novel processing of bioglass ceramics from silicone resins containing micro- and nano-sized oxide...

Novel processing of bioglass ceramics from silicone resins containingmicro- and nano-sized oxide particle fillers

L. Fiocco,1 E. Bernardo,1 P. Colombo,1* I. Cacciotti,2 A. Bianco,2 D. Bellucci,3 A. Sola,3 V. Cannillo3

1Dipartimento di Ingegneria Industriale, University of Padova, Via Marzolo, 9, 35131 Padova, Italy2Dipartimento di Ingegneria Industriale, INSTM UdR Roma Tor Vergata, University of Rome "Tor Vergata", Via del Politecnico

1, 00133 Rome, Italy3Dipartimento di Ingegneria “E. Ferrari”, Universit�a degli Studi di Modena e Reggio Emilia, Via Vignolese 905, 41125 Modena,

Italy

Received 17 April 2013; revised 29 July 2013; accepted 9 August 2013

Published online 29 August 2013 in Wiley Online Library (wileyonlinelibrary.com). DOI: 10.1002/jbm.a.34918

Abstract: Highly porous scaffolds with composition similar to

those of 45S5 and 58S bioglasses were successfully produced

by an innovative processing method based on preceramic

polymers containing micro- and nano-sized fillers. Silica from

the decomposition of the silicone resins reacted with the

oxides deriving from the fillers, yielding glass ceramic compo-

nents after heating at 1000�C. Despite the limited mechanical

strength, the obtained samples possessed suitable porous

architecture and promising biocompatibility and bioactivity

characteristics, as testified by preliminary in vitro tests. VC 2013

Wiley Periodicals, Inc. J Biomed Mater Res Part A: 102A: 2502–2510, 2014.

Key Words: polymer-derived ceramics, bioglasses, glass-

ceramics, porous scaffolds, bioactivity

How to cite this article: Fiocco L, Bernardo E, Colombo P, Cacciotti I, Bianco A, Bellucci D, Sola A, Cannillo V. 2014. Novelprocessing of bioglass ceramics from silicone resins containing micro- and nano-sized oxide particle fillers. J Biomed MaterRes Part A 2014:102A:2502–2510.

INTRODUCTION

Bioglasses, like all bioactive materials, stimulate a biologicalresponse from the body, i.e. they are able to bond with liv-ing tissues, thanks to the formation of a hydroxyapatite-likelayer on their surface when they are in contact with bodyfluids. In particular, 45S5 (45 wt % SiO2, 24.5 wt % CaO,24.5 wt % Na2O, 6 wt % P2O5) and 58S (58.2 wt % SiO2,32.6 wt % CaO, 9.2 wt % P2O5) compositions are “Class A”bioactive materials, so they are not simply osteoconductiveand capable of bonding to hard tissue (bone), but they arealso osteoproductive (they specifically stimulate the growthof new bone) and capable of bonding to soft tissue. In addi-tion, these two compositions of bioglass are bioresorbable,that is they dissolve in contact with body fluids and theirdissolution products are not toxic.1,2

The mechanism at the basis of the bone bonding abilityexhibited by bioglass consists of a rapid sequence of chemi-cal reactions occurring at the surface of the implant wheninserted into living tissues, involving chemical degradationwith release of ions such as Na, Si, Ca, and causing the con-version of the surface into a carbonated-substitutedhydroxyapatite-like layer.1–4 The controlled release of ionicdissolution products from bioactive glasses also providessome stimulation of cell genes towards a path of regenera-

tion and self-repair, allowing the use of bioglass scaffolds intissue engineering applications, aimed at regenerating dis-eased or damaged tissues, rather than simply in artificialprosthesis.5,6 Glass solubility increases as network connec-tivity is reduced, whereas crystallization inhibits the ionexchange, so that it is recognized that bioactive propertiesare enhanced by an amorphous structure.3

At present, the commercial use of bioactive glasses ismainly restricted to melt-derived components, such as pow-der, granules, or small monoliths. Complex shapes, such asthree-dimensional scaffolds, may be obtained by viscous flowsintering; however, it must be noted that, as sintering isaccompanied by partial crystallization, the bioactivity is some-what degraded, as reported above. Revised formulations (e.g.13–93 bioglass, with a more complex chemical formulation)lead to limited crystallization, but again compare negatively,in terms of bioactivity, with 45S5 and 58S. In this sense,in vitro cell culture showed no marked difference in the pro-liferation and differentiated function of osteoblastic cellsbetween dense disks of 45S5 and 13–93; however, 13–93degrades and converts to an HA-like material more slowlythan 45S5 glass.4 Noncrystallized porous scaffolds may beactually obtained by a different strategy, that is by applicationof the sol-gel technique, but they have not yet been approved

*P. Colombo is also with Department of Materials Science and Engineering, The Pennsylvania State University, University Park, Pennsylvania

16801, USACorrespondence to: E. Bernardo; e-mail: [email protected]

2502 VC 2013 WILEY PERIODICALS, INC.

for clinical use.1 Furthermore, sol-gel techniques are difficultto scale-up in the industrial system for several reasons, suchas the high cost of the raw materials, the use of largeamounts of flammable solvents, the associated drying prob-lems, the complexity and the long duration of the process.7,8

The present work aims at exploring a novel route toproduce highly amorphous bioceramic foams, having thecompositions of 45S5 and 58S bioglasses, alternative toboth conventional melting techniques (and subsequent sin-tering) and sol-gel. More precisely, bioceramic foams werederived from the thermal treatment of preceramic polymers,in the form of silicone resins, containing micro- and nano-sized filler powders. The technology based on preceramicpolymer and fillers, recently presented9 as an extension ofthe research on polymer-derived ceramics,10,11 has the fun-damental advantage of combining synthesis and shaping forthe production of silicate and silicon oxynitride ceramics. Inthe case of silicates, including bioceramics such as wollas-tonite,12 the synthesis is associated with the reactionbetween silica, provided by the thermo-oxidative decompo-sition of the silicone resins, and the oxide particles, pro-vided by the fillers, whereas the shaping, especially in theform of highly porous bodies, can be easily obtained usingwell-established and conventional polymer-forming technol-ogies (e.g. direct foaming or extrusion).7,13 The adopted for-mulations, containing more than one oxide (in addition tosilica) and not corresponding to the stoichiometry of anycrystalline silicate phase, were reputed to favor the amor-phous state, as observed for common glass compositions.

Although partially successful (only foams of 58S compo-sition were not highly crystalline), the approach herereported is reputed to be significant, because of its simplic-ity, the limited processing temperature (not exceeding1000�C), the microstructural homogeneity of selected sam-ples and the promising preliminary in vitro tests.

EXPERIMENTAL PROCEDURE

Several combinations of silicone resins and fillers weretested in order to obtain the final composition of 45S5 and58S bioactive glasses. Two commercial silicones, solid(SilresVR H44, Wacker-Chemie GmbH, M€unchen, Germany)and liquid (SilresVR H62C, Wacker-Chemie, and PDMSPS340.5, United Chemical Technologies, PA), were used assilica precursors, whereas the other oxides were introducedin the form of micro- and nano-sized fillers. Micro-sized fill-ers (reagent grade chemicals, all from Sigma Aldrich Ltd,Gillingham, UK, average diameter estimated using micros-copy as not exceeding 10 lm) comprised calcium carbonate(CaCO3), sodium carbonate (Na2CO3), and sodium phosphatedibasic heptahydrate (Na2HPO4�7H2O, later referred to asNaP7H). Nano-sized fillers consisted of calcium carbonate(average diameter of d50 590 nm, specific surface of 20m2/g, PlasmaChem GmbH, Berlin, Germany) and as synthe-sized tri-calcium phosphate (Ca3(PO4)2) precursor (laterreferred to as TCP-p). The latter was synthesized accordingto a previously presented procedure, and it was used in theform of agglomerates with a maximum diameter of about10 lm, as determined by microscopy.14

The calculations of the required quantity of precursorswere made on the basis of the desired final composition,the ceramic yield of the precursors and the stoichiometry ofthe employed fillers. Silicones and fillers were mixed withor without solvents. In the former case, the silicones werefirst dissolved in isopropyl alcohol, then fillers were addedunder magnetic stirring, and the mixture was ultrasonicatedfor 10 min. A thick paste was obtained after evaporation ofthe solvent (at 60�C, overnight). In the latter case, the mix-ing procedure changed depending on the nature of the poly-mer: solid H44 was manually mixed with the fillers, bypestle and mortar, whereas liquid H62C incorporated thefillers under magnetic stirring, helped by the addition oflow-viscosity polydimethylsiloxane (PDMS). Pastes wereobtained without any drying step.

For most formulations, dicarbamoylhydrazine (DCH, AlfaAesar GmbH, Germany) powder was added to act as a foamingagent (2.5 wt % related to the overall ceramic residue), as itdecomposes at around 250�C releasing a large volume of gas.

The silicone-based mixtures, in form of both powdersand pastes, were poured in small aluminum containers andsubjected to a low-temperature treatment, at 250–300�C,aimed at both cross-linking the polymers (occurring throughthe functionalities present in the polymer structure, such as–OH or vinyl groups) and foaming, thanks to the release ofgaseous products from the decomposition of some of thefillers or of DCH.

Samples from all formulations were fired at 1000�C for1 h in air, with a heating rate of 2�C/min; in some cases thesamples were heated with a rate of 5�C/min and subjectedto an intermediate holding stage at 550�C for 4 h, aimed atoptimizing the gas evolution and the polymer-to-ceramictransformation.

The resulting ceramic components were characterizedby optical stereomicroscopy, scanning electron microscopy(JEOL, JSM-6300F), X-ray diffraction (Bruker AXS D8Advance, Germany). The Match! software package (crystalImpact GbR, Bonn, Germany) was used for phase identifica-tion, supported by data from PDF-2 database (ICDD-Interna-tional Centre for Diffraction Data, Newton Square, PA, USA).

The bulk density was obtained by considering the mass tovolume ratio for 3–6 selected ceramic blocks with dimensionsof 9 mm 3 9 mm 3 8 mm, whereas the total porosity wascomputed by considering the true density measured by gaspicnometry (AccuPyc 1330, Micromeritics, Norcross, GA). Thesame blocks were subjected to compression testing at roomtemperature, using an Instron 1121 UTM (Instron Danvers,MA, USA) operating with a cross-head speed of 1 mm/min. Atleast five samples were tested for each formulation.

The biocompatibility and bioactivity of samples C, D1,and D2 were assessed by in vitro tests through immersionin a Simulated Body Fluid (SBF) solution, according to theprocedure proposed by Kokubo et al.15 Each sample wasimmersed in 25 mL of SBF in flasks, which were thenplaced in a controlled environmental chamber at a constanttemperature of 37�C. The solution was refreshed threetimes a week (after 2, 4, 7, 9, 11 days) to reproducedynamic conditions. The pH variation induced by the

ORIGINAL ARTICLE

JOURNAL OF BIOMEDICAL MATERIALS RESEARCH A | AUG 2014 VOL 102A, ISSUE 8 2503

samples was also monitored. The scaffolds were extractedfrom the SBF solution after given times of 1, 3, 7, and 14days. Once removed from the incubation medium, all sam-ples were gently rinsed using distilled water, and finally leftto dry at room temperature for 24 h. SEM (Quanta 2000;FEI, Eindhoven, the Netherlands) analyses were performedon the samples after in vitro studies. The SEM was operatedin low-vacuum mode with a pressure of 0.53 Torr. In addi-tion, a local chemical analysis was carried out by X-rayEnergy Dispersion Spectroscopy (Inca; Oxford Instruments,Buckinghamshire, U.K.).

Raman spectroscopy was performed by means of aJobin-Yvon Raman Microscope spectrometer (Horiba Jobin-Yvon, Edison, NJ). A 632.8 nm diode laser emitting, with anoutput power of 20 mW at the sample, was employed. Pho-tons scattered by each specimen were dispersed by meansof a 1800-lines/mm grating monochromator and then col-lected on a CCD camera. The collection optic was set at1003 ULWD objective; a spectrum collection setup of 15acquisitions, each of them taking 60 s, was employed.

RESULTS AND DISCUSSION

The formulations aimed at obtaining materials with thecomposition of 45S5 and 58S bioglasses are summarized inTable I. All formulations led to highly foamed bodies, asillustrated by Figure 1. The relatively dense surface of thesamples derives from the contact of the slurries with the Alcontainer into which they were poured.

Sample A allows to observe the effect of fillers, such asNaP7H, which accomplish two distinct functions. Besidesyielding oxides (Na2O, P2O5), the filler features a remarkablewater release (corresponding to �50 wt % of the startingcompound, in accordance with its stoichiometry) associated

to dehydration occurring in the temperature range 150–350�C (as inferred from DTA analysis), leading to theobserved significant foaming. In fact, bubbles generated bywater release, upon dehydration of NaP7H, were retainedwithin the body by the setting of the H62C polymer (occur-ring rapidly above 250�C).7

However, foaming in sample A was not associated to theproduction of the desired amorphous phase; on the con-trary, as testified by Figure 2, a glass-ceramic formed, fea-turing the presence of sodium-calcium silicate phases(Na2CaSiO4, PDF#731726, referred to as N2CS; Na2Ca2Si2O7,PDF#100016, referred to as NCS) and tri-calcium phosphate[Ca3(PO4)2, PDF#290359, referred to as TCP]. Althoughinteresting (N2CS is bioactive),16 sample A was not furtherinvestigated, as the basic goal of the present research wasthe achievement of mainly amorphous scaffolds.

Accordingly, the alternative formulations for ceramicsbased on 45S5 essentially aimed at obtaining more amor-phous materials and simplifying the process. Low viscositysuspensions were achieved by mixing H62C with PDMS, with-out the addition of isopropyl alcohol. This choice, if advanta-geous in avoiding the drying step, involved some additionalproblems in the selection of fillers. As an example, NaP7Pand Na2CO3 led to large agglomerates, resulting in poorlyhomogeneous dispersions (no subsequent ceramization ofthese samples was therefore carried out). The poor homoge-neity was attributed to the interaction of the fillers with thechemical structure of PDMS. As reported in Table I, alterna-tive fillers, such as sodium sulphate and TCP-p, were foundto lead to more homogeneous dispersions and thereforethey were considered for further processing (sample B2).

Figure 1(b,c) testify the foaming of samples from thenew formulations (B1 and B2). In this case, a significant

TABLE I. Summary of Formulations for Ceramics Based on 45S5 and 58S Composition

Sample Silica precursor Fillers Foaming agent Bulk density (g/cm3) Distinctive features

A (45S5) H62C m-CaCO3 None 0.54 6 0.03 - Solvent added (isopropyl alcohol)Na2CO3 - Foaming at 300�CNaP7H - Intermediate stage at 550�C

- Highly crystallized ceramic product

B1 (45S5) H62C 20% m-CaCO3 DCH (2.5 wt %) n.a. - No solvent addedPDMS 80% Na2CO3 - Foaming at 280�C

NaP7H - Highly crystallized ceramic product- Undesirable crystal phases

B2 (45S5) H62C 20% TCP-p DCH (2.5 wt %) n.a. - No solvent addedPDMS 80% m-CaCO3 - Foaming at 280�C

Na2SO4 - Highly crystallized ceramic product- Undesirable crystal phases

C (58S) H44 TCP-p DCH (2.5 wt %) 0.60 6 0.06 - No solvent addedm-CaCO3 - Dry powder mixing

- Foaming at 270�C

D1 (58S) H62C 80% TCP-p DCH (2.5 wt %) 0.57 6 0.11 - No solvent addedPDMS 20% m-CaCO3 - Foaming at 270�C

D2 (58S) H62C 80% TCP-p DCH (2.5 wt %) 1.48 6 0.07 - No solvent addedPDMS 20% n-CaCO3 - Foaming at 270�C

m 5micro-sized; n 5 nano-sized; n.a. 5 not available data because of excessively weak samples.

2504 FIOCCO ET AL. NOVEL PROCESSING OF BIOGLASS CERAMICS FROM SILICONE RESINS

contribution to porosity was given, in addition to that of thefoaming agent (DCH present in a low amount), by the samenature of PDMS. In fact, this silicone has a low ceramicyield, i.e. it is known to transform into silica with a remark-able gas release occurring in the 300–500�C range. In fact,considering its ceramic yield of only 23%, nearly 4/5 of thestarting weight of the polymer is lost as gaseous species,whereas H62C and H44 possess ceramic yields of 58% and84%, respectively.7

Figure 1(c,d) clearly show that the porosity in sampleB2 is uniformly distributed and well interconnected, asrequired for bioceramic scaffolds. The favorable morphology,however, contrasts again with the phase assemblage. Thediffraction pattern in Figure 2 shows the formation of apotentially bioactive phase, that is calcium silicate oxide(Ca3SiO5, PDF#840594, referred to as C3S)

17 together withundesirable phases, such as unreacted sodium sulphate (ofnot proven biocompatibility; Na2SO4, PDF#781883, referredto as N2S), unreacted CaO (not suitable for medical use;

PDF#750264, referred to as CaO) and cristobalite (thatcould be responsible for the microcracks observed in thesample; SiO2, PDF#760940, referred to as Cr).18

The problems found with Na-containing fillers forced usto consider the composition 58S, featuring only CaO, SiO2,and P2O5, as reported in Table I. Also in this case, the foam-ing was substantial, as illustrated by Figure 3 and Table II.In sample C, the pore diameter was estimated to bebetween 500 mm and 2 mm, according to the analysis ofoptical microscope images [see Figure 3(a); many intercon-nections are clearly visible]. Scanning electron microscopy[Fig 3(b)] revealed also the presence of finer pores, with adiameter below 2 mm. Sample D1, on the other hand, exhib-ited pores with diameter between 100 mm and 1 mm [Fig.3(c)]. Both types of ceramic foams feature a nearly identicalamount of total porosity (�80%), with at least a significantfraction of pores with a dimension in the optimal range forscaffolds for tissue engineering (100–300 mm).19,20 The highporosity reflects in a quite poor crushing strength, well

FIGURE 1. Porous ceramics based on 45S5 composition: (a) sample A (low magnification); (b) sample B1 (low magnification); (c) sample B2 (low

magnification); and (d) sample B2 (high magnification).

ORIGINAL ARTICLE


below 1 MPa (see Table II), which is known to be a thresh-old limit for scaffolds operating “in vivo”. However, lowerstrength values could be accepted in case the scaffoldsshould be used for tissue regeneration “ex vivo” (in thiscase only the mechanical properties of the final tissue-engineered implant are critical).1,4,21–23

Differently from the other cases, sample D2 [Fig. 3(d)]was produced starting from nano-sized CaCO3. This choicederived from the observation, based on very recent experi-ences, that nano-sized fillers usually enable a superiorchemical homogeneity in the component after ceramic con-version.13 It may be noted that the porosity was not uni-form and practically comprised only large voids, probablybecause of air bubbles trapped upon mixing (the totalporosity did not exceed 35%): the entrapment was likelyfavored by the fact that nano-sized particles considerablyincreased the viscosity of the suspension, limiting the possi-bility of achieving a significant foaming and leading to apoor homogenization of the slurry.

The last samples were found to be less crystalline thanthe previous ones. In fact, the diffraction patterns, as shownin Figure 4, contain an amorphous halo, located in theregion typical of silicate glasses (i.e. �21�), especially forsamples D1 and D2, below the main crystal peaks. The crys-tal phases consisted of calcium silicates (para-wollastoniteCaSiO3, PDF#760925, referred to as p-w; Ca2SiO4,PDF#860401, referred to as C2S,) and TCP (PDF#290359).

Interestingly, there was no difference in the crystallinephases present: this means that the choice of the silica pre-cursors (H44 or H62C-PDMS) and the granulometry of theCaCO3 (nano- or micro-sized) had no influence on the typeof crystalline phases formed. On the other hand, differenceswere observed in terms of the overall crystallinity and dis-tribution of phases, as reported in Table II. According to thesemi-quantitative analysis provided by the Match! programpackage, para-wollastonite was dominant in sample C (25wt % of the crystal phase), whereas TCP was the mainphase in samples D1 and D2 (respectively, 16 wt % and 20wt % of crystal phase). C2S was present in an amount of�5 wt % in sample C, and �2 wt % in samples D1 and D2.

The bioactive behavior of samples C, D1, and D2 was inves-tigated by means of immersion tests in SBF, which claims tomimic the acellular human blood plasma. In particular, as thebone bonding ability of a biomaterial, also termed osseointe-gration,24 is associated with the formation of an apatite layeron its surface, the ability to bond to bone can be preliminarilyassessed by monitoring the formation of such apatite layer invitro in SBF. In the present research activity, it was observedthat, for all samples, the deposition of an apatite layerpromptly began after 1 day in SBF, as shown in Figure 5(a) forscaffold C, where white globular precipitates with the typicalapatite morphology can be observed on the surface. This factis further confirmed by the results of the EDS analysis per-formed on the globular precipitates, which were substantiallylarger after 3 days [see in Fig. 5(c) the data for sample D2].The EDS spectrum revealed the presence of Na, O, Ca, Si, P, Cl;in particular, apart from local fluctuations, the Ca/P ratio wasabout 1.67, which is similar to that of stoichiometric apatite.25

Analogous results were obtained for samples C and D1 (datanot reported for the sake of brevity). The presence of Cl can beascribed to chloride compounds precipitated from SBF, aswidely reported in the literature,26 wheras Si is because of asilica gel formed underneath the apatite precipitates.27 Thedissolution of the sample in SBF resulted in the formation ofhighly corroded areas and surface roughness [Fig. 5(b)], whichare expected to favor cell adhesion, proliferation, differentia-tion, and detachment strength,28 thus promoting an adequateosseointegration of the material.

After 7 days in SBF, the surface of the samples was com-pletely covered by apatite [Fig. 5(d)] and, at the same time,it appeared to be further altered by the SBF action. In par-ticular, the internal pore structure of the samples was alsocoated with apatite precipitates. The EDS analysis carriedout on the precipitates covering the surface of sample D2after 7 days in SBF [Fig. 5(e,f)] suggests the presence of aparticularly thick apatite layer covering the silica gel, as thespectrum shows a lower Si amount with respect to the anal-ysis performed after 3 days in SBF [Fig. 5(c)]. A prolongedsoaking in SBF (14 days, data not reported for the sake ofbrevity) led to further apatite deposition. The cauliflower-like precipitates progressively grew and completely coveredthe surface of the scaffold struts [Fig. 5(e)].

Thanks to the particularly high intensity of the Ramanpeaks associated to P–O vibration modes, the developmentof an apatite film on the surface of samples could be

FIGURE 2. Diffraction patterns of selected ceramics based on 45S5

composition.


confirmed by Raman spectroscopy. The Raman spectraacquired on the spherical precipitates covering the samplesafter 7 days in SBF are shown in Figure 6. Vibration bandscharacteristic of apatite were observed at about 430 cm21

(m2PO432), 590 cm21 (m4PO4

32), 960 cm21 (m1PO432, the

most intense band), and 1005 cm21 (m3PO432).29,30 Raman

spectroscopy is particularly important for the analysis of

the apatite deposited in vitro because it emphasizes the C–Ovibrations, as the in vitro grown apatite is usually carbo-nated.31 In this respect, it is possible to observe a peaklocated at about 1070 cm21 in the spectra reported in Fig-ure 6, which can be ascribed to the stretching of carbonategroups, thus confirming that the apatite film on samples C,D1, and D2 is carbonated.

The obtained samples were particularly interesting alsoin terms of the pH variation induced in the SBF, as shown inFigure 7. This issue is extremely important for the biocom-patibility of glasses, glass-ceramics, and bioceramic coatingsin general. In fact, when these systems are soaked in physio-logical fluids, the rate and amount of ion release can engen-der an excessive rise in the pH level or abrupt changes inpH, which are incompatible with life. Moreover, it has beenreported in literature that the pH value influences the proteinadsorption on the surfaces of biomaterials,32 which is a fun-damental step among the biological reactions taking place atthe interface between biological environment and medicalimplant. For example, Chen et al. reported high pH values(between 8.5 and 9) for 45S5 BioglassVR scaffolds soaked inSBF and observed that the attachment and stability of colla-gen on these systems was reduced.33 A dramatic increase inpH (between 9 and 9.5) during the first days of exposure to

FIGURE 3. Microstructural details of ceramic samples based on composition 58S: (a,b) sample C; (c) sample D1; and (d) sample D2.

TABLE II. Properties and Distribution of Crystal Phases in

Ceramic Materials of 58S Composition

Sample C D1 D2

PropertiesBulk density (g/cm3) 0.60 6 0.06 0.57 6 0.11 1.48 6 0.07True density (g/cm3) 2.88 6 0.01 2.62 6 0.01 2.19 6 0.02Porosity (vol %) 79 78 33Crushing strength

(MPa)0.57 6 0.35 0.28 6 0.20 4.92 6 0.68

Crystallinity (wt %) 50 30 43Distribution of crystal phasespara-wollastonite

(wt %)25 12 19

C2S (wt %) 5 2 2TCP (wt %) 20 16 20

ORIGINAL ARTICLE


SBF for 45S5 BioglassVR samples was also reported by Bel-lucci et al., who observed that the pH value of the SBF stabi-lized near to physiological values only after three weeks ofimmersion.34 A rapid pH increase was also observed for thetreatment of 45S5 BioglassVR in DMEM tissue culture medium(pH�9 after 24 h) and in phosphate-buttered saline solution(pH�11).35,36 In this regard, stable pH values close to 7.8can be considered optimal for osteoblast adhesion and prolif-eration37 and many investigations have been reported in lit-erature in order to develop new compositions tailored toavoid too high pH levels in SBF.38,39 Figure 7 shows the pHtrend as a function of the immersion time in SBF for samplesC, D1, and D2. All samples induced lower pH values com-pared to those commonly reported for 45S5 BioglassVR -derived glass ceramics. This fact is attributable to a slowerion leaching in our samples. In particular, D1 and D2 samplesappear particularly promising, as the pH stabilizes near tophysiological values already after a few hours.

The SBF proposed by Kokubo et al. is an acellular solutionand, as such, it is not able to simulate the real complexity of abiological environment, which also includes proteins, cells, andso forth.15,40. Therefore, SBF tests mainly offer an insight intothe inorganic reactions that are expected to occur after thematerial is implanted into the human body, whereas the assess-ment of the biological response of cells needs further experi-mental trials, such as cytotoxicity tests.41 Nevertheless theimmersion of a new material in SBF is extremely useful toFIGURE 4. Diffraction patterns of selected ceramics based on 58S

composition.

FIGURE 5. Formation of an apatite layer on the surface of samples C, D1, and D2 after immersion in SBF for different lengths of time: (a) sample

C—1 day; (b) sample D2—3 days; (c) sample D2—results of the EDS analysis performed on the spot indicated in (b); (d) D2 samples—7 days;

(e) sample D2—7 days; (f) sample D2—results of the EDS analysis performed on the spot indicated in (e).


estimate its attitude to stimulate the formation of a surfacelayer of hydroxyapatite, which is similar to the mineral compo-nent of bone.40 Even if it is known from the literature that theapatite-forming ability in vitro does not always imply the bone-bonding ability in vivo,42 it is interesting to observe that all thescaffolds tested in SBF were able to support the developmentof a surface layer of hydroxyapatite. Moreover it is worth not-ing that samples C, D1, and D2 were produced without theaddition of solvents and were processed to obtain a silicateamorphous matrix with some crystalline phases (calcium sili-cates, tri-calcium phosphate, Fig. 4) that are by no means dan-gerous to the human body.43–45 Moreover no anomalouscompounds (possibly deriving from the materials themselvesor from unwanted reactions between the scaffolds and the liq-uid medium) were detected by SEM, X-EDS analysis, andmicro-Raman spectroscopy. In addition, the pH of the SBFnever exceeded the value of 8.5 (Fig. 7), which represents aninteresting result, since, as already mentioned, the reactionmechanisms of bioactive glasses (also sintered ones) usuallyinduce very basic pH conditions, potentially harmful forcells.33–39 For these reasons, the new scaffolds are suitable can-didates for tissue engineering applications and further investi-gations on the cytotoxicity and other biological performancewill be the target of future work.

CONCLUSIONS

We may conclude that:

� Mixtures of silicones and fillers enabled to produce highlyporous glass ceramic products using a simple processingprocedure, such as foaming at 200–300�C followed byheating at 1000�C.

� All samples of nominal composition corresponding to thatof 45S5 bioglass were highly crystalline and led to diffi-culties concerning the selection of a suitable sodiumoxide precursor. For a selected combination of siliconesand fillers (formulation A) the homogeneity of foamingwas accompanied by the formation of potentially bioac-tive phases, to be further investigated in the future.

� Silicone/fillers mixtures of nominal composition correspond-ing to that of 58S bioactive glass yielded partially amorphouscomponents. It was observed that porosity and mechanicalproperties could be adjusted by modifying the starting pre-cursors (micro- or nano-sized powders and silicones).

� Samples with 58S composition exhibited a quite low com-pressive strength, but they effectively promoted thedevelopment of a surface layer of hydroxyapatite whenimmersed in a SBF. The reactions occurring between thescaffolds and the soaking medium did not cause largechanges of the pH, as frequently observed for conven-tional bioglass-derived materials and scaffolds. This isadvantageous, as basic values of the pH or abruptchanges are known to be dangerous for cells.

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FIGURE 6. Raman spectra acquired on samples C, D1, and D2 immersed

in SBF for 7 days. The main peaks of apatite are indicated by arrows.

FIGURE 7. pH variation induced by samples in SBF, refreshing the

solution every 48 h.

ORIGINAL ARTICLE


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