Acta Biomaterialia 7 (2011) 4071–4080

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Acta Biomaterialia

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Influence of strontium on structure, sintering and biodegradation behaviourof CaO–MgO–SrO–SiO2–P2O5–CaF2 glasses

Ashutosh Goel a,⇑, Raghu Raman Rajagopal b, José M.F. Ferreira b

a Pacific Northwest National Laboratory, Richland, WA 99354, USAb Department of Ceramics and Glass Engineering, University of Aveiro, CICECO, 3810-193 Aveiro, Portugal

a r t i c l e i n f o

Article history:Received 4 May 2011Received in revised form 23 June 2011Accepted 27 June 2011Available online 2 July 2011

Keywords:StrontiumBioactive glassSurface reactivityStructureXRD

1742-7061/$ - see front matter Published by Elseviedoi:10.1016/j.actbio.2011.06.047

⇑ Corresponding author. Tel.: +1 509 371 7143; faxE-mail address: ashutosh.goel@pnnl.gov (A. Goel).

a b s t r a c t

The present study investigates the influence of SrO on structure, apatite-forming ability, physico-chem-ical degradation and sintering behaviour of melt-quenched bioactive glasses with the composition(mol.%): (36.07 � x) CaO–xSrO–19.24MgO–5.61P2O5–38.49SiO2–0.59CaF2, where x varies between 0and 10. The detailed structural analysis of the glasses is made by infrared spectroscopy and magic anglespinning–nuclear magnetic resonance spectroscopy. Silicon is predominantly present as Q2 (Si) species,while phosphorus is found as orthophosphate in all the investigated glasses. The apatite-forming abilityof glasses is investigated by immersion of glass powders in simulated body fluid for time durations vary-ing between 1 h and 7 days. While increasing the Sr2+/Ca2+ ratio in the glasses does not affect their struc-ture significantly, their apatite-forming ability is decreased considerably. Further, physico-chemicaldegradation of glasses is studied in accordance with ISO 10993-14 ‘‘Biological evaluation of medicaldevices – Part 14: Identification and quantification of degradation products from ceramics’’ in Tris–HCland citric acid buffer, and the possible implications of the ion release profiles from the glasses in differentsolutions are discussed. The addition of strontium to the glasses leads to a sevenfold decrease in chemicaldegradation of glasses in Tris–HCl. The sintering of glass powders renders glass ceramics (GCs) with vary-ing degrees of crystallinity and good flexural strength (98–131 MPa), where the mechanical propertiesdepend on the nature and amount of crystalline phases present in the GCs.

Published by Elsevier Ltd. on behalf of Acta Materialia Inc.

1. Introduction

Strontium is an abundant and widely distributed element in thegeosphere, natural water and human tissues. The amount of Sr inthe skeleton is only 0.335% of its Ca content [1]. The biological ef-fects of strontium are related to its chemical similarity to calciumand other elements in group 2A of the periodic table. Due to itssimilarity to calcium and its bone-seeking behaviour, strontiumaccumulates to a high degree in bone, can displace calcium in hardtissue metabolic processes and at high concentrations interfereswith normal bone development [2]. Because of this bone-seekingproperty, strontium drew attention as a drug for the managementof osteoporosis in the 1950s. The biological role of strontium in hu-man body has been reviewed by Nielsen [1], while its incorpora-tion and distribution in bone has been experimentally evaluatedin rats, monkeys and humans by Dahl et al. [3].

Owing to the above-mentioned beneficial aspects of strontiumin bone regeneration and considering the ion releasing ability ofglasses in aqueous medium, bioactive glasses incorporated withstrontium have gained considerable attention in the recent past

r Ltd. on behalf of Acta Materialia

: +1 509 372 5997.

for various orthopaedic applications. The increasing interest inSr-containing bioactive glasses has resulted in a multiple-fold in-crease in the number of scientific publications in this area in last3 years. To the best of our knowledge, the foremost studies onSr-containing bioactive glasses were published in 1995 by Gallianoet al. [4,5] (although they were not sure about the biocompatibilityof Sr-containing glasses). This was followed by a dormant period ofmore than a decade, when (from 2007 onwards) there was rejuve-nated interest in the role of Sr2+ in bioactive glasses, as is evidentby the boom in the number of scientific publications and patent lit-erature. Recently, Hill and Stevens [6] patented a series of stron-tium-containing bioactive glass compositions, among which aglass with composition (wt.%) 44.08SiO2–24Na2O–21.60CaO–4.43SrO–5.88P2O5 is being commercialized as Stron-Bone™ byRepRegen Ltd. Similarly, Jallot et al. [7,8] patented a series of Sr-doped SiO2–CaO–P2O5-based bioactive glasses for biomedicalapplications.

Despite the above-mentioned interesting studies, some lacunasstill exist which do not allow structure–property relationships tobe drawn with respect to the precise role of Sr2+ in glass chemistry,which in turn affects the glass dissolution, bioactivity and thermalproperties. For example, Lao et al. [9,10] studied the apatite-form-ing ability of sol–gel synthesized glasses in the SiO2–CaO–SrO and

Inc.

4072 A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080

SiO2–CaO–SrO–P2O5 systems by immersion of glass discs in stan-dard Dulbecco’s modified Eagle’s medium for a time duration vary-ing between 1 and 10 days, while Hesaraki et al. [11] followed asimilar experimental procedure in order to study the in vitro bio-activity of glasses in the SiO2–CaO–SrO–P2O5 system by immersionin simulated body fluid (SBF). The results of Lao et al. [9,10] show areduced dissolution and the growth of a newly formed phosphocal-cic surface layer. According to O’Donnell and Hill [12], this obser-vation may be explained on the basis of the design of the glasscompositions, as Lao et al. [9,10] varied the Sr2+/Ca2+ ratio on aweight basis, which would result in an increase in the silica con-tent (on molar basis), thus leading to an increase in the networkconnectivity of the glass, which would consequently slow its disso-lution and decrease its bioactivity. On the other hand, Hesarakiet al. [11] varied the Sr2+/Ca2+ ratio on a molar basis, but still theglasses exhibited a considerable delay in the formation of the apa-tite layer on the surface of Sr-doped glasses, thus raising a genuineconcern regarding the validity of the explanation provided byO’Donnell and Hill [12] as a sole reason behind the influence ofSr2+ on the dissolution kinetics of glass. Another major concernwith the above discussed studies [9–11] is that the in vitro bioac-tivity analysis was analysed in static medium, which in no mannersimulates the conditions encountered by a material in vivo.

Similarly, in order to fabricate porous scaffolds or coatings onmetallic implants from bioactive glasses, it is mandatory to under-stand the sintering conditions of the glass powders and the inter-action between the sintering and crystallization of the material.By knowing the structural transformations that occur during theheat treatment of glasses, the scaffold fabrication process can betailored, e.g. in terms of achieving the highest possible density ofthe sintered scaffolds and the required crystallinity, which itselfcontrols the material’s bioactivity [13]. However, to the best ofour knowledge, apart from a study by Lotfibakhshaiesh et al.[14], no other study depicting the influence of strontium on sinter-ing behaviour of bioactive glasses has been published to date.

In the light of above-mentioned facts, there is undoubtedlyroom to conduct a detailed analysis on the influence of strontiumon the structure, sintering and dissolution behaviour of bioactiveglasses which might help us to draw some structure–property rela-tionships in order to gain a better understanding about the mech-anism of this ion in glass chemistry.

2. Experimental

2.1. Design and synthesis of glasses

The glasses were synthesized by varying the Ca2+/Sr2+ ratio in aglass composition designed in the primary crystallization field ofdiopside (Di; CaMgSi2O6) and fluorapatite (FA; Ca5(PO4)3F)) andcorresponding to the nominal composition (wt.%): 70Di–10FA–20TCP (tricalcium phosphate; 3CaO�P2O5). The basis of designingthe parent glass composition is presented in our recent study [15].

A series of glasses with compositions (mol.%): (36.07 � x)CaO–xSrO–19.24MgO–5.61P2O5–38.49SiO2–0.59 CaF2, where x variesbetween 0 and 10, was prepared by the melt-quenching technique.The glasses were labelled in accordance with their respective SrOcontent, i.e. Sr-0, Sr-2, Sr-4, Sr-6, Sr-8 and Sr-10. High-purity pow-ders of SiO2 (purity >99.5%), CaCO3 (>99.5%), MgCO3 (BDH Chemi-cals Ltd., UK, purity >99.0%), SrCO3 (Sigma Aldrich, Germany,99.9+%), NH4H2PO4 (Sigma Aldrich, Germany, >99.0%) and CaF2

(Sigma Aldrich, Germany, 325 mesh, >99.9%) were used. Homoge-neous mixtures of batches (�100 g), obtained by ball milling, werepreheated at 1150 �C for 3 h for decarbonization (SrCO3 does notdecompose to SrO at lower temperature), then melted in Pt cruci-bles at 1550 �C for 1 h in air. The glasses were obtained in frit form

by quenching the glass melts in cold water, while monolithicglasses were obtained by casting the glass melts in a bronze mouldand further subjecting them to annealing at 500 �C for 1 h. The fritswere dried and then milled in a high-speed agate mill, resulting infine glass powders with mean particle sizes of �10–20 lm (deter-mined by the light scattering technique; Coulter LS 230, BeckmanCoulter, Fullerton, CA; Fraunhofer optical model). The amorphousnature of the glasses was confirmed by X-ray diffraction (XRD)analysis (Rigaku Geigerflex D/Max, Tokyo, Japan; C Series; Cu Karadiation; 2h angle range 10–80�; 0.02� s�1 steps).

2.2. Structural characterization

2.2.1. Infrared spectroscopyInfrared spectra of the glasses (before and after immersion in

SBF) were obtained by Fourier transform infrared spectroscopy(FTIR) using an infrared Fourier spectrometer (model Mattson Gal-axy S-7000, USA). For this purpose, glass powders were mixed withKBr in the proportion of 1/150 (by weight) and pressed into a pelletusing a hand press. A total of 64 scans for background and 64 scansper sample were made with signal gain 1. The resolution was4 cm�1.

2.2.2. Magic angle spinning (MAS)–nuclear magnetic resonance (NMR)spectroscopy

The 29Si MAS–NMR spectra were recorded on a Bruker ASX 400spectrometer operating at 79.52 MHz (9.4 T) using a 7 mm probe ata spinning rate of 5 kHz. The pulse length was 2 ls, and a 60 s delaytime was used. Kaolinite was used as the chemical shift reference.

The 31P MAS–NMR spectra of glasses were recorded on a BrukerASX 400 spectrometer operating at 161.97 MHz with 45� pulses,spinning rates of 12 kHz and a 60 s recycle delay, and the chemicalshift was quoted in ppm from phosphoric acid (85%).

2.3. Thermal analysis

The coefficient of thermal expansion (CTE) of the glasses wasobtained from dilatometry measurements, which were carriedout on prismatic samples with a cross-section of 4 mm � 5 mm(Bahr Thermo Analyse DIL 801 L, Hüllhorst, Germany; heating rate5 K min�1). The dilatometry measurements were made on a mini-mum of three samples from each composition and the standarddeviation for the reported values of CTE is in the range±0.1 � 10�6 K�1.

The crystallization behaviour of the glasses was studied usingdifferential thermal analysis (DTA; Setaram Labsys, SetaramInstrumentation, Caluire, France) of the glass powders, carriedout in air from room temperature to 1000 �C with a heating rate(b) of 10 K min�1. The glass powders (mean particle size: 10–20 lm), weighing 50 mg, were contained in an alumina crucibleand the reference material was a-alumina powder.

2.4. In vitro bioactivity analysis

The in vitro bioactivity of glasses, reflected in their capability ofinducing hydroxyapatite (HA) formation on their surfaces, wasinvestigated by immersion of the glass powders in SBF (0.10 g ofglass powder in 50 ml of SBF solution) at 37 �C. The SBF had an io-nic concentration (Na+ 142.0, K+ 5.0, Ca2+ 2.5, Mg2+ 1.5, Cl� 125.0,HPO�4 1.0, HCO2�

3 27.0, SO2�4 0.5 mmol l�1) nearly equivalent to hu-

man plasma, as discussed by Cüneyt Tas [16]. The powder–SBFmixtures were immediately sealed into sterilized plastic flasksand placed in an oven at 37 ± 0.5 �C. The sampling took place atdifferent times between 1 h and 7 days. The experiments were per-formed in duplicate in order to ensure the accuracy of the results.

A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080 4073

The apatite-forming ability on the glass powders was followed byXRD and FTIR analysis.

2.5. Physico-chemical degradation

The degradation tests were performed according to the stan-dard ISO 10993-14 ‘‘Biological evaluation of medical devices – Part14: Identification and quantification of degradation products fromceramics’’. The test comprises two parts. The first part deals withinvestigating the degradation behaviour of glasses/ceramics in acitric acid buffer solution with pH 3.0. Citric acid solution was usedbecause osteoclasts release this acid, while the extreme low pHsimulates a worst case low-end environment. The second part ofthe test simulates the more frequently encountered in vivo pH(7.4 ± 0.1) and therefore investigates the degradation of glasses/ceramics in freshly prepared Tris–HCl buffered solution. The testswere carried out without solution replacement at 37 �C and witha mixing speed of 120 rpm. The sampling was done after 120 h,when the solid and liquid phases were separated by filtering(0.22 lm, Millex GP, Millipore Corporation, USA). The solid sam-ples were then washed in deionized water and dried in an ovento a constant weight. The pH and ionic concentrations (Ca2+, Sr2+,Mg2+, P5+, Si4+; ICP-AES) of the soaking solutions were measured.The relative weight loss percentage (WL) of the glass samples after120 h of immersion in the solutions was calculated from the fol-lowing equation:

WL ¼W0 �W t

W0

� �� 100 ð1Þ

where W0 refers to the weight of the glasses before immersion andWt refers to their weight after immersion.

29Si

31P

-83

-82.75

-82.5

-82.25

-82

0 2 4 6 8 10

SrO (mol.%)

29Si

δ (

ppm

)

1.5

1.75

2

2.25

2.5

31P

δ (p

pm)

Fig. 1. Graph showing the variation in the peak positions of the 29Si and 31P MAS–NMR spectra of the investigated glasses with respect to their increasing SrOcontent.

2.6. Sintering and crystallization behaviour of glass powder compacts

The parallelepiped bars (3 � 4 � 50 mm3) were prepared fromglass powders by uniaxial pressing (80 MPa) and sintered undernon-isothermal conditions for 1 h at 850 �C at a heating rate of5 K min�1. The qualitative and quantitative analysis of crystallinephases in the GCs (crushed to particle size <25 lm) was made byXRD analysis using a conventional Bragg–Brentano diffractometer(Philips PW 3710, Eindhoven, the Netherlands) with Ni-filteredCu Ka radiation. The quantitative phase analysis of the GCs wasmade by the combined Rietveld reference intensity ratio (RIR)method. For this, 10 wt.% corundum (NIST SRM 676a) was addedto all the GC samples as an internal standard. The mixtures, groundin an agate mortar, were side loaded in am aluminium flat holderin order to minimize the problems due to non-random orienta-tions. Data were recorded in the 2h range of 5–110� (with a stepsize of 0.02� and 50 s of counting time for each step). The phasefractions were extracted by Rietveld RIR refinements, using GSASsoftware and EXPGUI as a graphical interface, rescaled on the basisof the absolute weight of the corundum originally added to themixtures as the internal standard, and thereby internally renor-malized. The background was successfully fitted with a Chebyshevfunction with a variable number of coefficients depending on itscomplexity. The peak profiles were modelled using a pseudo-Voigtfunction, with one Gaussian and one Lorentzian coefficient. Latticeconstants, phase fractions and coefficients corresponding to sam-ple displacement and asymmetry were also refined.

Microstructural observations were done on fractured GC sam-ples (etched by immersion in 2 vol.% HF solution for 2 min) byscanning electron microscopy (SEM; SU-70, Hitachi) with energy-dispersive spectroscopy (Bruker Quantax, Germany) to study thedistribution of elements in the crystals.

The mechanical properties were evaluated by measuring thethree-point bending strength of the rectified parallelepiped bars(3 � 4 � 50 mm3) of the sintered GCs (Shimadzu Autograph AG25 TA, Columbia, MD; 0.5 mm min�1 displacement).

3. Results and discussion

3.1. Structure of glasses

The bioactive properties of phosphosilicate glasses depend onthe presence of large amounts of network-modifying cations,whose charge balance is achieved by breaking Si–O–Si bridgesand creating correspondingly high concentrations of nonbridgingoxygen (NBO) atoms, and thus a significant fragmentation of thesilicate network. Indeed, experimental [17] and simulation data[18] show that the bulk structure of common melt-derived bioac-tive compositions is dominated by interconnected chains of Q2 sil-icate tetrahedra (the Qn notation denotes a tetrahedron with nbridging oxygens (BOs)) and predominantly isolated phosphategroups. Therefore, the molecular structure of the glasses undoubt-edly plays a crucial role in deciding their bioactivity and under-standing these details allows the design of new glasses withimproved chemical durability and tailored biodegradability forspecific applications.

In the present study, MAS–NMR results as presented in Fig. 1. Ingeneral, the 29Si spectra for all the investigated glasses depict thedominance of Q2 (Si) structural units in the glasses [17]. Only aslight shift in the peak positions of the spectra can be observed,and all the spectra are centred between �82.4 and �82.8 ppm,thus depicting no significant changes in the silicon coordinationin glass structure, as has also been reported by Fredholm et al.[19]. The initial addition of SrO (2 and 4 mol.%) to the parent glass(Sr-0) led to a shift in the peak position of 29Si spectra from�82.4 ppm to ��82.8 ppm (Sr-2 and Sr-4), thus signifying a slightincrease in the network connectivity of the glasses. The furtheraddition of SrO (6 mol.%) shifted the peak position of glass Sr-6to �82.4 ppm, implying decreasing connectivity in the glass struc-ture. In accordance with the 29Si NMR results, the CTE values of theinvestigated glasses were 11.4 � 10�6 K�1 (Sr-0), 11 � 10�6 K�1

(Sr-2), 10.6 � 10�6 K�1 (Sr-4), 10.7 � 10�6 K�1 (Sr-6), 11.3 �10�6 K�1 (Sr-8) and 11 � 10�6 K�1 (Sr-10). As is evident from theCTE values, the Sr-free parent glass (Sr-0) exhibits the highestCTE among all the investigated glasses, while the initial additionsof SrO (i.e. 2 and 4 mol.%) led to decrease in the CTE. However, fur-ther increasing the SrO in the glasses to 6 mol.% (Sr-6) increased

Sr-0

Sr-2

Sr-4

Sr-6

Sr-8

Sr-10

500

740

9201045

300 600 900 1200 1500

Wave number (cm-1)

Tra

nsm

ittan

ce (

a.u.

)

Glass

Sr-0

Sr-2

Sr-4

Sr-6

Sr-8

Sr-10470

565

865

1085

12401430

300 600 900 1200 1500

Wave number (cm-1)

Tra

nsm

ittan

ce (

a.u.

)1 h(b)

(a)

Fig. 2. FTIR spectra of glasses (a) before immersion and (b) after immersion in SBFsolution for 1 h.

4074 A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080

the CTE to 10.7 � 10�6 K�1. These results are in contradiction withFredholm et al. [19], who reported a monotonic increase in the CTEof glasses with increasing SrO/CaO ratio. Further, no appreciablechanges were observed in the glass transition temperature (Tg)due to variation in the SrO/CaO ratio as the Tg values obtainedfor all the glasses were 740 ± 2 �C at a heating rate of 10 K min�1.

The 31P MAS–NMR spectra of all the glasses show a predomi-nance of an orthophosphate-type environment (Fig. 1). In fact,the observed chemical shifts, 1–3 ppm, are close to that of the cal-cium orthophosphate (3.1 ppm) and that of the amorphous magne-sium orthophosphate (ca. 0.5 ppm) [17]. These results correlatewell with those reported by Lusvardi et al. [20] and Linati e al.[21] for the 45S5 glass, for which it has been deduced that the frac-tion of orthophosphate units is �82%, with the rest possibly com-prising meta- or pyrophosphates. Furthermore, the presence offluorine in the phosphosilicate glass network is known to inducesignificant changes in the local oxygen environment since the do-nor ability of fluoride is weaker in comparison to the oxide, makingthe possibility of Si–F formation unlikely. Also, according to Hayak-awa et al. [22], replacing Ca by Sr does not influence the number ofSi–F bonds, if any, present in the glass. Therefore, under such con-ditions, fluoride remains predominantly in the ionic state and pref-erentially forms ionic bonds with metal cations (Ca, Mg and Sr inpresent study), thus avoiding the depolymerization of silicate glassnetwork and forcing phosphate groups to link with silicate groups,leading to higher glass connectivity [20], which might further ham-per their bioactivity. However, it is noteworthy that, in the presentcase, the possibility of the formation of Si–O–P bonds is low, if notnegligible. In fact, according to the NMR results, phosphate groupsare not part of the actual glass network backbone.

3.2. Surface reactivity of glasses

The XRD patterns observed for all as-quenched glasses (i.e. be-fore soaking in SBF solution; not shown) exhibit a broad amor-phous halo depicting the absence of any crystallinity in glasses.The room-temperature FTIR transmittance spectra of all the inves-tigated glasses are shown in Fig. 2a. In general, FTIR spectra of allthe investigated glasses exhibit three broad transmittance bandsin the region of 300–1300 cm�1. This lack of sharp features is indic-ative of the general disorder in the silicate and phosphate networkmainly due to a wide distribution of Qn units occurring in theseglasses. The most intense bands in the 800–1300 cm�1 region cor-respond to the stretching vibrations of the SiO4 tetrahedron with adifferent number of bridging oxygen atoms [23]. Further, in all theglasses, this region (i.e. 800–1300 cm�1) is split into two transmit-tance bands, centred at �1045 and �920 cm�1. The high-frequencyband can be assigned to the Si–O asymmetric stretching mode ofBOs, whereas the �920 cm�1 one may be attributed to the Si–Oasymmetric stretching mode of the NBOs [24,25]. Furthermore,the 500 cm�1 band can be attributed to Si–O–Si bending modes[24], while the weak 740 cm�1 shoulder may be due to Si–O–Sisymmetric stretching with simultaneous Si cation motions [26].It is noteworthy that the high-frequency band at 1045 cm�1 mayalso be assigned to the asymmetric stretching of PO4 units, whichhas been reported to appear in crystalline fluorapatite at1038 cm�1 [27].

The immersion of glass powders in SBF solution for 1 h did notlead to the formation of any crystalline phases on their surfaces, asis evident from the broad amorphous halo exhibited in their XRDpatterns (Fig. 3a). However, the FTIR spectra of all the investigatedglasses showed considerable differences in comparison to the spec-tra of their respective parent glasses (Fig. 2a) after soaking in SBFsolution for 1 h (Fig. 2b). As is evident from Fig. 2b, a strong low-frequency band centred at �470 cm�1, ascribed to a deformationmode of silica layer that develops on the dissolving glass particles,

could be seen in all the glasses after immersion in SBF solution for1 h [28]. The main IR band now occurs at 1085 cm�1 and a nearbyshoulder, centred at �1240 cm�1 and attributed to Si–O–Si vibra-tion [29], can be observed in all the glasses, due to the interfacialformation of a silica gel layer, as postulated in Hench’s inorganicreaction set [30]. Further, a small sharp peak could be observedin all the investigated glasses at �565 cm�1. This is the most char-acteristic region for apatite and other phosphates as it correspondsto P–O bending vibrations in a PO3�

4 tetrahedron. A single peak inthis region suggests the presence of non-apatitic or amorphouscalcium phosphate, which is usually taken as an indication ofthe presence of precursors to HA. Apatitic PO3�

4 groups have

Sr-2

Sr-4

Sr-6

Sr-8

Sr-10

Sr-0

20 25 30 35 40

2θ (degrees)

Inte

nsity

(a.

u.)

1 h in SBF

Sr-0

Sr-2

Sr-4Sr-6Sr-8

Sr-10

20 25 30 35 40

2θ (degrees)

Inte

nsity

(a.

u.)

12 h in SBF

HA

C

Sr-0

Sr-2

Sr-4

Sr-6

Sr-8

Sr-10

HA

20 25 30 35 40

2θ (degrees)

Inte

nsity

(a.

u.)

7 days in SBF

(a)

(b)

(c)

Fig. 3. X-ray diffractograms of glass powders after immersion in SBF solution for (a)1 h, (b) 12 h and (c) 7 days.

A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080 4075

characteristic split bands at �560 and 600 cm�1, with a third signalat �575 cm�1 observed for crystallites of small size [31].Furthermore, a band at �1430 cm�1, along with another one at�865 cm�1 that is present in all glasses, might correspond to theformation of complex carbonate species connected with the pres-ence of Ca2+ ions on the surface [28] or to the incorporation of car-bonate into the phosphocalcic layer [31]. Since the XRD data didnot exhibit any calcite (CaCO3) in glasses after immersion in SBFsolution (Fig. 3a), the latter scenario seems the more feasible. Itshould be noted that the broad CO2�

3 band at �1440 cm�1 observedin most of the investigated glasses after immersion in SBF indicatesA-type substitution (i.e. carbonate replacing a hydroxyl group). TheCO2�

3 signal for B-type substitution (i.e. carbonate replacing phos-phate group) would be shifted to lower wave numbers, startingfrom �1410 cm�1 [31].

An increase in immersion time to 12 h led to the formation ofhydroxyapatite (HA; PDF card: 9–432) in the strontium-free parentglass composition, as is evident from the intense phase reflection at2h = 31.77� in the XRD data (Fig. 3b), along with the formation of

minor amounts of calcite (CaCO3), while no strontium-containingglass composition exhibited HA formation. A similar trend was ob-served for glasses after immersion in SBF for 24 and 72 h, with noother glasses besides the Sr-0 glass exhibiting apatite formation ontheir surfaces. Further, increasing the immersion time to 7 days ledto the precipitation of HA in the glasses, with its tendency decreas-ing with increasing SrO/CaO ratio, as evidenced by the XRD datapresented in Fig. 3c. These results are in agreement with those ofHesaraki et al. [11,32], who demonstrated that an increasing stron-tium content in glasses retards apatite formation on glass the sur-face in vitro. However, they contradict the results of Lao et al. [10]and O’Donnell et al. [33], who reported that strontium addition inphosphosilicate glasses enhances apatite formation on the glasssurface. As has also been mentioned in Section 1 of this paper,the major concern with these latter two studies is that thein vitro bioactivity analysis was analysed in static medium, whichin no manner simulates the conditions encountered by a materialin vivo.

Since intense ionic exchanges occur at the bioactive glass sur-face that cause major changes in the degree of supersaturationfor HA formation in biological fluids, the potential for each glassto form an apatite layer can be extrapolated from the correspond-ing evolution of the degree of supersaturation. According to Laoet al. [10], the degree of supersaturation may be defined as:SD = Q/Ksp, where Ksp is the solubility product of HA in aqueoussolution while Q is the ionic activity product for the formation ofHA. Therefore, the solution and HA mineral phase reach equilib-rium when SD = 1. For SD < 1, the dissolution of HA mineral phaseis favoured, while the solution is supersaturated with respect to HAmineral and its precipitation is favoured when SD > 1. In the pres-ent scenario, during the initial hours of immersion of the glass inSBF, the SD increases because of dealkalinization of the glass sur-face, while the formation of a silica-rich layer provides regions oflow interfacial energy, thus providing favourable sites for thenucleation of HA. However, the decreasing intensity of the XRDphase reflections for the HA phase in glass Sr-0 after 7 days ofimmersion in SBF (Fig. 3c) may be attributed to the changing de-gree of supersaturation owing to the refreshing of the SBF solutionafter every 48 h. The decreasing apatite-forming ability of stron-tium-containing glasses despite their structural similarity withthe strontium-free parent glass composition (Sr-0) may be ex-plained on the basis of the greater metal–oxygen bond strength(351 kJ mol�1 for Ca–O and 389 kJ mol�1 for Sr–O) and lower elec-tronegativity of Sr2+ (0.99) in comparison to Ca2+ (1.04), which con-sequently reduces the ability of Sr2+ to be exchanged with H+ fromthe solution [10].

It is noteworthy that although SBF tests are useful as an initialexperiment, recent evidence has indicated that the formation ofan HA layer is not a critical mechanism for bone regeneration[34] as the ionic dissolution products from bioactive glasses appearto stimulate the growth and differentiation of osteoblasts at the ge-netic level, an effect which has been found to be dose dependent[35]. Therefore, it becomes mandatory to study the ionic dissolu-tion and degradability for such glasses under normal and extremephysiological conditions so as to assess their applicability in hu-man biomedicine.

One of the relevant parameters in the study of glass dissolutionis the kind of medium used. The pH and ionic strength of the med-ium play important roles in the rate at which the glasses dissolve.The variation in pH of the two testing solutions (Tris–HCl and citricacid buffer) with respect to the SrO content in the glasses is pre-sented in Fig. 4. The introduction of Sr2+ in the parent glass (Sr-0)led to a drastic fall in the pH of Tris–HCl (Fig. 4a), while it remainedalmost unaffected by a further increase in strontium content. Thisdrop in pH may be attributed to the decrease in the chemicaldegradation of the glasses due to the addition of strontium, as is

pH

6.6

7

7.4

7.8

8.2

0 2 4 6 8 10SrO (mol.%)

pH

0

0.5

1

1.5

2

Wt.

loss

(%

)

Tris HCl

pH

6.6

7

7.4

7.8

8.2

0 2 4 6 8 10SrO (mol.%)

pH

-6

-4

-2

0

2

4

Wt.

loss

(%

)Citric acid buffer(b)

(a)

Fig. 4. Graphs depicting the change in solution pH and weight loss of glass samples(with respect to variation in SrO content in glasses) after immersion in (a) Tris–HCland (b) citric acid buffer for 120 h, respectively.

Ca

MgSr

P

Si

0

50

100

150

200

250

300

2 4 6 8

SrO (mol.%)

Con

cent

ratio

n (m

g/l)

Tris-HCl

Ca

Mg

SrP

Si

0

200

400

600

800

2 4 6 8

SrO (mol.%)

Con

cent

ratio

n (m

g/l)

Citric acid buffer

(a)

Fig. 5. ICP-AES plots of elemental concentration of Ca, Mg, Sr, P and Si in (a) Tris–HCl and (b) citric acid buffer (with respect to strontium content) after immersion ofglasses in these solutions for 120 h. (Error bars are masked by the data points in thefigure.)

4076 A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080

depicted in Fig. 4a, where a sevenfold decrease in weight loss canbe observed with the substitution of 2 mol.% CaO by SrO. Similarresults have also been presented in case of borosilicate glasses byZhang et al. [36], who found that increasing the Sr content inglasses led to a decrease in their chemical degradation. Since thepH of Tris–HCl is similar to that of the SBF solution, the results ob-tained here justify the low apatite-forming ability of strontium-containing glasses in SBF.

On the other hand, when the glasses were immersed in citricacid buffer, the pH of the solution exhibited a value of 7.27 for glassSr-0 (Fig. 4b) and remained almost constant with the addition ofstrontium to the glasses, with an average value of 7.28 ± 0.02.However, a significant variation in the chemical degradation ofthe glasses could be observed with increasing strontium/calciumratio, as is evident from the weight loss data presented in Fig. 4b.The immersion of the parent glass composition Sr-0 led to the for-mation of white-coloured crystalline products in the flasks (con-firmed by XRD), thus depicting the possibility of the occurrenceof a chemical precipitation reaction involving species dissolved incitric acid, which led to the weight gain (instead of weight loss)in the final remnant. This tendency of reaction among the ionic dis-solution products may be attributed to the presence of a fluoridecomponent in the glass composition. The mechanisms controllingaqueous corrosion of fluoride-containing glasses are complex andvary widely, depending on the testing environment. It has been re-ported that the solubility of fluoride glasses is greatly enhanced inacidic media, and a static liquid film surrounding the glass particlesis formed which becomes rapidly saturated with respect to lesssoluble species, thus causing precipitation, crystallization and theformation of an altered surface layer due to the preferentialleaching of the more soluble species [37]. However, this tendency

towards chemical reaction among ionic products decreased withincreasing strontium content in glasses, as glass compositions withSrO P4 mol.% exhibited a broad amorphous halo in XRD typical forglasses (not shown) and a monotonic weight loss, as is evidentfrom Fig. 4b.

With respect to the ion release profile of glasses in two solu-tions, previous studies have shown that Si4+ release levels in therange 0.1–100 ppm [35,38] from bioactive glass and other bioma-terials show stimulatory effects on osteoblasts and the expressionof TGF-b mRNA in human osteoblast-like cells. Additionally, Siadministration has been known to prevent trabecular bone loss,and a silicon-deficient diet results in impaired collagen synthesisand a defective skeletal structure [39]. Furthermore, the gradualrelease of soluble silica over time may not only increase cytocom-patibility but may also enhance bone bonding due to the increasedformation of Si–OH (silanol) groups, which are known to play anactive role in the precipitation of calcium phosphate [30]. In thepresent study, increasing strontium content in the glasses de-creased the release of Si4+ species in Tris–HCl as its concentrationdecreased from 85.1 ppm for glass Sr-2 to 79.5 ppm(1 ppm = 1 mg l�1) for glass Sr-8 (Fig. 5a), while in citric acid bufferits concentration remained almost constant (77.2 ± 1.2 ppm), inde-pendent of the SrO content in the glasses (Fig. 5b). This decrease inrelease of Si4+ ions in Tris–HCl with increasing Sr2+/Ca2+ ratio inglasses supports the results of lower degradation and decreasingapatite-forming ability of strontium-containing glasses. However,the concentration of Si4+ ions in Tris–HCl is still in the range atwhich beneficial properties can be attributed to various biologicalprocesses.

Further, Ca2+ ion release from experimental glasses in Tris–HClsolution decreased from 273 to 167 ppm (Fig. 5a), while it in-creased considerably in citric acid solution (Fig. 5b) with increasing

Sr-0

Sr-2

Sr-4

Sr-6

Sr-8

Sr-10

600 700 800 900 1000

Temperature (ºC)

ΔT

(μV

)

Exo

Tg

Tc

Tp10 K min-1

Tg

Tc

Tp

ΔT

700

800

900

1000

0 2 4 6 8 10

SrO (mol.%)

Tem

pera

ture

(ºC

)

130

140

150

160

ΔT (

ºC)

= (

Tc-

Tg)

(b)

(a)

Fig. 6. (a) DTA thermographs of glasses at a heating rate of 10 K min�1; (b)influence of SrO on different thermal parameters of glasses as obtained from DTAdata. (Error bars are masked by the data points in the figure.)

Table 1Thermal parameters for glasses as obtained from DTA (heating rate: 10 K min�1).

Tg (�C) Tc (�C) Tp (�C) S (�C)

Sr-0 745 862 912 117Sr-2 742 878 917 136Sr-4 742 877 920 135Sr-6 741 883 921 142Sr-8 739 884 921 145Sr-10 738 867 925 129

A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080 4077

SrO/CaO ratio. Although the significance of Ca in the process ofbone mineralization is well established, the ability of extracellularCa to regulate cell-specific responses has only recently been dem-onstrated [40]. Increased levels of extracellular Ca (from 13.1 to90 ppm) have been shown to induce osteoblast proliferation andchemotaxis through binding to a G-protein-coupled extracellularcalcium sensing receptor, and its gradual release over time may en-hance therapeutic efficacy [38,39].

With respect to Sr2+ ion release from the glass compositions, itincreased from 41 (Sr-2) to 122 ppm (Sr-8) in Tris–HCl (Fig. 5a) andfrom 77 to 312 ppm in citric acid (Fig. 5b) with increasing SrO con-tent in glasses. Sr2+concentrations from 8.7 to 87.6 ppm have beenshown to have stimulatory effect on osteoblasts, and inhibitory ef-fects on osteoclast action from 8.7 to 2102.8 ppm in vitro [41]. Theblood active Sr2+ concentration in postmenopausal osteoporoticpatients treated with strontium ranelate (Protelos�, Servier Labo-ratories, Ireland) has been measured to be 10.5 ppm [42]. It cansimilarly be perceived that any glasses releasing Sr2+ within theabove ranges will show stimulatory effects. Thus, all glasses mayhave therapeutic potential with respect to the Sr2+ release profilesmeasured herein.

Furthermore, increasing the strontium/calcium ratio in theglasses affected the ion release profile of phosphorus as its concen-tration increased from 6 (Sr-2) to 22 ppm (Sr-8) in Tris–HCl(Fig. 5a) while it decreased from 251 (Sr-2) to 187 ppm (Sr-6) incitric acid buffer (Fig. 5b). Phosphorus release from biomaterialsin controlled amounts (<30 ppm) has been shown to favour bio-mineralization and induce expression of osteogenic messengerRNA transcripts while an increase in concentration of phosphorusbeyond 30 ppm results in a decrease in cell viability [43]. Meletiet al. [43] showed that treatment of human osteoblast like cellswith 31–217 ppm phosphorus exhibited a dose- and time-depen-dent decrease in cell viability, as a dose of 217 ppm of phosphoruscaused an almost complete loss of osteoblast viability in 96 h. Inthe present study, the ionic concentration of phosphorus releasedfrom strontium-containing glasses lies within the dose limit in or-der to promote favourable biological activity, while the ionic con-centration of phosphorus in citric acid buffer might be a matter ofconcern.

Magnesium is the fourth most abundant cation in the humanbody, with an estimated 1 mol of magnesium being stored in thebody of a normal 70 kg adult, approximately half of which is storedin bone tissue [44]. Magnesium is a co-factor for many enzymes,and stabilizes the structures of DNA and RNA. The level of magne-sium in the extracellular fluid ranges between 17 and 25.5 ppm,where homeostasis is maintained by the kidneys and intestines[44,45]. In the present study, the concentration of magnesium re-leased from experimental glasses in Tris–HCl as well as in citricacid buffer decreased with increasing Sr content, as is evident fromFig. 5. However, the amount of magnesium released from theglasses is still significantly higher (80–130 ppm in Tris–HCl and318–377 ppm in citric acid buffer) than the dose required forfavourable biological activity. It is noteworthy that, although ser-um magnesium levels exceeding 25.5 ppm can lead to muscularparalysis, hypotension and respiratory distress [45], and cardiac ar-rest occurs for severely high serum levels of 145–170 ppm, theincidence of hyper-magnesium is rare due to the efficient excretionof the element in urine [44–46].

3.3. Sintering and crystallization behaviour

The DTA plots of glass powders, shown in Fig. 6a, feature anendothermic dip corresponding to Tg before the onset of crystalli-zation (Tc) and a well-defined single exothermic crystallizationcurve. The presence of a single crystallization exotherm anticipatesthat the GC is formed as a result of either single-phase crystalliza-

tion or the almost simultaneous precipitation of different crystal-line phases. Table 1 lists the values of different thermalparameters obtained for Sr-containing glasses. As has been alreadymentioned in Section 3.1, the Tg of the glasses did not show anysignificant alterations with increasing SrO/CaO ratio in thoseglasses. However, the values of Tc and the peak temperature ofcrystallization (Tp) did exhibit a tendency to increase with increas-ing strontium content in the glasses (Fig. 6b).

With respect to the sintering ability of glasses, Table 1 andFig. 6b present the values of the thermal stability parameter DT(=Tc � Tg) for all the investigated glasses. The higher values of DTcorrespond to delay in nucleation and thus provide a wider pro-cessing window for a glass composition to attain maximum densi-fication. In the present study, increasing the strontium content inthe glasses up to 8 mol.% led to an increase in the value of DT, thusimplying better sintering ability, while further increasing the SrOto 10 mol.% led to degradation of the sintering behaviour of the

Di

FA

Glass

Bendingstrength

10

20

30

40

50

0 2 4 6 8 10

SrO (mol.%)

Wt.%

80

100

120

140

160

Ben

ding

str

engt

h (M

Pa)

Fig. 8. Plot depicting the variation in crystalline/amorphous content in glassceramics along with their flexural strength with respect to SrO content in glasscompositions.

4078 A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080

glasses. However, the value of DT obtained for all the investigatedglasses was high enough to help the glass powders attain gooddensification during sintering.

In accordance with the DTA results, well sintered GCs were ob-tained after sintering of glass powders at 850 �C for 1 h, as is evi-dent from the SEM images of GCs presented in Fig. 7. Themicrostructure of the GCs did not exhibit any significant amountof porosity in the sintered glass powder compacts. The XRD resultsrevealed the presence of Di (CaMgSi2O6; ICDD card: 01-078-1390)and FA as the only crystalline phases in all the GCs. It should benoted that the GCs with SrO content varying between 0 and4 mol.% exhibited the crystallization of stoichiometric FA (Ca5(-PO4)3F; ICDD: 01-071-3848), while the GC compositions with SrOcontent >4 mol.% exhibited the crystallization of Sr-doped FA(Ca8.83Sr1.17(PO4)6F2; ICDD: 01-070-3522).

Fig. 8 presents the quantitative analysis of variation in crystal-line/amorphous content in all the investigated GCs with respectto their strontium content, while Fig. 9 shows the fit of a measuredXRD pattern of a sintered GC by using the GSAS-EXPGUI software.The fitting to the measured X-ray diagram was performed by aleast-squares calculation. The calculated diagram (Fig. 9) is basedon crystallographic structure models, which also take into accountspecific instrument and sample effects. The parameters of thismodel were refined simultaneously using least-squares methodsin order to obtain the best fit to all measured data. The differenceplot in Fig. 9 does not show any significant misfits. The differences

Fig. 7. SEM images of glass ceramics (a) Sr-2 and (b) Sr-6 after sintering at 850 �Cfor 1 h.

Fig. 9. Observed (crosses), calculated (continuous line) and difference curves fromthe Rietveld refinement of the GC Sr-6. Markers representing the phase reflectionscorrespond to corundum, fluorapatite and diopside (from top to bottom).

between the main peaks of Di and FA are caused by adjustment dif-ficulties based on the crystallinity of the phases.

As is evident from Fig. 8, Di crystallized as the primary phase inall the GCs (29–42 wt.%) followed by FA as the secondary phase,with its amount varying between 15 and 26 wt.%. The amount ofresidual glassy phase in all the investigated GCs varied between31 and 47 wt.%. The residual glassy phase in the GCs is of crucialimportance as it controls their apatite-forming ability [13,47],and the GC may turn bio-inert if there is only a small amount ofresidual glassy phase (<5 wt.%). Thus, the design of the GC compo-sitions needs to be such that equilibrium is achieved such that thefinal material still has biological and mechanical properties. In thepresent investigation, the amount of residual glassy phase is highenough to render good bioactivity to the resultant GC material.

With respect to the mechanical properties of GCs in general, allthe GCs exhibited enough mechanical strength for their applicationin the fabrication of scaffolds. The flexural strength of GCs variedbetween 98 and 131 Mpa, and was observed to depend on theamount and nature of the crystalline phases in the GCs, along withtheir amorphous/crystalline ratio, as presented in Fig. 8. The flex-ural strength increased with increasing Di/FA crystalline ratio inthe GCs. It is well known that Di-based ceramics and GCs are stron-ger in comparison to their wollastonite or FA-based counterpartsand have no general toxicity in cell cultures as they help in boneregeneration [48–50]. It is for this reason that GC Sr-4 exhibitedthe highest mechanical strength (131 MPa), owing to having thesmallest FA content, while GC Sr-10 exhibited the lowest mechan-

A. Goel et al. / Acta Biomaterialia 7 (2011) 4071–4080 4079

ical strength due to a low Di/FA crystalline ratio and the highestamorphous content.

4. Conclusions

An attempt has been made to study the influence of strontiumon the structure, surface reactivity and sintering behaviour of alka-li-free bioactive phosphosilicate glasses. The following conclusionscan be drawn from the results obtained:

(i) All the glasses exhibit a structure dominated by Q2 (Si) units,while the phosphorus was predominantly as orthophos-phate. An increase in the Sr2+/Ca2+ ratio did not affect theglass structure significantly. The CTE of glasses did notdepict a linear variation with increasing strontium content,while their Tg remained almost constant.

(ii) Strontium addition in glasses retarded their HA-formingability after immersion in SBF; however, the formation of asilica-rich gel-like layer and deposition of amorphous cal-cium phosphate on their surfaces could be detected evenafter 1 h of immersion in SBF.

(iii) The chemical degradation of glasses in Tris–HCl decreased�7-fold with substitution of 2 mol.% CaO, with SrO becom-ing gradually less accentuated with further increasing stron-tium content in the glasses. Further, increasing SrO in theglasses suppressed the weight gain resulting from the pre-cipitation reaction among leached ion products in citric acidbuffer solution and reversed this trend to weight loss. Theslow degradation of strontium-containing glasses in physio-logical fluids might be a matter of concern, considering theirlong-term applications in the human body.

(iv) The concentration of different ions released from glasses inTris–HCl and citric acid buffer solution (except phosphorus)are in the range that can provide therapeutic efficacy.

(v) The sintering of glass powders resulted in well sintered anddense GCs, with their flexural strength varying between 98and 131 MPa. The Di crystallized as the dominant phase inGCs followed by FA as the secondary phase. The residualglassy phase in GCs varied between 31 and 47 wt.%. Thelarge amount of residual glassy phase along with the goodflexural strength proves the potential of the developed GCsfor scaffold fabrication in bone tissue engineering.

Acknowledgements

The support of FCT-Portugal and CICECO is greatlyacknowledged.

Appendix A. Figures with essential colour discrimination

Certain figures in this article, particularly Figures 6, 8 and 9, aredifficult to interpret in black and white. The full colour images canbe found in the on-line version, at doi:10.1016/j.actbio.2011.06.047.

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