Acta Biomaterialia 6 (2010) 2797–2807
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Acta Biomaterialia
journal homepage: www.elsevier.com/locate/actabiomat
Influence of design on bioactivity of novel CaSiO3–CaMg(SiO3)2 bioceramics:
In vitro simulated body fluid test and thermodynamic simulation
M.A. Sainz *, P. Pena, S. Serena, A. Caballero
Instituto de Cerámica y Vidrio, CSIC, C/Kelsen 5, 28049, Madrid, Spain
a r t i c l e
i n f o
Article history:
Received 14 September 2009
Received in revised form 30 December 2009
Accepted 5 January 2010
Available online 11 January 2010
Keywords:
Diopside
Wollastonite
Microstructure
Bioactivity
Thermodynamic simulation
a b s t r a c t
A new type of bioactive ceramic has been designed and obtained from high-temperature phase information from the wollastonite (CaSiO3)–diopside (CaMg(SiO3)2) phase equilibrium diagram. The selected
composition was that corresponding to the eutectic point of the pseudobinary CaSiO3–CaMg(SiO3)2 system. The sintering behaviour, phase evolution, microstructural changes and in vitro bioactivity of CaSiO3–
CaMg(SiO3)2 eutectic bioceramics were analysed by differential thermal analysis, X-ray diffraction, field
emission scanning electron microscopy (FE-SEM) and image analysis. A simulation of the dissolution
properties of the different materials studied, in water as well as in simulated body fluid (SBF), was also
carried out by thermodynamic calculations, with the purpose of understanding the in vitro results
obtained. The results demonstrate that the CaMg(SiO3)2 is significantly less soluble than CaSiO3, developing an in situ porous structure (biomimetic porous bone material) with adequate biodegradation rate and
stability strength when immersed in SBF. The influence of the microstructure (porosity, grain size and
phase composition) on the in vitro bioactivity of the obtained bioceramics was also examined.
Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
1. Introduction
Bioactive glasses and ceramics are considered as potential
materials as bone substitutes, because they can form a direct bond
with living bone, without the formation of surrounding fibrous tissue. The requirement of a bioactive material is that in the presence
of the human physiological environment a biological active apatite
called hydroxycarbonate apatite layer is produced on their surfaces, which provides the bonding interface with tissues as well
as bone. Several types of ceramics and glasses, such as sintered
hydroxyapatite (HA) [1], sintered b-tricalcium phosphate [2], apatite/b-tricalcium phosphate [3], apatite/wollastonite [4], wollastonite [5,6] and bioglasses [7] have been shown to bond to living
bone and they are used clinically as important bone substitutes.
The mentioned materials are not mechanically compatible with
the surrounding bone; for example bioglass has a bending strength
lower (40–45 MPa) than that corresponding to dense bone
(50–150 MPa). Nowadays the development of new bone-bonding
materials with improved mechanical properties is required.
Besides, the porosity of these materials is an important factor to
ensure the assimilation of bioceramics and their intergrowth with
bone tissue. The interconnections in a porous biomaterial are the
pathways between the pores and are the way to conduct cells
and blood vessels between pores, thus favouring bone ingrowths
* Corresponding author. Tel.: +34 917355840; fax: +34 917355843.
E-mail address: [email protected] (M.A. Sainz).
inside the ceramic. However, the fraction of large porosity degrades the strength of the material. In general, it is accepted that
is necessary to obtain a highly porous structure for tissue ingrowths. Some studies have reported notable bone ingrowths for
pores smaller that 100 lm [8–10]. Lu et al. [11] noticed that
in vivo a 20 lm interconnection size only allows cell penetration
and chondroid tissue formation; the size of the interconnections
must be over 50 lm to favour new bone ingrowths inside the
pores. Other authors [12–14] consider that only pores larger than
100–150 lm facilitate cell colonization and bone ingrowths and a
reduction in macroporosity may have a negative bearing on the
biological properties and the ingrowths in bone. Macroporosity is
conducive to osteoconduction, but also has many effects on the
mechanical behaviour of bioceramics [15], for which optimal
macroporosity parameters have not yet been defined. In general,
in the past 10 years macroporous forms of biphasic calcium phosphate ceramics prepared by a close association of hydroxyapatite
and b-tricalcium phosphate have been used for bone substitution
and dental applications [16,17]. But these structures cannot resist
severe mechanical stress without surgical stabilization. Most applications of calcium phosphate ceramics consist in filling bone
defects in those areas where mechanical stability is not a decisive
factor [18,19].
A new way of producing strong porous ceramics with interconnected porosity has been developed by De Aza et al. [20]. They
proposed to obtain eutectic structures from the binary wollastonite–tricalcium phosphate (W–TCP) system, bearing in mind the
1742-7061/$ - see front matter Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
doi:10.1016/j.actbio.2010.01.003
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different bioactive behaviour of the phases (wollastonite is bioactive and tricalcium phosphate is resorbable). The material was obtained by slow cooling from eutectic liquid which was formed by
spherical colonies composed of alternating radial lamellae of wollastonite and tricalcium phosphate. Therefore, this bioeutectic
ceramic is a bioactive material which is totally replaced by HA in
simulated body fluid (SBF). The procedure used for the synthesis
of these W–TCP potential implants restricts their size and even
their shape.
The bioactivity of wollastonite was noticed by De Aza et al.
[21,22] and the bioactivity of diopside has been also reported lately
by several authors [23,24], but in the present work it is the first
time that a new biphasic biomaterial with eutectic composition
and obtained by solid state sintering in the CaSiO3–CaMg(SiO3)2
system [25,26] is proposed as candidates for artificial bone and
dental root.
The proposed method allows obtaining bioactive materials with
variable size and shape and with controlled microstructure,
according to conventional processing methods in advanced ceramics. The advantage of this material is to combine biological activity,
control of the porosity and enough mechanical stability. It is based
on designing dense bioactive ceramics materials with the ability to
develop, when they are implanted into a living body, a porous apatite-like/diopside structure which improves the ingrowth of new
bone into implants (osseointegration).
The effect of the microstructure (porosity, grain size and structure of the phases), on the in vitro bioactivity [27] of pseudowollastonite–diopside (pW–D) and wollastonite 2M–diopside (W2M–D)
bioceramics has been studied. A simulation of the dissolution of
the different materials studied, in water and in SBF, has also been
carried out to understand the in vitro results obtained.
2. Materials and methods
2.1. Materials
Chemicals used in the synthesis of diopside and wollastonite
were high-purity amorphous silica (99.7 wt.%, Strem Chemicals,
USA); calcium carbonate (>99.5 wt.%, Merck, Darmstadt, Germany)
and MgO (>99.9 wt.%, Merck, Darmstadt, Germany). Stoichiometric
quantities of the starting powders to obtain pure wollastonite (W,
CaSiO3) and diopside (D, CaMg(SiO3)2) were attrition-milled with
PSZ–zirconia balls.
2.2. Processing methods
The pseudowollastonite (psW) was obtained by a solid state
reaction from a stoichiometric mixture of high purity milled silica
and calcium carbonate. The diopside was synthesized by solid state
reaction from a mixture of high purity silicon oxide, calcium
carbonate and MgO. In both syntheses the powders were initially
calcined at 950 °C and immediately were attrition-milled with
PSZ–zirconia balls in isopropyl alcohol media, dried at 65 °C, sieved
through 100 mesh and isostatically pressed at 200 MPa.
Larges cylinders (10 2 cm) were heated over platinum sheet
at a rate of 5 °C min1 up to 1375 °C for 6 h in the case of psW,
and at a rate of 5 °C min1 up to 1300 °C for 6 h in the synthesis
of diopside (D). The materials obtained were ground after heat
treatment, and then pressed and reheated again. This procedure
was repeated until the X-ray diffraction pattern showed as unique phases present psW and D, respectively. Then the powders
thus obtained were characterized by X-ray fluorescence (XRF).
Pseudowollastonite and diopside powders obtained with a
specific surface area of 1.11 and 1.31 m2 g1 and a mean particle
size of 5.33 and 4.75 lm, respectively, were used as raw
materials.
Stoichiometric quantities of these powders to obtain the eutectic composition (63.23 CaMg(SiO3)2 (D) and 36.77 CaSiO3 (W)
(wt.%)) of the wollastonite–diopside system [25,26] were attrition-milled in an attritor for 15 h, with PSZ–zirconia balls in isopropyl alcohol media. Subsequently, the powders were dried at
65 °C and sieved through 60 mesh. These powders were analysed
by differential thermal analysis (TG-DTA; STA-409, Netzsch), hotstage microscopy equipped with an image analysis system and
an electrical furnace (model EM 201 Leica, Germany) and X-ray diffraction (XRD; D-5000 Siemens). To determine the processes that
occur during the sinterization these powders were also characterized by specific surface area measurement (Mo. Mastersizer, Malvern) and particle size distributions in powder aqueous
suspensions (Laser Diffraction, Mastersizer S, Malvern).
The psW–D mixture green powders were thereafter isostatically
pressed at 200 MPa to form green compacts. In order to optimize
the sintering parameters of green compacts of the homogeneous
mix of wollastonite–diopside were sintered in air atmosphere, at
1250, 1300 and 1350 °C/4 h at a heating rate of 5 °C min1, and
rapidly cooled at a rate of 10 °C min1, to obtain dense compacts.
The bulk densities of the sintered bodies were determined by
water displacement method.
2.3. Materials characterization
XRD analyses were carried out (Siemens D-5000) to determine
the crystalline phases of the different wollastonite–diopside composites obtained.
The microstructure of sintered samples wollastonite–diopside
was studied on specimens polished with diamond spray down to
1 lm and chemically etched with diluted acetic acid (1:5). The different samples were studied by image analysis software (model
LEICA Quin Pro, LMIS Ltd, UK) and field emission scanning electron
microscopy (FE-SEM) (S-4700 HITACHI, Japan). Quantitative analyses were made by energy-dispersive X-ray spectroscopy (EDS)
using the ZAF (atomic number, absorption, fluorescence) correction software and theoretical internal standards. Microanalysis
data represent the average of four independent determinations.
2.4. SBF in vitro test
In order to estimate the bioactivity (potential of apatite formation) of the samples, we used the corrected SBF proposed by Kokubo et al. [28,29], the Tris-buffered SBF No. 9 (Na+ 142.0, K+ 5.0,
2
1.0 and SO2
Mg2+ 1.5, Ca2+ 2.5, Cl 147.8, HCO
3 5.0, HPO4
4
3
0.5 mol m ).
Optimized wollastonite–diopside ceramics were cut as discs
from bars obtained, and they were 2 mm in thickness and 5 mm
in diameter. After being ultrasonically washed in isopropyl alcohol,
in acetone and deionized water, they were vertically mounted on a
nylon wire in polyethylene falcon test tubes containing 100 ml of
SBF at 37 ± 0.5 °C and pH = 7.25 ± 0.2, using HCl 0.1 N for pH
adjustment. The ratio of volume of SBF to area of ceramic was
equal to 0.5 cm3 mm2.
Discs were removed at 1, 2 and 3 weeks of soaking time, gently
rinsed with deionized water and acetone, and dried in air at room
temperature.
Sample surfaces and cross-sections, before and after the exposure to the SBF, were examined by SEM at 13 keV and microanalysis (Ca, Mg, Si and P), EDS elemental maps of the cross-sections
were also obtained. Silicon, calcium, magnesium and phosphorus
ion-release profiles, in SBF at 37 °C and pH = 7.25 ± 0.2, were determined for psW–D (sintered to 1300 °C) and W2M–D (sintered to
1350 °C) samples. The SBF was removed after several periods of
immersion and silicon, calcium, magnesium and phosphorus were
determined in the removed SBF by inductively coupled plasma
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
atomic emission spectrometry (Thermo Jarrell Ash, IRIS Advantages). The samples evolution was established by measuring the
thickness of the porous layer formed in the ceramic–SBF interface
by FE-SEM.
CaMg(SiO3)2+Liq.
pW-CaSiO3+Liq.
1400
1360 ºC
2M-Ca SiO3+Liq.
W
+p
1300
2M
A simulation of dissolution behaviour of W2M, psW and D in
SBF has also been carried out by thermodynamic calculations using
the HSC software [30]. The aqueous phase was described as an
ideal solution and ionic species in solution have been selected from
database [30–35], taking into account the ionic composition of the
SBF utilized. The heat capacity values of such ions at high temperatures (T > 25 °C) have been extrapolated by the Criss–Cobble
method [36–39]. The equilibriums were calculated using the routine GIBBS in the HSC software, which uses the Gibbs energy minimization method.
Liq.
1500
Temperature (ºC)
2.5. Thermodynamic simulation
1600
1200
2M
2M-Ca SiO3+ CaMg(SiO3)2
1100
1000
0
10
20
30
CaMg(SiO3)2
40
50
(weight %)
60
70
80
90 100
CaSiO3
3. Results
Fig. 1. Phase equilibrium diagram of the join W–D of the ternary system SiO2–CaO–
MgO [26].
3.1. Synthesis of CaSiO3 and CaMg(SiO3)2
The X-ray diffraction patterns of both synthetic wollastonite
and diopside obtained showed psW and D as the only phase being
present, respectively. Both compounds were ground to an average
particle size of 5 lm. The results of chemical analysis by XRF, for
both materials, are shown in Table 1. The physical characteristics
of both powders are also shown in Table 1.
3.2. Sintering of the samples
Stoichiometric quantities of the starting powders, corresponding to the eutectic composition (63.23 (D)–36.77 (W) (wt.%))
(Fig. 1), were attrition-milled, with PSZ–zirconia balls, for 15 h,
whereby homogeneous and partially amorphous powders with
an average particle size of 0.85 lm and a specific surface area of
1.68 m2 g1 were obtained.
XRD patterns of the starting psW–D powders are shown in
Fig. 2. The broadening and low intensity of the peaks indicates that
a significant decrease in the crystallite size and a partial amorphization occurs during the large milling (15 h) if they are compared
with the unmilling ones.
The thermal behaviour of the psW–D sample was studied by
TG-DTA and plotted in Fig. 3. The results were obtained under a dynamic air flow at a heating rate of 5 °C min1, on powdered specimens, from room temperature to 1500 °C. The small weight loss
(0.5 wt.%) between 290 and 520 °C (TG) with two slight exothermic effects (347–533 °C) in DTA is associated with the elimination
of the residual water and alcohol retained in the powder surfaces
as a consequence of the milling process. The endothermic peak
with maximum observed at 1369 °C was attributed to the congruTable 1
Physico-chemical characteristics of synthetic powders.
Chemical analysis (wt.%)
CaO
SiO2
MgO
Al2O3
Na2O
K2O
TiO2
Physical characteristics
Average grain size (lm)
Specific surface area (m2 g1)
CaSiO3
CaMg(SiO3)2
48.6 ± 0.3
51.1 ± 0.3
–
0.230 ± 0.005
0.025 ± 0.005
0.012 ± 0.005
0.030 ± 0.005
26.8 ± 0.3
55.1 ± 0.3
18.0 ± 0.3
0.019 ± 0.005
0.061 ± 0.005
0.013 ± 0.005
5.3 ± 0.1
1.1 ± 0.1
4.7 ± 0.1
1.3 ± 0.1
ent fusion of the eutectic composition, in agreement with the temperature of the eutectic point established by Schairer et al. [25] for
this system. During the cooling, two exothermic peaks at 1252 and
1214 °C were observed. The first was attributed to the recrystallization of diopside and the second one at 1214 °C was assigned to
the recrystallization of wollastonite from the liquid, to form the
stable phase at low temperature (wollastonite 2M).
Taking into account these studies, psW–D dense ceramics materials were obtained after heating bars of the eutectic composition
at 1250, 1300 and 1350 °C for 4 h, where a solid state reaction process occurs.
The thermal treatments generated shrinkage in all samples, as
shown in Table 2. The tendency for the density was to increase with
the increasing amounts of shrinkage; in this way it was possible to
observe that the sample that shows highest shrinkage, corresponding to that heated at 1350 °C, showed also the highest density.
3.3. Mineralogical composition
XRD data corresponding to the heated samples are shown in
Fig. 4 and Table 2. In samples sintered at 1250 and 1300 °C the
phases detected were metastable pseudowollastonite (high-temperature polymorph of wollastonite) and diopside whereas in the
sample sintered at 1350 °C the formation of the stable low temperature polymorph wollastonite 2M (Ca0.83Mg0.17SiO3) was observed.
Up to temperatures of 1300 °C the phases detected were those corresponding to the starting materials, which indicated that the process occurs by solid state sintering in this range of temperature.
The equilibrium phase assemblage in the pseudobinary system at
temperatures lower than 1370 °C were the corresponding solid
solution of Mg in wollastonite 2M and diopside CaMg(SiO3)2 (see
Fig. 1).
The 2M ? pseudowollastonite polymorphic transition in CaSiO3
with high purity takes place at 1130 ± 5 °C. However, the presence
of Mg is solid solution in wollastonite (Ca1xMgxSiO3; 0 6 x 6 0.17)
shifts the transition temperature from 1130 ± 5 to 1370 ± 20 °C;
this solid solution explains the presence of the wollastonite 2M
in the sample sintered at 1350 °C/4 h.
3.4. Microstructural analysis
The compositions selected, sintered at 1250, 1300 and 1350 °C/
4 h, show homogeneous and fine-grained microstructures (2–
7 lm) with densities ranging from 75 to 94% (q/qth) (Table 3).
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
pW
80
70
D
Intensity
60
pW
50
D
pW
40
D
30
pW
20
D
pW
D
D
pW
10
0
10
20
30
40
50
60
70
80
2θ
Fig. 2. XRD patterns of the starting pW–D powders milled during 15 h. pW, pseudowollastonite; D, diopside.
DTA μV
TG %
2,0
200
1,5
1,0
100
0,5
0
0,0
-0,5
Table 2
Density of samples after thermal treatments.
Temperature
(°C)
Shrinkage
(%)
Density
(g cm3)
q/qth
(%)
Porosity
(%)
Phases by
XRD
RT
1250
1300
1350
0
5.9 ± 0.6
12.9 ± 0.6
16.6 ± 0.6
2.29 ± 0.05
2.32 ± 0.05
2.53 ± 0.05
2.89 ± 0.05
72
74
80
92
28 ± 3
26 ± 3
20 ± 3
8±3
psW + D
psW + D
psW + D
W2M + D
Theoretical density = 3.15 g cm3. Density of D = 3.39 g cm3; psW = 2.90 g cm3
and W2M = 2.91 g cm3.
-100
-1,0
-1,5
-200
-2,0
0
200
400
600 800 1000 1200 1400 1600
Temperature ºC
Up
DTA μV
250
Down
T=1214
200
pWollastonite + Diópside
150
100
50
T=347
0
T=533
T=1252
-50
-100
-150
T=1369
-200
-250
0
200
400
600 800 1000 1200 1400 1600
Temperature ºC
Fig. 3. Differential thermal analysis curves, during heating and cooling, of the
starting pW–D powders.
For a comprehensive microstructural characterization of the
materials, microstructural parameters such as grain size, porosity,
volume fraction and distribution of phases by image analysis and
FE-SEM-EDS, in polished cross-sections of the samples, have been
established (Table 3).
Fig. 5 shows the typical microstructure corresponding to
samples sintered at 1300 and 1350 °C. At 1300 °C, the morphology
observed corresponds to an interconnected irregular lobular structure (grey phase) and an interlocking dark grey phase (Fig. 5a). FESEM-EDS microanalysis shows that the grey phase corresponds to
psW, the dark grey phase has a composition far apart from to the
stoichiometry of D (Table 3). The porosity was quantified as
<15 vol.% and the grain size distribution gave a range from 2.3 to
6.0 lm with a d50 = 4.9 for minor grain sizes and a range from
5.4 to 11.4 lm with a d50 = 7.8 for coarse grains, both values measured by image analysis. The results indicated an increase of the
grain size in relation to the sample treated to 1250 °C (Table 3),
which is due to the grain coalescence stage in the sintering process.
The volume fraction of each phase was determined by image
analysis, yielding 62 ± 4 vol.% for the grey phase (psW) and
38 ± 4 vol.% for the dark grey (D). These results are in agreement
with the phase diagram data when taking into account the solid
solution of MgO in psW at these temperatures, where the greater
volume fraction (60%) corresponds to psW.
However, the microstructure of the sample treated at 1350 °C
was very different. In Fig. 5b, the microstructure reveals a higher
degree of sinterization than that observed at 1300 °C and it is
also possible to observe a continuous matrix which appears as
coarse grains, with a grain size distribution from 2.6 to 6.4 lm.
The EDS analysis confirmed that the matrix is wollastonite with
a little percentage of MgO and the coarse grains are diopside
with a composition (Table 3) close to the stoichiometric
(55.6 wt.% SiO2, 25.9 wt.% CaO and 18.5% MgO). The presence of
MgO in the EDS analysis of wollastonite reveals the formation
of a solid solution of MgO in this phase (Ca1xMgxO). No signifi-
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
D
120
pW
1300°C
100
D
80
Intensity
pW
60
D
DpW
D
40
20
pW
D
DDpWpW
pWD
pW
pW pW
pW
0
10
20
30
40
50
60
70
80
70
80
2θ
D
120
2M
100
1350°C
Intensity
80
D
60
D
2MD2M
40
D
2M2M
2M2M2M
2M
D
20
0
10
20
30
40
50
60
2θ
Fig. 4. X-ray diffraction patterns of the pW–D ceramics sintered at 1300 and 1350 °C. pW, pseudowollastonite; 2M, wollastonite 2M; D, diopside.
Table 3
FE-SEM-EDS microstructural parameters of sintered samples.
SEM-EDS sample
Average grain size (lm)
Fine
Coarse
SiO2
CaO
MgO
SiO2
CaO
MgO
1250 °C
1300 °C
1350 °C
Theoretical
MgO solid solution
2.1 ± 0.5
4.8 ± 0.5
–
4.8 ± 0.5
7.8 ± 0.5
4.5 ± 0.5
58 ± 6
51 ± 5
54 ± 5
51.72
56.8
42 ± 6
46 ± 5
43 ± 5
48.28
37.8
n.d.
3±2
3±2
–
5.4
52 ± 5
54 ± 5
60 ± 5
55.6
40 ± 5
36 ± 5
24 ± 3
25.9
8±5
10 ± 3
16 ± 3
18.5
EDS microanalysis (wt.%)
CaMg(SiO3)2 dark grey phase
EDS microanalysis (wt.%)
Ca1xMgxSiO3 grey phase
cant glassy phases on grain boundaries or impurity concentration, at the level of resolution employed (FE-SEM-EDS), was
found in any sample.
The above results indicate that an increase in the temperature
of thermal treatment causes the densification of the samples, especially in the sample treated at 1350 °C, in which the process is assisted by the formation of small amounts of liquid phases due to
the minor impurities present in the raw materials used. This
produces a different microstructure and a decrease in the porosity
in these samples (Fig. 5b).
3.5. In vitro bioactivity
In order to study bioactivity, the wollastonite–diopside composites sintered at 1250, 1300 and 1350 °C/4 h were soaked in
SBF at 37.0 ± 5 °C for 7, 14 and 21 days.
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
Fig. 5. FE-SEM images of the sample treated at (a) 1300 and (b) 1350 °C, after chemical etching with dilute acetic acid.
3.5.1. Pseudowollastonite–diopside ceramics
FE-SEM micrograph of the polished cross-section of the psW–D
(sintered at 1300 °C/4 h) sample after soaking in SBF for 7 days is
shown in Fig. 6. The surface of the sample was eroded by dissolution
of the CaSiO3 grains in the SBF during the period of immersion, forming
a porous structure layer parallel to the SBF–ceramic interface. The
depth of the porous layer at this time was 94 lm. Fine small agglomerate particles were observed inside the pores formed on the surface of
the sample (Fig. 7). These small agglomerate particles were determined to be bone-like apatite although from SEM-EDS microanalysis
the Ca/P ratio of the apatite-like layer was 2.3, higher than that in
hydroxyapatite. This fact suggest that carbonate hydroxyapatite is
forming on the surface (Ca/P ratio value greater than pure hydroxyapatite) according to the results reported by other authors [40–42].
respectively. After immersion, the material surface was covered
by a layer of small agglomerates of globular particles with agglomerate sizes of 60, 70 and 125 nm in diameter corresponding to
3.5.2. Wollastonite 2M–diopside samples
Fig. 8 shows SEM images of the surfaces of W2M–D sample (sintered at 1350 °C/4 h) after soaking in SBF for 1, 2 and 3 weeks,
Fig. 6. SEM cross-section view of the pseudowollastonite–diopside sample
immersed in SBF solution for 7 days.
Fig. 7. SEM surface view of the pseudowollastonite–diopside sample treated at
1300 °C immersed in SBF solution for 7 days.
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
treatment for 7, 14 and 21 days. These small agglomerates of particles formed on the surface grew with immersion time in the SBF,
as can be clearly seen in Fig. 8.
The corresponding EDS analysis shows that the Ca/P ratio does
not always correspond to the hydroxyapatite ratio, which could be
indicative of the existence of other calcium phosphate salts. The
apatite-like layer presented Ca/P ratio 1.68 in the zone to get
away from the surface (which was similar to that of the stoichiometric HA (1.67)), while in the zone near to the surface in contact
with SBF the Ca/P ratio increased up to 2.3, suggesting the formation of carbonate hydroxyapatite. The CDHA formation is a consequence of the usually high levels of calcia in the solution near on
the surface, due to dissolution of wollastonite, and the enrichment
with carbonate ions present in SBF to give rise to carbonate
hydroxyapatite on the surface.
In Fig. 9 the release profiles of Si, Ca, Mg and P ions in SBF at
37 °C for psW–D and W2M–D samples are shown. Both ceramic
materials released Si, Mg and Ca ions, and removed P ions from
SBF in a similar way, but in different amounts. The specific concentration of Ca in SBF increased from 148 to 205 and 195 mg l1 cm2
for psW–D and W2M–D, respectively, when soaked at 37 ± 0.5 °C
for 21 days. The continuous Ca2+ ion dissolution from the wollastonite (principally) and diopside materials contributed to generate
an increase in the Ca2+ ion concentration in the SBF even though
the formation of Ca–P phase. In addition, SBF was enriched in Si
up to a value of 69 and 49 mg l1 cm2 for psW–D and W2M–D,
respectively, for the same incubation period. The concentration
of Mg in SBF slightly increased from 64 to 69 mg l1 cm2 for both
materials, indicating some partial dissolution of diopside. On the
other hand, phosphorous ions were removed from SBF by both sintered materials, because the formation of Ca–P phase on the surface of the diopside grains and its concentrations decreased from
69 to 42 and 54 mg l1 cm2 for psW–D and W2M–D, respectively.
After 14 days in soaking, the formation rate gradually slowed down
due to depletion of P and Ca ion concentration from the SFB solution and diffusion processes of ions across the new apatite-like
layer, which becomes very difficult and completely stops the apatite-like formation process. Other authors [43,44] have reported
that these ions produce osteostimulation when presented at a
particular ratio of ions and in a particular concentration range
(15–30 ppm Si and 60–90 ppm Ca). These ion ratios and concentrations are achieved for wollastonite–diopside materials studied.
SEM photograph and EDS analyses of polished cross-sections of
W2M–D sample (sintered at 1350 °C) after soaking in the SBF solution at different times are shown in Fig. 10. The surface of samples
was partially dissolved in the SBF, forming a porous layer parallel
to the SBF–ceramic interface. The depths of the porous layers were
86 lm, 100 and 105 lm after 1, 2 and 3 weeks, respectively. The
EDS microanalyses established that the dissolved phase was W2M,
whereas D remained as an interconnected porous structure. The
porosity was produced by the preferential dissolution of W2M
grains in the SBF solution. In Fig. 10 it is also possible to observe
a densified zone, enriched in silicon, which could be produced by
the low diffusing rate of the silicon ions to the surface. Similar effects were observed in all samples studied.
2
Elemental ionic concentrations (mg/L cm )
200
Si pW-D
CapW-D
PpW-D
Mg pW-D
Si W2M-D
Ca W2M-D
P W2M-D
Mg W2M-D
150
100
50
0
0
7
14
21
Soaking time (days)
Fig. 8. SEM surfaces views of wollastonite 2M–diopside (sintered at 1350 °C/4 h)
samples after soaking in SBF solution for (a) 7, (b) 14 and (c) 21 days.
Fig. 9. ICP element concentrations results, in SBF at 37 °C and pH 7.25 ± 0.2, caused
by immersion of pW–D and W2M–D ceramic as a function of soaking time.
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
Fig. 10. SEM micrographs of the cross-sections of SBF wollastonite 2M–diopside sample (sintered at 1350 °C/4 h) interfaces after soaking for (a) 7, (b) 14 and (c) 21 days.
In Table 4 we can see the depth of penetration values obtained
from the different samples and different times of soaking in SBF;
the measurements were made on SEM micrographs. These results
indicate that psW is more easily reabsorbed that W2M and D and
that the material with more initial porosity presents more reactivity (bioactivity). Moreover, the highest interaction occurs in the
first week, and during the second one the reaction only advanced
a few micrometers, being finished practically at the end of the third
week.
Fig. 11 shows the relationships between soaking time and thickness of the porous layer developed on the surface of the ceramic
substrates. The penetration depth as a function of the soaking time
reveals that kinetics of the formation of the porous layer followed a
logarithmic law. At the beginning, the porosity of the microstruc-
Table 4
Penetration depth of the interaction developed into the surface for the different
samples by immersion in SBF.
Time of immersed in
SBF (weeks)
pW–D
1250 °C (lm)
pW–D
1300 °C (lm)
W2M–D
1350 °C (lm)
1
2
3
90 ± 2
115 ± 2
118 ± 2
94 ± 2
108 ± 2
110 ± 2
86 ± 2
100 ± 2
106 ± 2
ture facilitates ion exchange with the solution and resulted in
faster dissolution and formation of apatite. Afterwards, the differences observed between the samples studied can be explained in
terms of the differences in solubility and initial porosity. The formation of a porous layer in the surface is attributed to a higher rate
of dissolution of psW and W2M vs. diopside on the psW–D initial
surface ceramics.
3.6. Thermodynamic simulation
A simulation of dissolution of psW, W2M and diopside was estimated by thermodynamic calculations using the HSC software
[30]. The composition of the original simulated SBF has been obtained using as raw materials: NaCl, NaHCO3, KCl, K2HPO4,
MgCl26H2O, CaCl2, Na2SO4 while HCl(a) was added to fit the initial
pH to 7.25. The species considered in the modelization of the aqueous phase were: H2O, Cl(), HCO3(a), HPO4(2a), K(+a), Mg(+2), Na(+a),
OH(a), H(+a), CaOH(+a), MgOH(+a), HSO4(a), SO4(2a), HSiO3(a),
SiO4(4a) and electrons (e) to maintain the electronic neutrality
of the system.
The result of the simulation of dissolution in 1 L of SBF of psW,
W2M and D can be observed in Fig. 12a–c, respectively. According
to the figure, the limit of solubility in SBF at 37 °C of the phases can
be set at 3.4 105 g l1 for psW (pH 10.3), at 1.9 105 for
M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
2805
Penetration depth (microns)
120
80
pWD 1250
pWD 1300
W2MD 1350
40
0
0
5
10
15
Soaking time (days)
20
25
Fig. 11. Thickness change of the formed porous layer as a function of soaking time
in SBF.
W2M (pH 10.1) and at 1.8 106 g l1 for D (pH 9.3), so that
D is by far the less soluble phase, and psW the most soluble phase.
On the other hand, all the phases produce a basic pH in the equilibrium which decreases the solubility limit of HA and facilitates the
formation of HA. The simulated solubilities are in agreement with
the experimental results obtained in this work. This could confirm
the differences observed in the dissolution rate between the various phases studied.
4. Discussion
The biphasic compounds studied were constituted by two crystalline phases (as determined by XRD and FE-SEM-EDS microanalysis), pseudowollastonite–diopside and wollastonite 2M–diopside,
as a function of sintering temperature, in agreement with the temperature of the eutectic point established by Schairer et al. [25].
Moreover, by controlling the microstructure and composition, we
were able to design porous materials for psW–D samples treated
at 1250 and 1300 °C and dense materials for W2M–D treated at
1350 °C.
The in vitro SBF studies showed the formation of a porous surface structure constituted by the remaining diopside phase with
small globules of apatite-like particles (CDHA (Ca10x(HPO4)x
(PO4)6x(OH)2x (0 < x < 1)) inside the pores. These particles show
different values of the Ca/P ratio because it is possible that initially
in the interface SBF/material the pH conditions were the required
for the formation of CDHA (pH 6.5–9.5). Subsequently, the pH conditions at the interface SBF/CDHA can allow the formation of HA
(pH 9.5–12).
The interaction with simulated body fluid is higher for the porous microstructure, pseudowollastonite–diopside obtained at
1300 °C/4 h, than for the dense microstructure wollastonite 2M–
diopside obtained at 1350 °C/4 h. Moreover, the simulation of the
solubility calculated using thermodynamic functions has allowed
establishing, in the present work, that the order of solubility in
SBF–water is: psW > W2M > diopside as experimentally observed.
Taking into account both factors (initial porosity and phase solubility) the reactivity of the samples in SBF solution was the following:
psW 1250 °C > psW 1300 °C > W2M 1350 °C.
The results evidence that the microstructure morphology generated in the simulated physiological media was controlled by the
solubility of the different phases, which generates changes of the
surface chemistry and the surface topography. This mechanism
leads to the in situ formation of an interconnected porous struc-
Fig. 12. Thermodynamic simulation of the dissolution behaviour of (a) pseudowollastonite, (b) wollastonite 2M and (c) diopside in 1 L of SBF at 37 °C. The
composition of the aqueous phase has been represented as a function of the solid
phase added.
ture of diopside which could be expected to promote bone
growth. The phase with the highest solubility (wollastonite) controls the formation of the HA layer on the surface of D, which is
the phase that produced the structural support of the
biomaterials.
The mechanism of the formation of the porous diopside/apatitelike layer on diopside–wollastonite ceramics can be summarized in
the following steps:
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M.A. Sainz et al. / Acta Biomaterialia 6 (2010) 2797–2807
– Dissolution of wollastonite in SBF with release of Ca2+ and HSiO3
as majority ions, and increase in Ca2+, HSiO3 , OH, ionic activities at the neighbourhood of the reacting surface until they
exceed the solubility product of the HA:
CaSiO3 ðsÞ þ H2 O ! Ca2þ ðaq:Þ þ HSiO3 ðaq:Þ þ OH ðaq:Þ
With the release of the calcium ions from the wollastonite solid
solution, many Si–OH groups are formed on the surfaces of
W–D composites. These silanol groups induce heterogeneous
nucleation of apatite, and the released calcium ions increase the
ionic activity product of apatite, enhancing apatite nucleation.
– Partial dissolution of diopside (Figs. 9 and 10) in SBF with
release of Ca2+, Mg2+ and HSiO3 :
CaMgðSiO3 Þ2 ðsÞ þ 2H2 O ! Ca2þ ðaq:Þ þ Mg2þ ðaq:Þ
þ 2HSiO3 ðaq:Þ þ 2OH ðaq:Þ
An ionic exchange of Ca2+ from diopside network for 2H3O+ from
SBF and only in some regions a silica layer was found.
– Next the nucleation of a calcium-deficient hydroxyapatite layer
takes place on the porous surface of the material, by reaction of
the SBF phosphate ions with the excess of calcium ions liberated
to SBF by the ceramic. The silicate ions liberated by the ceramic
can produce a silicon hydroxyapatite:
4
2þ
ð6 xÞHPO2
ðaq:Þ þ 8OH ðaq:Þ
4 ðaq:Þ þ ySiO4 ðaq:Þ þ 10Ca
calculations which allowed established the mechanism of bioactivity in these materials. Controlled rates of dissolution of these biphasic bioceramics were possible to design a porous morphology
similar to porous bone with adequate biomechanical properties.
The in situ formation of an interconnected porous apatite-like/
diopside layer at the ceramic–SBF interface was controlled by the
high solubility of wollastonite in SBF. This porous structure generated is a three-dimensional architecture ceramic matrix and is expected to be useful for the invasion of cells and new bone
formation.
By means of processing technology, it is possible to tailor ceramic biomaterials with controlled microstructures and higher porous sizes in order to improve their osteointegration.
Acknowledgements
This work was funded by MICINN (Ministry of Science and
Innovation, Spain) through research projects MAT2006-12749C02-01-02 and MAT2007-65857. Financial support of program
MP-1732-mf-2008 is also recognized. We are very grateful to
Professor S. De Aza for his interesting comments on our work.
Appendix. Figures with essential colour discrimination
Certain figures in this article, particularly Figures 1–4, 6, 9, 11,
12 are difficult to interpret in black and white. The full colour
images can be found in the on-line version, at doi:10.1016/
j.actbio.2010.01.003).
! Ca10x ðHPO4 Þx ðPO4 Þ6y ðSiO4 Þy ðOHÞ2x ðsÞ þ 6H2 O
References
Data from FE-MEB-EDS suggest a very small quantity of silicon
substitution on the apatite phase according to 40.45/21.92/0.18
(Ca/P/Si atomic ratio).
– Once apatite nuclei are formed on the surface of the porous
layer, they can grow spontaneously by consuming calcium and
phosphate ions from the surrounding simulated body fluid.
The potential of these materials was to combine biological
activity with adequate mechanical stability and controlled porosity. The mechanical stability of the composition study (63.23
CaMg(SiO3)2 (D) and 36.77 CaSiO3 (W) (wt.%)) can be estimated.
An elastic modulus of 130 GPa has been calculated, taking into
account the elastic modulus of starting materials (diopside
170 GPa and wollastonite 75 GPa) and the porosity. This value is
higher than the reported for other bioceramics (BioglassÒ 35 GPa,
Glass–ceramic A-W 118 GPa, HAp 47 GPa) [45]. The interconnected
porous structure generated (self-assembled into a three-dimensional architecture creating a ceramic matrix) is expected to be
useful for the invasion of cells and new bone formation when the
ceramics are implanted into bone defects. The ionic dissolution
products (Si and Ca ions) showed a range of concentration that will
lead to enhanced proliferation of osteoblast (osteostimulation).
5. Conclusions
Pseudowollastonite–diopside and wollastonite 2M–diopside
compacts with different porosities were obtained by a carefully
controlled process, attrition-milling followed by heating of
homogeneous compacts, of fine and pure CaMg(SiO3)2 and CaSiO3
synthetic powders.
The results obtained indicated that CaSiO3–CaMg(SiO3)2 ceramics present high reactivity in SBF and might be used as bioactive
implant materials. A simulation of dissolution behaviour of W2M,
psW and D in SBF, has been also carried out by thermodynamic
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