Ca²⁺, TiO₃²⁻ and PO₄³⁻ ions diffuse from solution toward the nuclei. For bigger pores and cracks, Ca²⁺, TiO₃²⁻, and PO₄³⁻ ions can easily get through the concentration gradient region, which improves the formation of stable crystal layers.

3.5. Cytotoxicity assay

Fig. 10 shows the total internal reflection fluorescence images of SiHa cancer cell incubated with N-STC samples and determined by 0.4% trypan blue staining. As can be seen from Fig. 10, there were no killed SiHa cells after incubated with N-STC samples for 4, 8, and 12 h, respectively. When the incubation time reached to 24 h, only a few cells were killed and displayed a blue color, implying a low cytotoxicity. These results indicated the potential of N-STC samples as robust implant materials.

Fig. 10 Images of SiHa cell incubated with N-STC samples for 0 h (a), 4 h (b), 8 h (c), 12 h (d), 24 h (e) and (f).

4. Conclusions

In this study, the N-STC with evenly distributed and interconnected pores was successfully fabricated by a simple approach. We demonstrated that the sintering temperature and bioglass additive significantly affect the pore distribution, the size of pores, and the strength of the N-STC. The bending strength of the samples with 15 wt% bioglass addition and sintered at 950 °C (about 82.75% porosity) reached 40.2 MPa. When soaked in SBF, the CT and HAP were deposited on the surface of the NSTC samples. The N-STC samples possess excellent porous structure, mechanical strength, bioactivity, and low cytotoxicity, making itself a potential candidate for bone tissue engineering applications.

Acknowledgements

We express our appreciation for the financial support, National Basic Research Program of China (no. 2012CB932302).

References

1. L. Shi, L. X. Cao, W. Liu, G. Su, R. J. Gao and Y. L. Zhao, Ceram. Int., 2014, 40, 4717–4723.
2. R. Dominko, E. Baudrin, P. Umek, D. Arcon, M. Gaberscek and J. Jamnik, Electrochem. Commun., 2006, 8, 673–677.
3. Y. X. Li, M. H. Li, F. Du, Y. Y. Zhang, W. Y. Fu and H. B. Yang, J. Jilin Univ., Earth Sci. Ed., 2010, 48, 121–123.
4. J. R. Salgado, E. Djurado and P. Fabry, J. Eur. Ceram. Soc., 2004, 24, 2477–2483.
5. J. Yang, D. Li and X. Wang, J. Mater. Sci., 2003, 38, 2907–2911.
6. I. Becker, I. Hofmann and F. A. Muller, J. Eur. Ceram. Soc., 2007, 27, 4547–4553.
7. M. Holzinger, J. Maier and W. Sitte, Solid State Ionics, 1997, 94, 217–225.
8. P. Yu, C. Ning, G. Tan, Y. Zhang, J. Liao, J. Sun, W. Peng, M. Zhong, Z. Yu and G. Ni, Surf. Coat. Technol., 2013, 235, 267–272.
9. S. Baliteau, A. L. Sauvet, C. Lopez and P. Fabry, Solid State Ionics, 2007, 178, 1517–1522.
10. L. Wang, T. Zhang, Q. Qi, J. Hu, Y. Zeng and G. Lu, Mater. Lett., 2009, 63, 903–904.
11. R. A. Aziz, I. I. Misnona, K. F. Chong, M. M. Yusoff and R. Jose, Electrochim. Acta, 2013, 113, 141–148.
12. H. Zhang, X. P. Gao, G. R. Li, T. Y. Yan and H. Y. Zhu, Electrochim. Acta, 2008, 53, 7061–7068.
13. A. F. Habibe, L. D. Maeda, R. C. Souza, M. J. R. Barboza, J. K. M. F. Daguano, S. O. Rogero and C. Santos, Mater. Sci. Eng., C, 2009, 29, 1959–1964.
14. S. I. Roohani-Esfahania, S. Nouri-Khorasani, Z. F. Lu, R. C. Appleyard and H. Zreiqat, Acta Biomater., 2011, 7, 1307–1318.
15. D. Bellucci, A. Sola, M. Gazzarri, F. Chiellini and V. Cannillo, Mater. Sci. Eng., C, 2013, 33, 1091–1101.
16. H. Ghomi, M. H. Fathi and H. Edris, Mater. Res. Bull., 2012, 47, 3523–3532.
17. K. Lin, L. Chen, H. Qu, J. Lu and J. Chang, Ceram. Int., 2011, 37, 2397–2403.
18. K. Lin, J. Chang, Z. Liu, Y. Zeng and R. Shen, J. Eur. Ceram. Soc., 2009, 29, 2937–2943.
19. H. Dai, X. Wang, Y. Han, X. Jiang and S. Li, J. Mater. Sci. Technol., 2011, 27, 431–436.
20. Z. Fang, Q. Feng and R. Tan, Ceram. Int., 2013, 39, 8847–8852.
21. E. Lee, Y. Koh, B. Yoon, H. Kim and H. Kim, Mater. Lett., 2007, 61, 2270–2273.
22. Z. Huang, K. Zhou, D. Lei, Z. Li, Y. Zhang and D. Zhang, Ceram. Int., 2013, 39, 6035–6040.
23. H. C. Li, D. G. Wang, J. H. Hu and C. Z. Chen, Mater. Sci. Eng., C, 2014, 35, 171–178.
24. K. Lin, J. Chang, J. Lu, W. Wu and Y. Zeng, Ceram. Int., 2007, 33, 979–985.
25. T. Kokubo, H. M. Kim, T. Nakamura and F. Miyaji, J. Am. Ceram. Soc., 1996, 79, 1127–1129.
26. J. J. Lin, J. X. Shen, R. J. Wang, J. J. Cui, W. J. Zhou, P. G. Hu, D. Liu, H. Liu, J. Y. Wang, R. I. Boughton and Y. Z. Yue, J. Mater. Chem., 2011, 21, 5106–5113.
27. S. S. R. Aryal, R. Bhattarai, M. S. Khil, D. Lee and H. Y. Kim, Mater. Sci. Eng., A, 2006, 426, 202–207.
28. T. Kokubo and H. Takadama, Biomaterials, 2006, 27, 2907–2915.
29. F. Zhang, S. G. Chen and C. Lin, Appl. Surf. Sci., 2011, 257, 3092–3096.
30. M. Sun, X. Tang and Z. Wu, Acta Univ. Med. Anhui, 2013, 48, 1324–1327.
31. M. N. Rahaman, D. E. Day, B. Sonny Bal, Q. Fu, S. B. Jung, L. F. Bonewald and A. P. Tomsia, Acta Biomater., 2011, 7, 2355–2373.

---