Influence of SrO substitution for CaO on the propertiesof bioactive glass S53P4

Jonathan Massera • Leena Hupa

Received: 22 August 2013 / Accepted: 6 December 2013 / Published online: 12 December 2013

� Springer Science+Business Media New York 2013

Abstract Commercial melt-quenched bioactive glasses

consist of the oxides of silicon, phosphorus, calcium and

sodium. Doping of the glasses with oxides of some other

elements is known to affect their capability to support

hydroxyapatite formation and thus bone tissue healing but

also to modify their high temperature processing parame-

ters. In the present study, the influence of gradual substi-

tution of SrO for CaO on the properties of the bioactive

glass S53P4 was studied. Thermal analysis and hot stage

microscopy were utilized to measure the thermal properties

of the glasses. The in vitro bioactivity and solubility was

measured by immersing the glasses in simulated body fluid

for 6 h to 1 week. The formation of silica rich and

hydroxyapatite layers was assessed from FTIR spectra

analysis and SEM images of the glass surface. Increasing

substitution of SrO for CaO decreased all characteristic

temperatures and led to a slightly stronger glass network.

The initial glass dissolution rate increased with SrO con-

tent. Hydroxyapatite layer was formed on all glasses but on

the SrO containing glasses the layer was thinner and con-

tained also strontium. The results suggest that substituting

SrO for CaO in S53P4 glass retards the bioactivity. How-

ever, substitution greater than 10 mol% allow for precipi-

tation of a strontium substituted hydroxyapatite layer.

1 Introduction

The potential of bioactive glasses S53P4 (BonALive�) and

45S5 (Bioglass�) as bone regeneration grafts has been

verified in several clinical studies. Today, these FDA

approved glasses are commercially available as particulates

for certain clinical applications. The oxide formulation of

S53P4 and 45S5 is within the four component system SiO2,

CaO, Na2O and P2O5. The bioactivity, expressed as the

capability of these glasses to support bone regeneration is

closely related to their low silica content giving the glasses

a special reaction pattern in biological solutions. The bio-

activity has been related with the capacity of the glasses to

release ions of calcium, sodium and soluble silica followed

by repolymerization of a silica rich inner and hydroxyap-

atite (HA) outer surface layers. The HA layer then pro-

motes new bone formation and bonds the glass with

biological bone [1]. Today, the calcium and silicate ions

released from the glass are known to promote the bone

formation on the cellular level [2]. In addition, doping the

glasses with other ions such as strontium, magnesium,

copper and zinc are known to affect healing and growth of

vascularized bone tissue [3].

In the recent years, much attention has been focused on

the processing of bioactive glass fibers and 3D scaffolds for

various biosensing, wound healing and tissue engineering

applications [4, 5]. The strong tendency of bioactive glas-

ses to crystallize at typical glass forming operations chal-

lenges the manufacture of 3D scaffolds and continuous

fibers. Recently, the crystallization mechanisms of bioac-

tive glasses 45S5 and S53P4 were discussed in detail [6].

45S5 showed a crystallization mechanism that could not be

described by the Johnson–Mehl–Avrami model and, hence

was more complicated than a simple nucleation and

growth, while preferential surface crystallization was

measured for S53P4. In addition, S53P4 had narrower

crystallization domain than 45S5. In heating the crystalli-

zation commenced by nucleation and growth of Na2Ca2-

Si3O9 followed by the formation of Ca3(PO4)2 and/or

J. Massera (&) � L. Hupa

Process Chemistry Centre, Abo Akademi University,

Biskopsgatan 8, 20500 Turku, Finland

e-mail: jmassera@abo.fi

123

J Mater Sci: Mater Med (2014) 25:657–668

DOI 10.1007/s10856-013-5120-1

Na2Ca4(PO4)2SiO4 at temperatures higher than 850 �C

[6, 7].

Glass 45S5 containing 45 wt% SiO2 shows more rapid

formation of typical reaction layers in vivo and in vitro

than S53P4 containing 53 wt% SiO2 [8]. The higher silica

content in S53P4 gives higher network connectivity and

thus suggests for easier hot-working into various shapes.

The more controllable crystallization of the glass S53P4

encouraged us to study how the in vitro reactivity and

crystallization tendency are affected by replacing CaO in

S53P4 by other alkaline earth oxides. The results of

substituting MgO for CaO up to 4 mol% were published

recently [9]. While the addition of MgO efficiently exten-

ded the high temperature working range domain, the

reactivity in simulated body fluid (SBF) slightly reduced

[9]. When compared to S53P4, the progressive substitution

of MgO for CaO slowed down the hydroxyapatite layer

formation as demonstrated by thinner reaction layer at the

glass surface and lower pH increase of the solution with

increasing immersion time [9]. Similarly to soda-lime

glasses the effect of MgO on the high temperature prop-

erties and the in vitro bioactivity could be partly related to

the increased network connectivity of the glass structure

induced by MgO [9, 10].

Strontium ions are known to have therapeutic effect on

osseous healing. Today it is used as strontium ranelate for

treating osteoporosis [11]. SrO has beneficial effects in vitro

and in vivo for the replication of pre-osteoblastic cells while

decreasing the activity of osteoclasts [12]. SrO is a glass

network modifier similar to CaO in soda-lime glasses.

Substituting strontium oxide for calcium oxide in bioactive

glass 45S5 reduced the glass transition temperature and

increased the onset of crystallization [13]. Goel et al. reported

that up to 10 mol% of SrO substituted for CaO decreased

significantly the apatite forming ability. In contrast, Fredholm

et al. demonstrated that full substitution of SrO for CaO allows

for enhanced apatite formation [14, 15]. In both studies the

glass compositions were close to that of 45S5.

This study was undertaken to evaluate the effect of

replacing CaO with SrO on the in vitro, thermal, and

structural properties of S53P4. Characteristic temperatures

were measured using differential thermal analysis and hot

stage microscopy. The crystallization behavior of the

glasses was assessed from the calculated values of the

activation energy for crystallization. The glass structure

was defined in terms of the attenuated total reflectance

measurements. In vitro reactivity of the glasses was per-

formed in SBF. The ions released in solution were analyzed

and change in the solution pH was recorded as functions of

immersion time. The layers formed at the surface of the

glass particulates were determined via SEM/EDX analysis

and changes in the glass surface structure were recorded

using ATR-FTIR.

2 Experimental procedure

2.1 Glass processing

The effect of the substitution of SrO for CaO on the high

temperature and in vitro properties of the bioactive glass

S53P4 was studied for glasses melted from batches mixed

of sand (99.4 % pure SiO2), and analytical grades of

Na2CO3, SrCO3, CaCO3, and CaHPO4�2H2O. The nominal

oxide compositions of the experimental glasses are in

mol % 53.85SiO2–22.66Na2O–1.72P2O5–(21.77–x)CaO–

xSrO with x varying from 0 to 21.77 (Table 1). The glasses

were coded according to the SrO content: S53P4, Sr5,

Sr10, Sr15, and Sr21.77. The glasses were melted in air in a

platinum crucible between 1,360 �C (S53P4) and 1,450 �C

(Sr21.77) for 3 h and then re-melted to ensure good

homogeneity. After casting in a graphite mold, the glasses

were annealed at 520 �C for 1 h and then allowed to cool

slowly to room temperature in the annealing furnace. The

annealing temperature and time were chosen to allow

removal of internal stress due to quenching while mini-

mizing the risk for surface nucleation. After melting the

composition of the glasses measured via EDX agreed with

the nominal composition within the accuracy of the

measurements.

2.2 Thermal, physical and structural properties

The glass transition temperature Tg and the crystallization

temperature Tp of the glasses were determined using differ-

ential thermal analysis (TGA/SDTA, Mettler Toledo TGA/

SDTA851e) at various heating rates (10, 15, 20, and

30 �C min-1). The measurements were performed on 50 mg

samples using the particle size 300–500 lm in platinum pans

in an N2 atmosphere. The glass transition temperature Tg was

taken at the inflection point of the first endotherm, obtained

by taking the first derivative of the DTA curve and Tp was

taken at the maximum of the exothermic peak. All DTA

Table 1 Oxide compositions of the experimental glasses in mol%

(wt%)

Oxide S53P4 Sr5 Sr10 Sr15 Sr21.77

Na2O 22.65

(23.00)

22.65

(22.14)

22.65

(21.34)

22.65

(20.59)

22.65

(19.67)

CaO 21.77

(20.00)

16.77

(14.83)

11.77

(10.03)

6.77

(5.57)

0.00

(0.00)

SrO 0.00

(0.00)

5.00

(8.17)

10.00

(15.75)

15.00

(22.80)

21.77

(31.60)

P2O5 1.72

(4.00)

1.72

(3.85)

1.72

(3.71)

1.72

(3.58)

1.72

(3.42)

SiO2 53.85

(53.00)

53.85

(51.01)

53.85

(49.17)

53.85

(47.46)

53.85

(45.32)

658 J Mater Sci: Mater Med (2014) 25:657–668

123

scans were repeated 3–4 times and the values obtained for the

characteristic temperatures are given as average values. The

thermal properties were studied also using a hot stage

microscope (HSM/ODHT, Misura Expert Systems). Sam-

ples were prepared by pressing glass powder (\45 lm) into

cylinders (height = 3 mm, diameter = 2 mm) with a

spring-loaded hand-press to ensure identical pressure for

each sample. The samples were heated to 480 �C at

40 �C min-1 and then to 1,200 �C at 5 �C min-1. They were

imaged at room temperature and at every 5 �C from 500 �C

to record changes in the sample height and shape at

increasing temperature. The sintering curve given by the

sample height as a function of temperature was used to define

the sintering temperature, Tsi, the starting and ending of the

plateau, Txi and Txf, corresponding to the glass crystallization

domain and the temperature at which the crystallized glass

goes through fusion process, Tf [16].

The activation energies associated with the glass tran-

sition temperature Ea and the crystallization temperature Ec

were determined using the Kissinger equation (Eq. 1) by

measuring Tg and Tp at different heating rates [17]:

lnq

T2g

!¼ � Ea

RTg

þ Constant; ð1Þ

where q is the heating rate (10, 15, 20, and 30 �C min-1);

Tg is the glass transition temperature; and R is the gas

constant. The value of Ec was calculated by replacing Tg

with Tp in the equation.

Thermal treatments were performed on 1 cm3 glass

cubes placed on a platinum foil in a preheated furnace. The

heat treatments were performed at 750 and 850 �C. These

temperatures were chosen within the main crystallization

domain based on DTA measurements. After 1 h at the

preset temperature, the samples were cooled rapidly in

static air. The samples were cut and polished for XRD

analysis. The crystalline phases were identified using an

X-ray diffraction analyzer (Philips X’pert) with Cu Karadiation (k = 1.5418 A). The scans were performed from

2h = 0�–60� with a step size of 0.02�.

The IR absorption spectra of the glasses were recorded

using Fourier transform infra-red spectroscopy (FTIR, Per-

kin Elmer) in attenuated total reflectance (ATR) mode on

powdered glasses. All spectra were recorded within the range

650–1,600 cm-1, corrected for Fresnel losses and normal-

ized to the absorption band showing the maximum intensity.

SEM/EDX (Leo 1530 Gemini from Zeiss and EDXA

UltraDry from Thermo Scientific) was used to image and

analyze the composition of the samples. The accuracy of

the elemental analysis is *0.1 wt%.

The density of bulk glass materials was measured by

Archimedes’ principle using deionized water. The accuracy

was better than 0.02 g cm-3.

2.3 In vitro bioactivity

Glass particulates with grain size 300–500 lm were

immersed in 50 ml of simulated body fluid (SBF) for 6, 24,

48, 72, and 168 h at 37 �C in an incubating shaker (Stuart

orbital incubator SI500). In the shaker, the orbital speed of

120 RPM was chosen to give laminar flow mixing of the

solution without inducing particle movement. The SBF was

prepared using the protocol developed by Kokubo et al.

[18]. Two series of experiments were conducted. In the

first series the mass of sample to volume of SBF ratio

was kept constant (75 mg/50 ml). In the other series the

mass of sample immersed in the solution was adjusted

to give a constant surface area to volume ratio (Sa/V &8.5 9 10-2 cm2 ml-1). As the glass was crushed by hand

and then sieved, the Sa was calculated assuming that the

average particles size was 400 lm in diameter and using

the glasses density. The masses varied from 75 mg for

glass S53P4 to 85.4 mg for the CaO-free glass. For each

time point four parallel samples of each glass were studied.

The change in the solution pH was recorded for each

immersion time and compared to a blank sample contain-

ing only SBF. After testing, the powder was washed with

acetone and dried. Part of the powder was embedded in

resin and polished to reveal the particulates’ cross section.

The composition and structure of the glass powder were

analyzed with SEM/EDX and FTIR-ATR.

Inductively coupled plasma-Optical emission spec-

trometer (ICP-OES; Optima 5300DV, Perkin Elmer) was

employed to quantify the amount of P, Sr, Ca and Si ions in

SBF after immersion of the glasses.

3 Results

3.1 Thermal properties

Figure 1 shows the DTA thermograms of the glasses recor-

ded at the heating rate 30 �C min-1. The glass transition

temperature decreased with increasing SrO content. The

main exothermic peak, corresponding to the glass crystalli-

zation first shifted to lower temperatures for Sr5 and then

shifted to higher temperatures for glasses with higher SrO

content. A shoulder appears on the DTA thermogram of Sr10

at around 750 �C. This shoulder shifts to lower temperatures

with increasing SrO content in the glass. The endothermic

event occurring at higher temperatures has been ascribed to

the melting of the crystals. The onset of melting was not

clearly distinguishable for all glasses. The thermograms

suggest that the liquidus temperature, Tl, taken as the endset

of the endothermic event shifted to lower temperatures when

SrO was substituted for CaO.

J Mater Sci: Mater Med (2014) 25:657–668 659

123

Figure 2 presents the HSM curves of the glasses. The

sintering (Tsi), onset of crystallization (Txi), end of crys-

tallization (Txf) and glass fusion (Tf) temperatures have

been assigned in the figure. The sintering temperature is

defined in the DIN51730 standard as shrinkage of the ini-

tial height to 95 %. All characteristic temperatures

decreased with increasing SrO content. Tsi shifted from

(625 ± 5) to (565 ± 5) �C. Txi and Txf, the temperatures at

which the particles start and cease crystallizing, decreased

from (663 ± 5) to (629 ± 5) �C and from (1054 ± 5) to

(926 ± 5) �C, respectively. Finally, Tf decreased from

(1,110 ± 5) to (975 ± 5) �C with progressive substitution

of SrO for CaO.

Figure 3a presents DT = Tg - Txi as a function of the

SrO content in the glass. DT, the temperature difference

between the onset of crystallization and glass transition, is

a gauge to the resistance of the glass to crystallize upon

heating; the smaller the difference the easier the crystalli-

zation. Txi was taken from the HSM curves while Tg was

based on the DTA trace obtained at a heating rate of

10 �C min-1. For all glasses the DT value was greater than

100 �C. Glass Sr5 showed the lowest value, around

105 �C, while for the other glasses DT was around 116 �C.

Figure 3b presents the activation energies associated

with the glass transition temperature (Ea) and with the main

crystallization peak temperature (Ec). Ea and Ec were cal-

culated using Eq. (1) and the measured values of Tg and Tp

for the different heating rates in DTA. The activation

energy for the glass transition temperature first drastically

decreased from (790 ± 30) for S53P4 to (442 ± 30)

kJ mol-1 for Sr5. With further increase in the SrO content

Ea first slightly decreased and then increased for the

composition in which SrO was totally substituted for CaO.

The activation energy for crystallization decreased with

increasing SrO. A local minimum was measured for the

glass Sr5. The activation energy for the crystallization

corresponding to the shoulder, in the DTA thermograms of

Sr10, Sr15 and Sr21.77, could not be calculated, as the

shoulder was not always well resolved for the lowest

heating rates.

Monoliths (1 cm3) of each glass were heat treated in air

at 750 and 850 �C. According to the DTA and HSM

curves, these temperatures are within the crystallization

Fig. 1 DTA thermograms of the investigated glasses

Fig. 2 Sintering curves for the investigated glasses expressed as the

change in sample height as a function of temperature obtained from

HSM

Fig. 3 DT (a) and the activation energies Ea and Ec (b) versus SrO

content in the glasses

660 J Mater Sci: Mater Med (2014) 25:657–668

123

domain of the glasses. The XRD pattern prior to heat

treatment (not shown) of each glass only exhibited a broad

band characteristic of the amorphous state. Figure 4 pre-

sents the XRD patterns after the heat treatment at 750 �C

(4a) and 850 �C (4b). Sharp peaks were obtained for all

glasses at 750 �C. At increasing SrO content the peaks

shifted toward lower 2H angle and their number and

intensity decreased. After the heat treatment at 850 �C the

XRD patterns of the glasses (Fig. 4b) exhibited an increase

in the intensity of the sharp peaks and new peaks appeared.

For S53P4 the crystals at 750 �C were characteristic of

Na2Ca2Si3O9 (ICCD card #01-075-1687) and at the higher

temperature Ca3(PO4)2 crystals (ICCD card #00-06-0200)

formed [6]. Most likely some Na2Ca4(PO4)2SiO4 crystals

were also present [7]. For glasses containing up to

10 mol% of SrO the peaks are similar but they shift to

lower angles with increasing SrO content. For the glasses

Sr15 and Sr21.77 some new peaks that could not be related

to either Ca3(PO4)2 or Na2Ca4(PO4)2SiO4 were recorded.

3.2 Structural properties

The influence of the gradual substitution of SrO for CaO on

the structural properties of S53P4 was investigated using

infrared spectroscopy. The FTIR spectra of the glasses are

shown in Fig. 5. All spectra were normalized to the band

with the maximum intensity at 930 cm-1. The spectra

show absorption bands also at 748, 1,023 cm-1 and in the

1,400–1,515 cm-1 region. The band at 930 cm-1 can be

attributed to Si–O- in [SiO4] and the band peaking at

1,023 cm-1 to Si–O–Si asymmetric stretching in [SiO4]

units [19–21]. The band at 748 cm-1 is due to Si–O

bending [22]. The band located within 1,400–1,515 cm-1

is related to carbonate in the glass structure [23].

No significant change in the absorption bands could be

recorded with progressive replacement of SrO for CaO.

3.3 In vitro properties

The glasses were immersed in simulated body fluid (SBF)

for 6, 24, 48, 72, and 168 h. Figure 6 presents the average

pH values of the solutions as functions of immersion time.

Figure 6a shows the pH for the samples with a constant

mass of glass to volume of SBF ratio, whereas Fig. 6b

presents the pH for samples having constant surface area to

volume ratio (Sa/V).

When using constant mass to volume ratio the pH of all

samples showed similar pH pattern up to 24 h of immer-

sion. For longer immersion times, all samples but Sr21.77

gave lower pH than S53P4. The pH differences between

the samples are small compared to the accuracy of the

measurements.

Density was measured for all samples. Density increased

from (2.66 ± 0.02) for S53P4 to (3.03 ± 0.02) g cm-3 for

Sr21.77 (Table 2). The density values were used to

Fig. 4 XRD patterns of the glasses after 1 h at a 750 �C and b 850 �C. (#: aluminum sample holder, *: Na2Ca2Si3O9, ?: Ca3(PO4)2, 0: SrSiO3,

!: Na2SrSi2O6, ?: Si2SiO4 and or Sr3(PO4)2)

Fig. 5 FTIR/ATR spectra of the investigated glasses

J Mater Sci: Mater Med (2014) 25:657–668 661

123

calculate the masses of the samples needed for maintaining

constant Sa/V ratio (Table 2). For comparison the theo-

retical mass to be weighed in order to keep a constant

molar amount of Na, Ca ? Sr, P and Si amount in contact

with SBF is also provided in the table.

Figure 6b indicates that the pH during the first 24 h

depends strongly on the SrO content in the glass. pH increase

is faster with increasing SrO content. The pH values for the

two test series after 1 week of immersion are given in

Table 2. (The final pH also increases with SrO content). In

the samples with constant Sa/V ratio, the molar amount of

total alkalis is constant but the pH value gradually increases

with the SrO content. For the series with constant mass to

volume ratio, the final pH value is of the same level for all the

glasses, although the molar amount of total alkalis decreases

in these samples with increasing SrO content.

Figure 7 presents the evolution of the concentrations of

the Sr, Ca, P and Si ions in SBF as a function of immersion

time for constant Sa/V. These ions are first released from the

glasses into the solution and later consumed when the typical

surface layers consisting of silica and hydroxyapatite are

formed in vitro. The concentration of Na is not reported here

as the high initial concentration of sodium in SBF does not

allow accurate measurement. Sr5 and S53P4 released iden-

tical amounts of Si ions in the solution at each time points

(Fig. 7a). Glasses with higher SrO content showed slightly

faster Si release during the first 72 h. At longer immersion

times the Si release leveled off in all glasses and the final Si

content in the SBF was identical. Figure 7b depicts the

changes in the Ca concentration of SBF. All glasses con-

taining CaO exhibited a release of Ca ions peaking in the first

6–24 h. Highest Ca concentration was measured at the

lowest immersion times and decreased in intensity with

increasing SrO content. The Ca concentration decreased

with prolonged immersion for S53P4 and Sr5, while in

immersion of the glasses with higher SrO content the Ca

concentration remained constant. The changes in the P

concentration in SBF are shown in Fig. 7c. The P concen-

tration was constant, in the accuracy of measurement, during

the first 6 h of immersion and then decreased with increasing

immersion time. The more SrO in the glass, the less the

decrease in P concentration during the immersion. Glass Sr5

behaved differently compared to all other samples. Sr5

exhibited the fastest decrease in the P concentration and P

ions were totally consumed after 168 h of immersion. For

S53P4 the P concentration was 4 mg l-1, while as much as

30.9 mg l-1 P was measured for the solution containing

Sr21.77 glass. Lastly, the concentration of Sr in SBF

increased with increasing immersion time and SrO content in

the glass (Fig. 7d).

Figure 8 presents the ATR spectra of the samples after

168 h in SBF. The spectra were normalized to the absorp-

tion band with major intensity. Major structural changes can

be seen between the ATR spectra of the untreated glass

particles (Fig. 5) and the particles immersed in SBF. The

intensities of the bands located at 748 and 930 cm-1

decreased with increasing immersion time. Simultaneously,

the broad band at 1,023 cm-1 increased in intensity,

sharpened and shifted to higher wavenumbers. A shoulder

at 1,212 cm-1 appeared after immersion and its intensity

increased with immersion time, as reported in our previous

study of glass S53P4 with MgO substituted CaO [9]. All

glasses showed similar changes but the amplitude of the

absorption peaks differed. The differences in the amplitude

may partly depend on the technique employed. In the ATR

analysis the sample is pressed against the crystal. This may

cause breakage of the upmost surface layer, thus revealing

the underneath reaction layer with another composition.

Accordingly, the amplitude of a peak partly depends on the

ratio of different surface layers analyzed.

Figure 9 presents SEM images and line scans of the

cross-sections of the glass particles immersed for 72 h in

SBF. SEM images were taken at 91,000 magnification for

Fig. 6 pH of SBF as a function of immersion time for samples with

constant mass to volume ratio (a), and constant Sa/V ratio (b)

662 J Mater Sci: Mater Med (2014) 25:657–668

123

particles of S53P4, Sr5 and Sr10 but 95,000 magnification

was used for Sr15 and Sr21.77. All glasses show a darker

and a brighter layer at the outermost surface. The thickness

of the brighter layer and the overall layer (dark ? bright

layer) decreased with increasing SrO content in the glass.

The line scan shows formation of the typical inner SiO2

rich and outermost Ca and P rich layers at the surface of the

glasses. Sr was present in the outermost layer of the glasses

with more than 5 mol% SrO.

4 Discussion

New SrO substituted bioactive glasses were prepared in

order to understand the impact of the replacement of Ca for

Sr on the thermal and structural properties of the glasses.

The impact of the SrO substitution on the reactivity in SBF

was also investigated.

4.1 Thermal, physical and structural properties

SrO and CaO are glass modifiers and are expected to

assume the same position in the glass network. The density

of the glasses increased from 2.66 to 3.03 g cm-3 with

increasing SrO content. This increase can be attributed to

the higher molecular mass of Sr compared to Ca. When

SrO was substituted for CaO the ATR spectra were similar.

The similar intensities and shapes of the bands at 930 and

1,030 cm-1 in the FTIR spectra (Fig. 5) confirm that all

investigated glasses have similar ratio between non-bridg-

ing and bridging oxygen in their network. Thus, the sub-

stitution of SrO for CaO did not markedly impact the glass

structure. In contrast, MgO was found to partly act as glass

modifier, partly as network former in MgO substituted

S53P4 [9, 24]. The band at 1,370–1,560 cm-1 is related to

carbonate species in the glass [25]. In this study the

intensity of this band did not change, thus indicating that

the concentration of carbonate was similar in all samples

[25]. The FTIR analysis suggested no structural modifica-

tions when replacing SrO for CaO. However, the greater

radius and lower filed strength of Sr2? compared to Ca2? is

likely to give a glass with weaker bonds.

Figure 1 suggests a decrease in Tg with increasing

substitution of SrO for CaO. This is in agreement with

Peterson et al. [26] who suggested that the glass transition

temperature decreases with decreasing the field strength of

the cations. The decrease in the sintering temperature Tsi

according to the HSM height curve (Fig. 2) is in agreement

with the DTA measurements (Fig. 1) and the field strength

effects. Furthermore with increasing SrO content the

change in the height of the sample between the sintering

temperature and the temperature were crystallization ini-

tiate is more pronounced. The higher decrease in theTa

ble

2M

easu

red

den

sity

val

ues

and

calc

ula

ted

mo

lar

amo

un

tso

fo

xid

esin

the

gla

sses

for

sam

ple

sw

ith

con

stan

tm

ass

tov

olu

me

rati

o(m

/V)

and

con

stan

tsu

rfac

ear

eato

vo

lum

era

tio

(Sa/

V)

Gla

ssco

de

S5

3P

4S

r5S

r10

Sr1

5S

r21

.77

Den

sity

(gcm

-3)

2.6

62

.72

2.8

22

.92

3.0

3

mg

/50

ml

SB

F7

5.0

75

.07

5.0

76

.77

5.0

79

.57

5.0

82

.37

5.0

85

.4

Mo

lN

a 2O

2.7

89

10

-4

2.7

89

10

-4

2.6

89

10

-4

2.7

49

10

-4

2.5

89

10

-4

2.7

49

10

-4

2.4

99

10

-4

2.7

49

10

-4

2.3

89

10

-4

2.7

19

10

-4

Mo

lC

aO2

.67

91

0-

42

.67

91

0-

41

.98

91

0-

42

.03

91

0-

41

.34

91

0-

41

.42

91

0-

47

.45

91

0-

58

.18

91

0-

50

0

Mo

lS

rO0

05

.91

91

0-

56

.05

91

0-

51

.14

91

0-

41

.21

91

0-

41

.65

91

0-

41

.81

91

0-

42

.29

91

0-

42

.61

91

0-

4

Mo

lP

2O

52

.11

91

0-

52

.11

91

0-

52

.03

91

0-

52

.08

91

0-

51

.96

91

0-

52

.08

91

0-

51

.89

91

0-

52

.08

91

0-

51

.81

91

0-

52

.06

91

0-

5

Mo

lS

iO2

6.6

29

10

-4

6.6

29

10

-4

6.3

79

10

-4

6.5

19

10

-4

6.1

49

10

-4

6.5

19

10

-4

5.9

29

10

-4

6.5

09

10

-4

5.6

69

10

-4

6.4

49

10

-4

pH

(16

8h

)7

.66

7.6

67

.64

7.6

67

.63

7.6

87

.65

7.7

07

.67

7.7

2

Mea

sure

dp

Ho

fS

BF

afte

rim

mer

sio

no

fea

chsa

mp

lefo

r1

68

h

J Mater Sci: Mater Med (2014) 25:657–668 663

123

sample height can suggest that the glass achieves a lower

viscosity before crystallization [27]. Accordingly, viscous

flow sintering of glasses containing SrO is easier.

Fredholm et al. [25] reported a decrease in the crystal-

lization peak temperatures in substituting SrO for CaO in

glass with composition close to 45S5. In this study the shift

of the crystallization peaks in compositions based on

S53P4 was more complex. In S53P4 two crystalline phases

were present after thermal treatment at the temperature

range of the crystallization domain, even though only one

crystallization peak could be seen in the DTA thermogram

[7].When 5 mol% of SrO was substituted for CaO the main

crystallization peak shifted to lower temperatures (Fig. 1).

For higher substitutions the main crystallization peak

shifted to higher temperatures while an additional crystal-

lization peak appeared at lower temperatures. It was

assumed that the single crystallization peak in the DTA

thermograms of S53P4 and Sr5 corresponds to the overlap

of two phases crystallizing in the temperature range of

700–950 �C. Using HSM, the onset and endset of crystal-

lization were found to decrease with increasing SrO con-

tent. This is further confirm in Fig. 1, where, the onset

(starting of the exothermic peak) of crystallization and

endset (end of the exothermic peak) is found to follow a

similar trend. However, the main crystallization peak in the

DTA thermogram shifted to higher temperatures with

increasing SrO content. This discrepancy between the

higher temperature for the maximum of the crystallization

peak and the decrease in the temperature at which the

Fig. 7 Concentrations of Si (a), Ca (b), P (c), and Sr (d) in the SBF solution for different immersion times of the experimental glasses (here Sa/V

is maintained constant between samples)

Fig. 8 FTIR/ATR spectra of the experimental glasses after immer-

sion in SBF for 168 h

664 J Mater Sci: Mater Med (2014) 25:657–668

123

Fig. 9 SEM images and line

scans across the reaction layer

formed at the surface of S53P4

(a, b), Sr5 (c, d), Sr10 (e, f),Sr15 (g, h), and Sr21.77 (i,j) after immersion in SBF for

72 h

J Mater Sci: Mater Med (2014) 25:657–668 665

123

crystallization ends in the HSM, can be attributed to a

sharpening of the main crystallization peak. The activation

energy for the main crystallization Ec was calculated using

the Kissinger equation. The activation energy decreased

with increasing SrO content. The decrease of the Ec for the

main crystallization for glasses with more than 5 mol%

SrO was assumed to depend partly on the separation of the

peak into two peaks, one minor at a lower temperature and

the main peak at a higher temperature. As the Ec calcula-

tion was based on peak maximum data, two overlapping

peaks lead to a higher value than a single peak. Heat

treatments performed on bulk glass (1 cm3) at 750 and

850 �C were used to get more information on the crystal-

lization characteristics. Before the heat treatment all

glasses were amorphous and presented no signs of crys-

tallization. After 1 h at 750 �C all glasses had partially

crystallized. Eventhough glass Sr10 does not exhibit any

crystallization peaks on the XRD pattern, significant

amounts of crystallite could be seen at the surface of the

glass. The lack of diffraction peak was attributed to the

large amorphous volume compared to the low crystal

content at the glass surface. With increasing SrO content

the intensity and number of diffraction peaks decreased and

shifted toward lower angles (Fig. 4b). Glass S53P4 forms

preferentially Na2Ca2Si3O9 crystals [6, 9]. The XRD pat-

tern of the glass Sr21.77 containing no CaO indicates

formation Na2SrSi2O6 crystals. For high Sr substitution Ca

will substitute Sr in Na2SrSi2O6 crystals as reported for

SrO substituted bioactive glass 45S5 [28]. This explains the

progressive shift of the diffraction peaks. At higher tem-

peratures S53P4 crystallizes into Ca3(PO4)2 and/or

Na2Ca4(PO4)2SiO4 [6, 7, 9]. The small shift of the dif-

fraction peaks toward lower angles when SrO content

increase up to 5 mol% SrO suggests that Sr ions substitute

the Ca in these crystals. For higher SrO contents, the peaks

corresponding to Ca3(PO4)2 decreased in intensity and new

peaks corresponding to SrSiO3 (ICDD card # 00-034-0099)

appeared. The remaining diffraction peaks were believed to

correspond to Sr3(PO4)2 and/or Si2SiO4.

Both the DTA and HSM data suggest a decrease in the

melting temperature of the glasses with increasing SrO.

This is in agreement with the lower bond strength of the

Sr–O compared to Ca–O induced by the higher ionic radius

of Sr.

4.2 In vitro properties

The influence of gradual substitution of SrO for CaO on the

dissolution and in vitro reactivity of glass S53P4 was

studied in SBF. It is well known that the surface area to

volume ratio (Sa/V) should be controlled when performing

dissolution tests [29, 30]. The surface area mainly depends

on the sample size, shape and density. However, as small

changes in glass composition affect the density only

slightly, often a constant mass to volume ratio is used for

in vitro studies of powdered fractions. In this study the

gradual substitution of SrO for CaO lead to large changes

in the density of the glass. In order to get a better under-

standing of the influence of sample amount on the in vitro

reactivity, the dissolution studies were done using (i) con-

stant mass to volume ratio and (ii) constant surface area to

volume ratio [9, 15]. In both cases, the pH of the solution

rapidly increased during the first 48 h and then stabilized at

longer immersion times. However, while no real differ-

ences between the pH of the solution can be seen in

Fig. 6a, a clear increase in the pH and thus in the disso-

lution of the glasses with increasing SrO content can be

seen for samples with constant Sa/V (Fig. 6b). With

increasing SrO content the rise in pH is steeper during the

first 48 h. In addition, the pH after 168 h of immersion was

higher for glasses with higher SrO contents. The amount of

glasses in each experiment was calculated by taking into

account the density. This experimental way of defining the

mass for each composition gives samples with almost

identical mol % of Na2O, CaO ? SrO, P2O5 and SiO2. Due

to the smaller amount of glass in samples with constant

mass to volume ratio, the increase in the pH of SBF was

slightly less than for the samples with constant Sa/V ratio.

However, both approaches showed that SrO increased the

dissolution rate in SBF. Although the highest difference in

the mass of sample per 50 ml of solution is 8 wt% (for

Sr21.77), the difference in the pH values of the buffered

immersion solution SBF is not high. However, when using

smaller fractions the differences in the pH of the solution

may be greater. For molar substitution of SrO for CaO, the

constant Sa/V gives more reliable information of the

impact of the glass composition on the glass dissolution

and reactivity.

The immersion solutions and glass samples from

experiments with constant Sa/V ratio were utilized to

measure the effect of CaO/SrO molar ratio on the disso-

lution degree of the glasses and hydroxyapatite formation

at the surface. The pH values and the steeper initial slope of

the Si concentration indicated that SrO increases the dis-

solution rate in SBF. These results are also in agreement

with previous studies demonstrating an expansion of the

structural units with progressive replacement of SrO for

CaO in silicate based glasses [15].The final Si content was

identical in all solutions suggesting that Si reaches

saturation.

The consumption of Ca and P for samples S53P4 and

Sr5 (Fig. 7b, c) can be correlated with precipitation of

hydroxyapatite at the sample surface. Interestingly, con-

sumption of Ca and P was suppressed in glasses with more

than 5 mol% SrO. For these glasses, the Sr concentration in

666 J Mater Sci: Mater Med (2014) 25:657–668

123

the immersion solution increased to relatively high levels

suggesting that formation of hydroxyapatite is limited on

glasses with high SrO content.

SEM/EDX images show reaction layers at the surface of

all glases. The line scan analyses along the arrows (Fig. 9)

indicate that the darker layer in the SEM images is com-

posed of mainly SiO2 while the bright layer consists of Ca

and P (CaP). While both layers are present in all samples,

they become thinner with increasing SrO content in the

glass. Evenmore, samples S53P4 and Sr5 show clear SiO2

rich and CaP layers, while in the other glasses the silica

rich layer also contained Ca, P and Sr. The higher the SrO

content in the glass the more Sr in the SiO2 rich layer. In

the outermost surface the content of Ca, P and Sr increased

but some Si remained. This indicates formation of a Sr

substituted (CaP) similarly to the layers that form in

phosphate glasses [24]. In general, while the initial disso-

lution seems to be increased SrO reduces the ability to form

a distinct hydroxyapatite layer but instead a thin Sr

substituted hydroxyapatite layer forms at the glass surface.

To confirm this conclusion XRD analysis was performed

on glass particles immersed for 1 week in SBF (Fig. 10).

The XRD pattern of glass S53P4 exhibit peaks at 2h =

26�, 29�, 32� and 33�. These peaks are typical of hydroxy-

carbonate apatite (ICDD card#0019-2720) and/or hydro-

xyapatite (0009-4320). With increasing the SrO content the

peak intensity decreases to almost completely disappear for

glass Sr10 and Sr15. One sharp peak can be seen on the

XRD pattern of the glass Sr21.77. The decrease in intensity

of the diffraction peak may be related to the decrease in the

layer thickness with increasing the SrO content and/or a

lower rate of transformation from amorphous CaP to its

crystalline form. Lastly, in Fig. 10, it is seen that the peak

position seem to shift to lower angles with increasing the

SrO content. This can be attributed to the partial Sr sub-

stitution for Ca in the CaP layer. Pan et al. [31] have shown

that strontium-substituted apatite precipitated from solution

only when the Sr content was greater than 0.3 mM. In our

work the Sr concentration in SBF after 168 h of immersion

was 12.4 mg l-1 (0.14 mM) for Sr5 showing no Sr in the

CaP layer, while 49 mg l-1 (0.56 mM) was measured for

Sr10 having some Sr present in the CaP layer. Despite the

high concentration of Sr in the immersion solution of

Sr21.77, Sr was not fully substituted for Ca in the CaP

layer. Interestingly, some Ca has precipitated/diffused

from the immersion solution in the surface layers of this

CaO free glass. Thus, a high Sr content in the interfacial

solution due to dissolution of glass seems not to prevent

formation of Ca containing apatite. Sr-apatite has a sig-

nificantly higher solubility than typical Ca-apatite [32].

This can be attributed to significantly higher solubility of

Sr-apatite compared to typical apatite at physiological

pH [32]. Pan et al. [31] suggest that Sr-apatite is not

thermodynamically favored and hence the Sr incorporation is

only kinetically driven in SBF with a high Sr supersaturation.

The general shape of the IR spectra giving the surface

composition is identical for all glasses, but small differ-

ences can be noticed. The shoulder at 959 cm-1 attributed

to C–O vibration modes in CO32- and to P–O–P bonding

[33, 34] of S53P4 and Sr5 shifts to 941 cm-1 for Sr10 and

then further to 930 cm-1 for Sr15 and Sr21.77. The Si–O

vibration band at 930 cm-1 may also come from the mixed

surface layer of silica and Ca–Sr–P. Furthermore, the

absorption band in the 1,360–1,500 cm-1 region remains

strong in intensity for S53P4 while it almost completely

disappears for the glasses containing SrO. This band is

usually attributed to carbonate group and may indicate that

of these glasses only S53P4 forms a carbonated hydroxy-

apatite at its surface in SBF after 168 h of immersion [23].

This suggests that SrO in the glass slows down or prevents

formation of carbonated hydroxyapatite.

5 Conclusions

Substitution of SrO for CaO in bioactive glass S53P4 did

not change the general glass structure expressed as network

forming and modifying units. The substitution induced

small changes in the thermal properties of the glasses, but

did not give essentially larger temperature windows for e.g.

fiber drawing or sintering.

In vitro reactivity in SBF studied using constant surface

area to volume ratio gave results that could be correlated to

the glass structure, while maintaining the mass of sample to

volume ratio constant did not lead to as evident pH vari-

ation as a function of SrO substitution. The initial disso-

lution increased with SrO content in the glass and lead to

slightly higher pH values of the immersion solution. Glass

Fig. 10 XRD pattern of the investigated glasses immersed for

1 week in SBF

J Mater Sci: Mater Med (2014) 25:657–668 667

123

with 5 mol% SrO formed a similar thick and distinct silica-

rich and CaP layers at the surface as the SrO free glass

S53P4. However, the higher the strontium contents in the

glass the thinner and more blurred surface layers formed at

the glass surfaces. In these glasses the Sr content gradually

decreased in the silica rich layer. Interestingly, some Ca

from the immersion solution also was found in the outer-

most surface layers. The results suggest that substituting

SrO for CaO leads to thin Sr-substituted apatite layers at

the glass surface.

Acknowledgments The Academy of Finland is gratefully

acknowledged for the financial support of J. Massera.

References

1. Xynos ID, Hukkanen MVJ, Batten JJ, Buttery LD, Hench LL,

Polak JM. Bioglass� 45S5 stimulates osteoblast turnover and

enhances bone formation in vitro: implications and applications

for bone tissue engineering. Calcif Tissue Int. 2000;67:321–9.

2. Hench LL. Genetic design of bioactive glass. J Eur Ceram Soc.

2009;29:1257–65.

3. Hoppe A, Mourino V, Boccaccini AR. Therapeutic inorganic ions

in bioactive glasses to enhance bone formation and beyond.

Biomater Sci. 2013;1:254–6.

4. Gholamzadeh B, Nobovati H. World of academy of science. Eng

Technol. 2008;42:297–307.

5. Rahaman MN, Day DE, Sonny Bal B, Fu Q, Jung SB, Bonewald

LF, Tomsia AP. Bioactive glass in tissue engineering. Acta

Biomater. 2011;7:2355–73.

6. Massera J, Fagerlund S, Hupa L, Hupa M. Crystallization

mechanism of the bioactive glasses, 45S5 and S53P4. J Am

Ceram Soc. 2012;95:607–13.

7. Fagerlund S, Massera J, Moritz N, Hupa L, Hupa M. Phase

composition and in vitro bioactivity of porous implants made of

bioactive glass S53P4. Acta Biomater. 2012;8:2331–9.

8. Hupa L, Karlsson K, Hupa M, Aro HT. Comparison of bioactive

glasses in vitro and in vivo. Glass Technol. 2010;51:89–92.

9. Massera J, Hupa L, Hupa M. Influence of the partial substitution

of CaO with MgO on the thermal properties and in vitro reactivity

of the bioactive glass S53P4. J Non-Cryst Solids. 2012;358:

2701–7.

10. Watts SJ, Hill RG, O’Donnell MD, Law RV. Influence of mag-

nesia on the structure and properties of bioactive glasses. J Non-

Cryst Solids. 2010;356:517–24.

11. Reginster JY, Seeman E, De Vernejoul MC, Adami S, Compton

J, Phenekos C, Devogelaer J-P, Diaz Currel M, Sawicki A, Go-

emaere S, Sorensen OH, Felsenberg D, Meunier PJ. Strontium

ranelate reduces the risk of nonvertebral fractures in postmeno-

pausal women with osteoporosis: treatment of peripheral osteo-

porosis (TROPOS) study. Endocr Care. (2005);90:2816–22.

12. Gentleman E, Fredholm Y:C, Jell G, Loftibakhshaiesh N,

O’Donnell MD, Hill R:G, Stevens MM. The effects of strontium-

substituted bioactive glasses on osteoblasts and osteoclasts

in vitro. Biomaterials. 2010;31:3249–52.

13. Loftibakhshaiesh N, Brauer DS, Hill RG. Bioactive glass engi-

neered coatings for Ti6Al4V alloys: influence of strontium sub-

stitution for calcium on sintering behaviour. J Non-Cryst Solids.

2010;356:2583–90.

14. Goel A, Rajagopal RR, Ferreira JMF. Influence of strontium on

structure, sintering and biodegradation behaviour of CaO–MgO–

SrO–SiO2–P2O5–CaF2 glasses. Acta Biomater. 2011;7:4071–80.

15. Fredholm YC, Karpukhina N, Brauer DS, Jones JR, Law RV, Hill

RG. Influence of strontium for calcium substitution in bioactive

glasses on degradation, ion release and apatite formation. J R Soc

Interface. 2012;9:880–9.

16. Arstila H, Vedel E, Hupa L, Ylanen H, Hupa M. Measuring the

devitrification of bioactive glasses. Key Eng Mater. 2004;

254–256:67–70.

17. Kissinger HE. Reaction kinetics in differential thermal analysis.

Anal Chem. 1957;29:1702–6.

18. Kokubo T, Kushitani H, Sakka S, Kitsugi T, Yamamuro T.

Solutions able to reproduce in vivo surface-structure changes in

bioactive glass-ceramic A-W3. J Biomed Mater Res. 1990;24:

721–34.

19. Stoch L, Sroda M. Infrared spectroscopy in the investigation of

oxide glasses structure. J Mol Struct. 1999;511–512:77–84.

20. Szumera M, Waclawska I, Olenjniczak Z. Influence of B2O3 on

the structure and crystallization of soil active glasses. J Therm

Anal Calorim. 2010;99:879–86.

21. Bentrup GJ, Moawad HMM, Santos LF, Almeida RM, Jain H.

Structure of Na2O–CaO–P2O5–SiO2 glass–ceramics with multi-

modal porosity. J Am Ceram Soc. 2009;92:249–52.

22. Serra J, Gonzalez P, Liste S, Serra C, Chiussi S, Leon B, Perez-

Amor M, Ylanen HO, Hupa M. FTIR and XPS studies of bio-

active silica based glasses. J Non-Cryst Solids. 2003;332:20–7.

23. Queiroz AC, Santos JD, Monteiro FJ, Prado da Silva MH. Dis-

solution studies of hydroxyapatite and glass-reinforced hydroxy-

apatite ceramics. Mater Charact. 2003;50:197–202.

24. Massera J, Petit L, Cardinal T, Videau JJ, Hupa M, Hupa L.

‘‘Thermal properties and surface reactivity in simulated body

fluid of new strontium ion-containing phosphate glasses’’. J Mater

Sci. 2013;24:1407–16.

25. Fredholm YC, Karpukhina N, Law RV, Hill RG. Strontium

containing bioactive glasses: glass structure and physical prop-

erties. J Non-Cryst Solids. 2010;356:2546–51.

26. Peterson IM, Tien T-Y. Thermal expansion and glass transition

temperatures of Y–Mg–Si–Al–O–N glasses. J Am Ceram Soc.

(1995);78:1979–95.

27. Vedel E, Arstila H, Ylanen H, Hupa L, Hupa M. Predicting

physical and chemical properties of bioactive glasses from

chemical composition. Part 1: viscosity characteristics. Glass

Technol. 2008;49:251–9.

28. Fujikura K, Karpukhina N, Kasuga T, Brauer DS, Hill RG, Law

RV. ‘‘Influence of strontium substitution on structure and crys-

tallisation of Bioglass� 45S5. J Mater Chem. 2012;22:7395–402.

29. Dimbley V, Turner WES. The relationship between chemical

composition and the resistance of glasses to the action of

chemical reagents. Part I. J Soc Glass Technol. 1926;10:314–58.

30. El-Shamy TM, Douglas RW. Kinetics of the reaction of water

with glass. Glass Technol. 1972;13:77–80.

31. Pan HB, Li ZY, Lam WM, Wong CT, Darvell BW, Luk KDK,

Hu Y, Lu WW. Nucleation of strontium-substituted apatite. Cryst

Growth Des. 2009;9:3342–5.

32. Pan H, Li Z, Lam WM, Wong JC, Darvell BW, Luk KD, Lu WW.

Solubility of strontium-substituted apatite by solid titration. Acta

Biomater. 2009;5:1678–85.

33. Saiz E, Goldman M, Gomez-Vega JM, Tomsia AP, Marshall

GW, Marshall SJ. In vitro behavior of silicate glass coatings on

Ti6Al4V. Biomaterials. 2002;23:3749–56.

34. Rehman I, Knowles JC, Bonfield W. Analysis of in vitro reaction

layers formed on Bioglass using thin-film X-ray diffraction and

ATR-FTIR microspectroscopy. J Biomed Res 1998;41:162–6.

668 J Mater Sci: Mater Med (2014) 25:657–668

123