1. Introduction
Continuous human evolution related with materials and technology has led to major improvements, such as considerable prolonged human life expectancy. Some consequences from such development are related with diseases of an ageing population or injuries from trauma accidents, which render strategies to repair, replace, and even regenerate damaged osseous tissue necessary. Tissue engineering researches new biomaterials that could meet these demands. In this group of materials, ceramics have received more attention given their good properties: for example, their ability to favor bone adhesion and bone ingrowth when implanted into animal and human bodies. The use of calcium phosphates is very popular because their compositions are similar to the mineral bone part [
1,
2], but P
2O
5-free calcium silicates have also been able to develop a hydroxyapatite layer on their surfaces [
3,
4], which is key to bone–implant interactions.
The success of a biomaterial to be used as an implant is related not only to its intrinsic properties but also to the required hierarchical structure. In addition to surface connections to bone tissue, it must act as a temporary template by stimulating and guiding bone natural regeneration, and it degrades at a rate that guarantees mechanical support [
5]. The porous structure has been developed by different methods, such as the polymeric sponge method, the use of porogens, foaming techniques, 3D printing, and so on [
6,
7,
8,
9]. Although there is some controversy as to the ideal pore size [
10,
11,
12], it is clear that interconnection and the macro- and microporosity combination are critical factors to gain scaffolds capable of mimicking cancellous bone.
Another goal to achieve with this porous structure is to maintain mechanical support in the early bone regeneration stage. A good approach to overcome this problem has been devised by De Aza et al. by the in situ generation of a porous template by soaking a biphasic ceramic of CaSiO
3 (W)-Ca
3(PO
4)
2 (C
3P) of a eutectic composition in simulated body fluid (SBF), and by the dissolution and transformation of a dense ceramic into a porous one [
13,
14]. Nevertheless, when the ceramic was implanted in vivo, transformation only took place on the surface of the material at a depth of 200 µm because the pores left by the dissolution of the W phase in the eutectic structure were very narrow (~1.5 µm). Thus bone tissue was unable to completely osteointegrate the implant [
15].
However, by means of processing, it is possible in material science to adapt the bioceramic potential implant size, shape, and microstructure. To overcome this problem, in the present work we studied a hypereutectic composition that was rich in Ca
3(PO
4)
2 (the phase in which the material with a eutectic composition transforms pseudomorphically into hydroxyapatite (HA) after soaking in SBF) and with two grain sizes of CaSiO
3 powder because it is the biodegradable phase [
16]. We hypothesized that the new ceramic would develop in situ a structure with large enough pores for cells to colonize it completely and would consequently present total implant osteointegration in vivo.
In order to check our initial hypothesis, the new hypereutectic ceramic was tested in vitro in SBF. Adult human mesenchymal stem cells (ahMSCs) were used to study their interaction and biological response when in contact with the prepared ceramics to study their biocompatibility, attachment, adhesion, and proliferation.
2. Materials and Methods
2.1 Materials Preparation
The starting material included laboratory wollastonite and tricalcium phosphate synthesized by a solid-state reaction according to a procedure described in a previous study [
17,
18]. C
3P was milled in an attrition mill for 3 min using isopropanol and 20 mm diameter PSZ-zirconia balls. The obtained powder was dried at 60 °C/24 h and sieved through a 30 µm mesh.
The W material was ground in a tungsten carbide mill and sieved. The powder that was left in sieves between 63–100 µm and 100–150 µm was chosen to prepare the two ceramics.
The hypereutectic ceramic materials were prepared by a mixture of W and C3P at a ratio of 30:70 wt %, referred to hereafter as W-C3P 1 (63–100 µm of W) and W-C3P 2 (100–150 µm of W). The mixture was homogenized with acetone, and pressed into bars (~7 mm ∅ × 80 mm long), with isostatic pression at 200 MPa to form green compacts.
Green compacts were sintered at 1325, 1350, and 1375 °C/6 h in air at heating and cooling rates of 5 °C/min to determine the optimum sintering temperature. From the sintered bars, diamond machining was performed for the prepared characterization.
2.2. Ceramics Characterization
In order to study the densification degree of the sintered hypereutectic ceramics, bulk density was performed using Archimedes´s method and porosity by mercury porosimetry (Poremaster-60 GT, Quantachmore Inst., Boyton Beach, FL, USA). Five samples for each sintering temperature were studied and standard deviation was calculated.
X-ray diffraction (XRD) was used for the mineralogical analysis done on the 120 µm powder samples. A Bruker AXS D8-Advance powder X-ray diffractometer (Karlsruhe, Germany) by λCuKα1 radiation (1.5418 Å) and a secondary curved graphite monochromator were used for obtaining XRD patterns. Standards from the Joint Committee on Powder Diffraction Standards (JCPDS) database were used as a reference in relation to the obtained diffractograms, α-C3P (JCPD-09-0348) and α-CaSiO3 (JCPD-74-0874).
Confocal Raman measurements were taken using a 532 nm excitation laser (green laser) and a 20X objective lens. Raman spectra were recorded within a spectral range, which went from 100 to 1200 cm−1, and were analyzed by the Witec Control Plus software (Witec alpha − 300R, Witec, Ulm, Germany).
The obtained ceramic materials were embedded in a vacuum in an epoxy resin and were then progressively polished down to 0.1 μm by diamond paste and etched with acetic acid (1% concentration, 10 s). Ultrasonic bath with distilled water was used for cleaning and, after drying, they were palladium-coated for scanning electron microscopy (SEM, Hitachi S-3500N, Ibaraki, Japan), including a wavelength dispersive spectroscopy system (WDS, INCA-Oxford, UK).
A Zeta-20 3D imaging and metrology microscope (Zeta-20 Optical Profilometer, Zeta Instruments, San Jose, CA, EEUU) was used to acquire the height data presented herein. The 3D images of the surface texture were acquired at the 50X optical magnification.
The mechanical properties of the bioceramics were studied using the Brazilian test, namely, the diametrical compression of discs test (DCDT). Circular discs were used with a ratio diameter/thickness (D/T) of ~0.3, and they were placed between two stainless steel loading plates with their faces perpendicular to the loading plates in a universal testing machine (Microtest, Spain) The load was applied with a displacement rate of 5 mm/min until the sample cracked. Fifteen specimens of WTCP 1 were tested as representative one, and strength was calculated using Equation (1):
where P is the applied load, D is the diameter, and T is the thickness of the disc.
For brittle materials, the weakest link theory [
19] was applied to analyze the distribution of strength values. The Weibull parameters were determined using the ENV-843-5 procedure through the Weibull function represented in Equation (2):
where P
f (σ) is the failure probability, σ
0 is the strength (failure probability 63.2%), and m is the Weibull modulus.
According to this standard, the probabilities of failure were calculated using Equation (3):
where N is the total number of specimens tested and n is the specimen rank in ascending order of failure stress.
2.3. SBF Test
An in vitro bioactivity analysis was done according to Kokubo’s protocol [
20]. The ceramic bars obtained after thermal treatment were cut as discs (7 mm in diameter, 1.5 mm thickness). Polystyrene bottles were used to immerse the discs hooked onto a nylon thread in SBF (50 ml), which has an ion concentration that is almost equal to human blood plasma, pH = 7.4. The samples immersed in the bottles with SBF were shaken in a water bath at 37 °C for pre-established periods of time. After the corresponding soaking period (up to four weeks), samples were removed from SBF, rinsed with deionized water, and dried at room temperature. Sample surfaces and cross-sections were examined by SEM and a microanalysis by WDS. For the cross-section analysis, 10 measures per sample were taken in five samples. The silicon, calcium, and phosphorus ion release profiles were determined for all the samples studied by inductively coupled plasma optical omission spectroscopy (ICP-OES, Perkin Elmer Optima 2000
TM, PerkinElmer Inc., Waltham, MA, USA).
For the transmission electron microscopy (TEM, JEM-2010 JEOL, Tokio, Japan) study, the precipitated layer on the scaffold’s surface was carefully removed using a razor blade and dispersed on the surface of a Petri dish with ethanol. After drying, a powder specimen was collected on a carbon-coated TEM copper grid (300 mesh). Electron beam transparent particles were chosen for the TEM examination by selected area diffraction (SAD).
2.4. Biocompatibility, Adhesion, and Proliferation Assays
The
ahMSCs used for the study were isolated as previously described [
21] according to the International Society of Cell Therapy (ISCT) criteria [
22] for their characterization. The assays were approved by the Institutional Ethical and Clinical Trials Committee (V. Arrixaca University Hospital of Murcia). Informed consents were obtained from all the volunteers.
The ahMSCs of third passage (P3), after cell expansion, were prepared for use in all the cellular studies. Biocompatibility, adhesion, and proliferation studies were performed by seeding a density of 5000 cells/cm2 of ahMSCs on top of the W-C3P 2 ceramics, which had been previously pre-incubated for 24 h with 500 μL of fetal bovine serum (FBS), and placed in 48-well culture plates with 500 µL of growth medium (GM) which consisted of Dulbecco’s Modified Eagle Medium (DMEM) with 10 % of FBS and 1 % penicillin-streptomycin (all from Sigma-Aldrich, St. Louis, MO, USA). All the samples were incubated at 37 °C in a 5 % CO2 atmosphere with 95 % relatively humidity. Medium was changed every three to four days.
Field emission scanning electron microscopy (FESEM, Merlin VP Compact, Carl Zeiss Microscopy S.L., Oberkochen, Germany) was used to examine cell adherence growth and morphology at 1, 7, 14, 21, and 28 days in GM.
2.5. Cell Viability Assay
The reagent Alamar Blue (Invitrogen, Carlsbad, CA, USA) is non-radioactive, safe, and stable in environment assays. This was why it was used to measure the cellular well-being and viability of ahMSCs on the W-C3P 2 surface. The top of the samples was seeded with 5000 cells/cm2 in a 48-well plate and incubated under the same conditions described above for 1, 7, 14, 21, and 28 days with GM.
Medium was replaced every three days and a positive control was also used by seeding ahMSC onto tissue-treated polystyrene (TCPS) culture plates. When each culture period ended, the corresponding medium was discarded and wells were washed twice with phosphate buffer solution (PBS). Each well was then filled with 500 μL of fresh medium containing 10 % (v/v) of Alamar Blue reagent to be incubated at 37 °C for 4 h. After the reaction time, 200 µL of dissolution were transferred to a 96-well plate and the fluorescence analysis was carried out directly in a Synergy MX ultraviolet visible (UV-Vis) reader (BioTek Instruments Inc., Winooski, VT, USA) at excitation and emission wavelengths of 560 nm and 590 nm, respectively. The results are reported as arbitrary units (a.u.)
All the experiments were performed at least in triplicate. Student’s t-test (p-value <0.05), was used in case of any possible significant differences between groups. All the data are reported as mean ± standard deviation (SD).
4. Discussion
In agreement with information on the pseudobinary system Ca
3(PO
4)
2-CaSiO
3 [
27], the phases formed at room temperature would be β-C
3P and wollastosnite-2M. However, the studied biphasic ceramic was constituted by two metastable crystalline phases as determined by XRD (
Figure 1), a SEM-WDS microanalysis (
Table 2), and µ-Raman (
Figure 2), namely, α-wollastonite and α-tricalcium phosphate. This discrepancy between the phases found in the real materials, and those in the thermodynamic equilibrium, can be explained by the kinetic effects on high-temperature reactions and transformations during cooling. Some authors have reported that small quantities of Si in C
3P could stabilize the high-temperature form of C
3P [
33,
34].
Both ceramics showed significantly increased shrinkage at 1350 °C (
Table 1) while sintering the green compacts, which can be explained by the existence of a liquid phase. In the materials with certain quantities of impurities, a liquid phase can form at temperatures lower than those of the equilibrium of pure materials [
35]. With WC
3P 1 and WC
3P 2, liquid might appear at temperatures as low as 1402 °C (invariant eutectic temperature of the pseudobinary system). As checked, the density of the ceramics sintered at 1375 °C was considerably higher than that of those sintered at 1350 °C, which is consistent with the presence of a glass phase. In any case, the amount of liquid formed at 1375 °C (slightly below the eutectic temperature) was so small that its presence in the microstructure of both ceramics was not detected, in which the grain boundaries between phases CaSiO
3 and Ca
3(PO
4)
2 were conventional and no other secondary phases could be related with liquid (
Figure 3 and
Figure 4).
The swallowing of the WC
3P 2 ceramic sintered at 1375 °C shown in
Figure 3c could be due to the formation of a low-viscosity calcium-silicophosphate liquid that penetrated open pores by capillarity, which led to their closing as observed by the large round pores presented in the microstructure.
The enhanced ceramic was that sintered at 1375 °C, which presented a homogeneous microstructure that consisted in phases α-W and α-C
3P. The α-C
3P grains presented microcracks through grains (
Figure 3a) and with a larger quantity in ceramic WC
3P 2 with a bigger wollastonite grain size (
Figure 3b).
The cause of the microcracks observed in both ceramics sintered at 1375 °C could be due to the stresses generated by the difference in the thermal expansion coefficient between α-W (10 × 10
−6 K
−1) and α-C
3P (60 × 10
−6 K
−1) [
36,
37,
38] and/or by the crystalline anisotropy of both phases. Fine microstructures would limit cracking because the critical cracking size could not be achieved. Cracks would form for the grains larger than the critical size (
Figure 3).
Diametrical compression test realized for WC
3P 1 as the representative one (
Figure 6) showed 32.5 MPa as the greatest strength value. Additional mechanical studies are in progress, but this result is promising for getting adequate mechanical support during bone regeneration in our future research goal: sinus lift procedures.
The ions measured by ICP from the SBF medium was done to control the possible ion release, which would occur in vivo and attempt to clarify the bioactivity mechanism. The concentration of the different ions present in SBF throughout the experimental time (
Figure 10) showed differences that could be attributed to the wollastonite particle size. For the
in vitro analysis by immersion in SBF, both ceramics displayed increased calcium and silicon ion concentrations during the first week of the experiment, which indicated the partial surface dissolution of the α-W phase. The degradation rate of a material in a living body should equal the bone tissue repair process for bone regeneration applications. In addition, the precipitate layer of apatite over the material is usually considered a sign of bioactivity, which was studied in-depth in this research study. The α-W phase showed a significantly higher dissolution rate than the α-C
3P (
Figure 10) like the SBF enriched in the Si ion. These processes occurred initially on the grain boundaries; subsequently, the diffusion process occurred intergranularly, and porosity was sometimes generated in the material in the first process stages (
Figure 7). This might be due to the constant solubility product of α-W (2.53 10
−8) being much bigger than that of α-C
3P (2.03 × 10
−29) [
39], which suggests a faster dissolution rate for the α-W phase. Behavior was slightly different in both ceramics: after four weeks, the presence of an amorphous silica phase prior to precipitate deposition was especially visible for sample WC
3P 2 (
Figure 9 bottom, and
Figure 11c). The presence of the amorphous silica layer, as reported in the scientific literature for such reactions [
35], was not detected in the SEM/WDS cross-section studies of WC
3P 1 (
Figure 8). This would explain the depletion of the Si ion in the SBF of WC
3P 2 at two weeks. It remained in the form of amorphous Si on the ceramic surface, whereas all the Si ions in WC
3P 1 passed to the medium, and no maximum value was reached during the experiment.
When analyzing P ion evolution, both ceramics showed a gradual decrease throughout the assay. This was probably due to the reaction of α-C
3P with calcium and hydroxyl ions at the medium, which became HA according to the following: 3(Ca
3(PO
4)
2)+ Ca + 2OH → Ca
10(PO
4)
6(OH)
2. As no increase in Ca had taken place after the first week, most Ca had to react with P at the interface, which would lead to the nucleation on the surface of the Ca–P powder particles (
Figure 7b), which transformed into a new Ca–P rich layer with time.
In the second immersion stage, a globular layer, which almost covered the entire surface, was observed for both ceramics (
Figure 7) as a Ca–P rich layer. This precipitate grew with exposure time (
Figure 8 and
Figure 9) and developed a continuous layer, which, with WC
3P 1, presented porosity (
Figure 7c (upper) and
Figure 8). Thickness was measured from one to four weeks (representative periods) with values of 14.62 μm for WC
3P 1 and 17.55 μm for WC
3P 2 (
Figure 8 and
Figure 9, respectively). The thickness data represent the average of 10 independent determinations. It was possible to control the dissolution of the bioceramic by adjusting the raw particle size of α-W. However, more in vivo degradation studies must be conducted to confirm the changes noted in the degradation properties of the WC
3P bioceramics.
These preliminary studies have shown that both cell viability and biocompatibility were good [
40]. However, it is necessary to do more assays to confirm this. Ions released from bioceramics, like calcium, silicon, and phosphate groups, can determine cell functions, such as adhesion and proliferation.
The FESEM images and alamarBlue results showed the biomaterials’ good biocompatibility. Between day 1 and 7, the number of cells was small and metabolic activity was low. This could be explained by the self-renewal capacity of ahMSCs, which is related to their growth and proliferation properties. The bioceramic surface is a new environment for cells, which need some time to adapt to it. However, metabolic activity and the number of the cells increased from day 14 until the end of the study time, when cells occupied the whole biomaterial’s surface.
With respect to the limitations of the experiment for obtaining a material with open porosity, it was observed that one of the constituent phases of the material (W) was not completely dissolved, because of which our starting hypothesis was not fulfilled. We are currently working on a new experimental design which will change the composition allowing wollastonite to be the major phase and a heat treatment that will favor nucleation but not growth, for a more homogeneous distribution of this phase that allows smaller grains of Wollastonite to be generated, but grains large enough for cell colonization.
5. Conclusions
In this study, dense WC3P (30–70 wt %) ceramics were successfully prepared by sintering WC3P powder with different particle size ratios at 1375 °C for 6 h. The obtained ceramic presented a homogeneous biphasic microstructure that included grains from a resorbable phase, α-W (30 wt %, and 63–100 µm or 100–150 μm), and another corresponding to α-C3P (70 wt %, ∼30 µm) that pseudomorphically transformed into a HA-like formation. Thermal expansion led to the microcracking of the α-C3P phase in the ceramic while cooling down from the sintering temperature.
Immersion in SBF allowed a globular apatite layer to develop on its surface. Apatite formation took place through the dissolution of the α-W phase with subsequent transformation of α-C3P into HA. After nucleation, HA growth developed a similar layer to bone-like apatite on its surface.
Although both ceramics presented in vitro bioactivity and cellular biocompatibility, our initial hypothesis about open porosity formation by complete α-W dissolution was not fulfilled. It is known that ceramics with an α-W phase of a larger grain size dissolve more quickly.
The in vitro cell culture experiments showed good morphology, adhesion, and proliferation of ahMSCs on the WC3P ceramics. It would be necessary to analyze the in vivo response of this biomaterial, because of the limitations in the in vitro tests. Nonetheless, the obtained results were sufficient to lead us to believe that it could act as an effective scaffold for bone tissue engineering.
More in vivo degradation studies must be conducted to confirm the changes in degradation properties of the W-C3P bioceramics, and as well a comprehensive mechanical study comparing our scaffold with a dense material.