bioactive glass nanocomposite foams on their bioactivity and mechanical properties

Materials Research Bulletin 47 (2012) 3523–3532

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Effect of the composition of hydroxyapatite/bioactive glass nanocomposite foams on their bioactivity and mechanical properties H. Ghomi *, M.H. Fathi, H. Edris Biomaterials Group, Department of Materials Engineering, Isfahan University of Technology, Isfahan 8415683111, Iran

A R T I C L E I N F O

A B S T R A C T

Article history: Received 24 September 2011 Received in revised form 1 June 2012 Accepted 26 June 2012 Available online 2 July 2012

Nanocomposite foams were prepared on addition of 63S bioactive glass (BG) to pure hydroxyapatite (HA), in proportion of 0, 25, 50, 75, and 100 wt.%, and their mechanical properties and bioactivity were compared. The prepared nanocomposite foams have a grain size in the range 24–38 nm and pore size in the range 100–400 mm. The compressive strength and elastic modulus increased with increasing the amount of BG addition up to 25 wt.% and then decreased by more addition of BG. The maximum values of compressive strength and elastic modulus were found to be about 2.78 MPa and 219 MPa, respectively. The mean values of total and interconnected porosity were calculated in the range 84–88% and 57–76%, respectively. The obtained composite foams have chemical composition similar to the mineral phase of bone and by changing the ratio of HA/BG it can reach the appropriate bioactivity and biodegradability level needed for different biomedical applications. ß 2012 Elsevier Ltd. All rights reserved.

Keywords: A. Nanostructures B. Sol–gel chemistry C. Electron microscopy C. X-ray diffraction D. Mechanical properties

1. Introduction The new challenge in biomaterials is to enhance the body’s own regenerative capacity by stimulating genes that initiate repair at the site of damage or disease. Third generation bioactive glasses (BGs) and macroporous foams are being designed to activate genes that stimulate regeneration of living tissues [1]. Highly porous bioceramic scaffolds (foams) supply a framework for enhanced cell infiltration and migration throughout the scaffold [2], and act as a template for bone growth in three dimensions [3]. Cell spreading and proliferation with bone progenitors were capable of filling 400 mm pores within two weeks [2]. There are many criteria for an ideal tissue-engineering scaffold. It has been recognized that the structure of the scaffold should consist of a highly interconnected porous network. Interconnections greater than 50 mm in diameter [4] and pores larger than 100 mm in diameter [5] should allow for cell penetration, tissue ingrowth, vascularization, and nutrient delivery to the regenerating tissue and assure mineralized bone formation. The mechanical properties of the scaffold should be sufficient to provide mechanical stability in load bearing sites prior to regeneration of new tissue [1,3,6]. Bioactive ceramics, such as BGs and hydroxyapatite (HA) have been developed over the last two decades. Their accomplishments in the field of biomedical applications have attracted wide

* Corresponding author. Tel.: +98 910 3141226; fax: +98 311 3912752. E-mail addresses: [email protected], [email protected] (H. Ghomi). 0025-5408/$ – see front matter ß 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.materresbull.2012.06.066

attention. HA is characterized by its high biocompatibility and close chemical similarity to biological apatite present in hard human tissues [7]. However, although HA is bioactive, its reactivity with existing bone, and the rate at which bone apposes and integrates with HA is relatively low [8]. BGs are more reactive, degradable, osteoconductive, and show better bioactivity than HA [7,9]. Recent studies showed that the degradation products of BGs could stimulate the production of growth factors, cell proliferation and activate the gene expression of osteoblast [10]. In addition, bioactive glass is the only one, which could bond in a few hours to hard and soft tissue [11]. The limiting factor in the use of BGs as tissue engineering scaffolds is the inherent brittleness of glass [1]. If HA and BG are to be combined in an optimized tissue engineering scaffold, then the designed composite offers an exceptional opportunity. It allows for the creation of bioresorbable and bioactive scaffolds with tailored physical and mechanical properties, and similar chemical composition to biological apatite present in hard human tissues. Moreover, the designed composite can be engineered in such a way that its bioresorption rate in the human body matches the formation rate of new tissue. Nowadays, there are different technologies for manufacturing highly porous body of bioceramics. Important advantages have been found in the production of macroporous bioceramics through the route of gelcasting of foams [12–14]. Gelcasting method is an appropriate technique for fabrication of ceramic materials where reliability is required for fabrication of components with high strength, good dimensional accuracy, uniform structure, and varying thickness or complex shapes [15–17]. In this process, foaming is conducted by addition of a foaming agent and vigorous

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agitation of aqueous ceramic suspensions. Then, in situ polymerization of an organic monomer [12,16] or gelation of a gelling agent [13,15,18] causes to form a sufficiently strong gel to withstand the body weight, even at the typically low solids loading used in such suspensions. In the early gelcasting systems, acrylamide monomers were used as gelling agents in organic solvents. Owing to the toxicity of these systems, other gelling systems using proteins and naturally occurring polysaccharides such as agarose have been evaluated for the gelcasting of porous foams [13,15]. Recent studies have shown that increasing specific surface area and pore volume of BGs might greatly accelerate the deposition process of HA. Webster et al. revealed that the biomaterials in nanoscale could stimulate the reaction between materials and cells. It was demonstrated by in vitro tests of nano bioceramics that osteoblast proliferation and long-term functions could be enhanced when the size of bioceramics grains or fibers were less than 100 nm. Furthermore, nanomaterial was one of the components of natural bones, so nanomaterials have been paid much attention in bionics and bone tissue engineering field [19]. In our previous work, bioceramic foams consisting of 50 wt.% HA and 100 wt.% HA were prepared via gelcasting method [20,21]. The aim of present work was fabrication, in vitro characterization, and optimization of the compressive strength of nanostructure bioceramic composite foams consisting of different degrees of HA and BG. For this purpose, HA and BG nanopowders were synthesized by the sol–gel method and porous body of HA/BG composite was fabricated by the gelcasting technique using agarose as a gelling agent. 2. Experimental procedure 2.1. Starting materials In this study, tetraethyl orthosilicate (TEOS, Merck), triethyl phosphate (TEP, Aldrich), calcium nitrate tetrahydrate (Ca(NO3)24H2O, Merck), phosphoric pentoxide (P2O5, Merck), absolute ethanol (C2H5OH, Merck), and hydrochloric acid (HCl, Merck) were used as HA and BG precursors. HA and BG nanopowders were prepared based on the procedures used and reported before at the Biomaterials Group, Department of Materials Engineering, Isfahan University of Technology [22,23]. In addition, agarose powder (Merck), Tergitol Np-9 (Aldrich), and tripoly phosphate sodium (TPP) were used as the gelcasting method components. 2.2. Preparation of HA nanopowder The sol–gel method was used for preparing the HA nanopowder as follows: first, the appropriate amounts of Ca(NO3)24H2O and P2O5 were separately dissolved in ethanol to form 1.67 and 0.5 mol/l solutions, respectively. The prepared solutions with Ca/P molar ratio of 10:6, which is the observed Ca/P ratio in hydroxyapatite, were mixed together. The mixture was then stirred at room temperature using a magnetic stirrer for 24 h to obtain a clear gel. The prepared gel being aged at ambient temperature for 24 h, was dried at 80 8C in an electrical air oven for 24 h. Heat treatment of prepared powders was carried out at 600 8C for 20 min in a muffle furnace by a heating rate of 5 8C/min. 2.3. Preparation of BG nanopowder The composition of the BG nanopowder used in this study belongs to the system CaO–SiO2–P2O5, 63S with 65% SiO2, 31% CaO, and 4% P2O5 in molar percentages. Proper amounts of TEOS (as a silica precursor), 2 N HCl (as a catalyst), and deionized water were dissolved in ethanol and stirred at room temperature for 30 min.

Then, TEP was dissolved into the acid silica sol. After 20 min of stirring, Ca(NO3)24H2O was added. After that, the solution was stirred for an additional 60 min. The clear solution was placed in an electrical air oven for aging at 60 8C for 48 h. The aged gel was then dried at 120 8C for 48 h to obtain a white powder. The obtained powder was calcined at 600 8C for 2 h in a muffle furnace by a heating rate of 10 8C/min. 2.4. Preparation of HA/BG composite foams Different powder compositions were prepared on addition of BG nanopowder to pure HA, in proportion of 0, 25, 50, 75, and 100 wt.% by ball milling (B/P: 5/1, rotational speed: 175 rpm, and time: 45 min). The prepared powders, by means of 1 wt.% TPP as a dispersant with regard to solids loading, were dispersed in deionized water by ultrasonic and mechanical agitation, to reach a solid loading of 60 wt.%. Simultaneously, agarose solutions (7 wt.%) as a gelling agent were prepared by mixing agarose powder with deionized water under stirring using a magnetic stirrer and heating up to 130 8C in a sealed beaker. The agarose solution was added to the stirring slurry of the powders at 80 8C to obtain a suspension with 50 wt.% powders and 1.2 wt.% active gelling agent based on water. The prepared slurry was foamed by vigorous agitation using a triple-blade mixer at 80 8C with the addition of 1.5 vol.% Tergitol as a surfactant. The foamed slurry was then poured into the mold and the gelling reaction was catalyzed by cooling the foam to 0 8C. The gelation process provides permanent stabilization for the bubbles. The samples were then de-molded, dried at room ambient, and thermally stabilized at 900 8C for 4 h. The obtained foams with 50 wt.% BG were sintered at different temperatures to evaluate the effect of sintering temperature. 2.5. Materials characterization 2.5.1. Specific surface area The specific surface area of the prepared powders and composite foams were calculated from the N2 adsorption isotherms using the multipoint Brunauer–Emmett–Teller (BET) technique (Micromeritics Instrument Corp., Gemini). The average particle sizes of the prepared powders, assuming that the synthesized particles were spheroids, were calculated as shown in Eq. (1) [24]: d¼

6  103 SBET  r

(1)

where d is the average particle size (nm), SBET is the specific surface area (m2 g1), and r is the theoretical density (g cm3) of the synthesized powders. Textural pore size distribution and the modal textural pore diameter were derived using the Barrett–Joyner–Halenda (BJH) method on isotherms obtained from N2 adsorption analysis. 2.5.2. X-ray diffraction analysis In order to analyze the phase structure of the prepared powders and composite foams, X-ray diffraction (XRD, Philips Xpert) analysis with Cu Ka radiation at 40 kV and 40 mA, with a step size of 0.058 and a count rate of 1 s per step, was performed and the diffraction patterns were recorded from 208 to 708. The obtained experimental patterns were compared to the standards, compiled by joint committee on powder diffraction and standards (JCDPS). The crystallite sizes of the initial powders and prepared foams were determined by Scherrer’s equation [25]:

b¼

K l t  cos u

(2)

H. Ghomi et al. / Materials Research Bulletin 47 (2012) 3523–3532

where b is the full-width at half maximum intensity (rad), l is the wavelength (l = 0.154056 nm), u is the Bragg’s angle (8), K is a constant (K = 0.89), and t is the apparent crystallite size (nm). 2.5.3. Scanning and transmission electron microscopy analyses Scanning electron microscopy (SEM) evaluations were performed using a Philips XL30 to investigate the morphology and size of the pores before and after immersing the composite foams in the simulated body fluid (SBF). Image analysis method was used for the measurement of pore size distribution. The diameters of minimum 200 pores were measured and the pore size distributions were calculated. Transmission electron microscopy (TEM, Philips CM-120) was used to study the morphology and particle size of the prepared powders. 2.5.4. Porosity measurement In order to determine the porosity of the prepared composite foams, three samples of the foams were selected for each sintering temperature and each composition and the porosity of the bulk specimens was measured by Archimedes method. The apparent porosity, which measures the interconnected porosity, was determined by weighing the dry ceramic (Wd), and then reweighing the ceramic both when it is suspended in water (Ws) and after it is removed from the water (Ww). Then   Ww  Wd (3)  100 Apparent porosity ¼ Ww  Ws The true porosity includes both interconnected and closed pores. The true porosity, which better correlates with the properties of the ceramic, is:   rB True porosity ¼ (4)  100

r

where B¼

Wd Ww  Ws

(5)

B is the bulk density and r is the theoretical density or specific gravity of the ceramic. The bulk density is the weight of the ceramic divided by its volume [26]. The percentages of shrinkage of the foams upon drying and sintering were analyzed to determine the shrink range of the foams after drying and sintering. The dimensions of three samples of the foams were measured before and after drying and sintering. This study was done on cylinder shape samples where the height and diameter were measured using a digital caliper rule. The percentage of shrinkage was measured as shown in Eq. (6): % Shrinkage ¼

L1  L2  100 L1

(6)

where L1 and L2 are the lengths before and after drying or sintering process, respectively. Furthermore, to measure green density, mass of each foam (m) was measured in a digital balance after drying and the value of green density (rg) was determined as can be seen in Eq. (7).

rg ¼

m

pr 2 h

(7)

where r is the radius (cm) and h is the height (cm) of the foam after drying. 2.5.5. Mechanical testing To measure the compressive strength of the porous samples, a crosshead speed of 0.5 mm/min was used and the compressive tests were performed on cylindrical samples (10 mm in diameter

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and 20 mm in length) using a universal testing machine (zwick, material prufung, 1446–60) with 10 kN load cell. The compressive strength was estimated from the maximum load registered during the test divided by the original area. The elastic modulus was calculated as the slope of the initial linear portion of the stress– strain graphs. The results were finalized based on an average of five specimens of each type of the foams. 2.5.6. In vitro evaluation To evaluate the bioactivity and bioresorbability of the prepared foams, the standard SBF was prepared according to Kokubo’s protocol by dissolving appropriate amounts of the relevant reagent grade chemicals in distilled water and the samples were immersed in the prepared solution [27]. Each prepared specimen was placed in one sterilized polyethylene bottle and SBF with a solid/liquid ratio of 10 mg/ml was added to these bottles. The bottles were held in a water bath at 36.5  0.5 8C for various immersion times (7, 14, 21, and 28 days) without refreshing the soaking medium. The specimens were then removed from the solutions, rinsed with deionized water, and dried at room condition. During this time (28 days), the changes in pH value of the solutions were evaluated by a pH meter. The concentration of the calcium and phosphorous ions in the filtered solutions were determined by inductive coupled plasma optical emission spectroscopy (ICP-OES, Perkin Elmer). The bone-like apatite formation on the surface of the samples and pores filled by formation of apatite were investigated by SEM technique. The Fourier transform infrared spectroscopy (FTIR, Bomem, MB-100) in the range of 400–4000 cm1 was utilized to confirm the formation of apatite layer on the samples. 3. Results and discussion The specific surface areas obtained from BET plots were found to be 99.3 and 189.7 m2 g1 for HA and BG powders with corresponding particle sizes of 19 and 11 nm, respectively. Particle sizes obtained from the results of image analysis of TEM micrographs also confirmed the above results. Fig. 1 shows TEM micrographs of the prepared BG and HA powders. TEM micrographs illustrate that the particle size is smaller than 30 nm. The XRD patterns of the starting HA and BG powders, composite foams with 50 wt.% BG at different sintering temperatures, and foams with 100 and 0 wt.% BG sintered at 900 8C are shown in Fig. 2. Good agreement was found between the XRD pattern of the prepared HA powder (Fig. 2(c)) and the stoichiometric HA [28]. However, the XRD pattern of the HA foam (foam with 0 wt.% BG) sintered at 900 8C showed the partial decomposition of HA to beta tricalcium phosphate (b-TCP) phase. The spectrum of the foam with 100 wt.% BG sintered at 900 8C (Fig. 2(b)) showed the peaks that were indicative of Larnite (Ca2SiO4) [29]. It should be noted that after heat treatment at 900 8C, the crystallinity of BG has been increased. This is due to the formation of crystalline phase (Larnite) at approximately 900 8C. Crystallization of the BG will affect its bioactivity and resorbability [30]. However, the BG reinforced HA composite foam sintered at 900 8C contained a HA phase and variable amounts of b-TCP and amorphous phases, depending on the amount of BG added and did not have any peaks of Ca2SiO4. This is in agreement with the results of other researchers and indicates that BG behaves more as a sintering aid and promotes the conversion of HA to b-TCP [31]. Fig. 2(e)–(g) shows XRD spectra for composite foams with 50 wt.% BG sintered at 700, 800, and 900 8C. Spectra of foams sintered at 700 and 800 8C showed that there was no phase transformation detected when the sintering temperature was less than 900 8C. However, when the sintering temperature was as high as 900 8C, the X-ray analysis showed the presence of b-TCP phase, which indicates the decomposition of HA and BG. The presence of b-TCP is very important, because it is a very soluble

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Fig. 1. TEM micrographs of the prepared (a) BG and (b) HA nanopowders.

phase [32], which favors scaffolds resorption. By increasing the sintering temperature, the XRD patterns of the composite foams exhibited an increase in peak height and a decrease in peak width, thus indicating an increase in crystallinity and crystallite size. The crystallite sizes of the composite foams with 50 wt.% BG at different sintering temperatures were estimated in the range 24–38 nm. This is in agreement with obtained results of BET and TEM analyses. Fig. 3 shows a photograph of typical ceramic foams after sintering and also the test specimens used for measuring the compressive strength and elastic modulus of the foams. Fig. 4 shows SEM micrographs of the composite foams with different compositions sintered at 900 8C for 4 h. SEM micrographs provided information about size distribution, morphology, and interconnectivity of the pores.

Fig. 2. XRD patterns of the prepared (a) BG nanopowder, (b) foam with 100 wt.% BG sintered at 900 8C for 4 h, (c) HA nanopowder, (d) foam with 0 wt.% BG sintered at 900 8C for 4 h, and (e–g) composite foams with 50 wt.% BG sintered at 700, 800, and 900 8C for 4 h, respectively.

SEM micrographs showed that the obtained porous bioceramic structure consisted of a highly interconnected spherical porous network with the pore size between 100 and 400 mm. Furthermore, the prepared foams apart from the macropores, which provide the potential for tissue ingrowth, exhibit a large amount of micropores, which enhance bioactivity and release of ionic products. Fig. 4 also shows the sintered structure of the pore walls and pores at struts for 75 wt.% BG scaffold indicating some residual porosity was revealed in the struts of the sintered body due to the incomplete densification of the matrix, which can reduce mechanical strength and increase biodegradability of the structure. Such analysis of the bioceramic scaffold or foam is important; because it is related to osteoconductivity, resorbability, permeability, and mechanical properties. Interconnected pores are very important to allow circulation and exchange of body fluids, ion diffusion, nutritional supply, osteoblast cell penetration, and vascularization [1–3,33]. Fig. 5(a) shows the adsorption–desorption isotherm of nitrogen at 196 8C on the composite foam with 25 wt.% BG. Fig. 5(b) shows the textural pore size distribution of the foam with 25 wt.% BG, obtained by the BJH method from nitrogen adsorption isotherm. Fig. 5(b) indicates that the textural porosity was within the

Fig. 3. Photographs of typical ceramic foams after sintering and of the compressive test specimens.

H. Ghomi et al. / Materials Research Bulletin 47 (2012) 3523–3532

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Fig. 4. SEM micrographs of the composite foams with different compositions sintered at 900 8C for 4 h at different magnifications.

mesoporous range (2–50 nm). The modal mesopore diameter is about 50.93 A˚. Fig. 5(c) shows the BET plot for nitrogen adsorbed at 196 8C on the composite foam with 25 wt.% BG. The BET specific surface area and C value obtained from BET plot were found to be 40.65 m2 g1 and 107.1, respectively. In combination with the macroporous structures that provide a possibility for bone ingrowth, the high surface area of the prepared foams provides an advantage toward supplying greater interfacial contact between implant and host tissue, thus improving fixation of the implant and, consequently, reducing the chances of implant mobility. In addition, the high specific surface area of the prepared foams by increasing the bioreactions and leaching of ions from the surface during contact with a corrosive medium, such as blood, leads to stimulation of surrounding bone mineralization, which aids bone repair and fixation [12]. The mean values of compressive strength and elastic modulus of the composite foams with different compositions sintered at 900 8C were measured in the range of 0.92–2.78 MPa and 57– 219 MPa, respectively. In fact, sintering at 900 8C for 4 h provides an optimal combination of compressive strength together with the macroscopic structural features appropriate for bone ingrowth and angiogenesis [5]. The foams with compressive strength of 2.78 MPa were fabricated while maintaining interconnected pore size in excess of 100 mm. Also, in order to determine the effect of sintering temperature on the mentioned mechanical properties, the

compressive tests were performed on composite foams with 50 wt.% BG, at different sintering temperatures (7008, 8008, and 900 8C for 4 h) and the mean values of compressive strength and elastic modulus were measured in the range 0.87–1.95 MPa and 92–204 MPa, respectively. It could be observed that the compressive strength of the foams was increased with increasing the sintering temperature due to less porosity and pore size, which is in good agreement with other reports [33]. Table 1 summarizes physical and mechanical properties of nanocomposite foams as a function of foams composition and sintering temperature. Fig. 6 shows the effect of foams composition and sintering temperature on the compressive strength and elastic modulus of the composite foams. Fig. 6 shows that the compressive strength of the composite foams was increased with increasing the amount of BG to 25 wt.%, and more addition of BG leads to decrease in the mentioned mechanical properties. The compressive strength of the composite foams with 25 wt.% BG, after sintering at 900 8C for 4 h, was 2.78 MPa, which is close to the standard for a porous bioceramic bone implant (2.4 MPa) [34] and within the lower limit of the compressive strength of trabecular bone (2–12 MPa) [35]. In addition, bone ingrowth increased the compressive strength of porous implants [36]. The obtained compressive strength, in comparison to compressive strength results of the other researchers for bioceramic foams with different compositions [12–14,18], demonstrates that the foams compressive strength has increased with decreasing particle size and production of BG reinforced HA.

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Fig. 5. (a) The adsorption–desorption isotherm of nitrogen at 196 8C on the composite foam with 25 wt.% BG, (b) textural pore size distribution of the foam with 25 wt.% BG, obtained by the BJH method from nitrogen adsorption isotherm and (c) the BET plot for nitrogen adsorbed at 196 8C on the composite foam with 25 wt.% BG. Table 1 Physical and mechanical properties of the nanocomposite foams as a function of foams composition and sintering temperature. Composition

Sintering temperature

Green density (gr cm3) (S.D.)

Drying shrinkage (%) (S.D.)

Sintering shrinkage (%) (S.D.)

True porosity (%) (S.D.)

Apparent porosity (%) (S.D.)

Compressive strength (MPa) (S.D.)

Elastic modulus (MPa) (S.D.)

0 wt.% BG 25 wt.% BG 50 wt.% BG 75 wt.% BG 100 wt.% BG 50 wt.% BG 50 wt.% BG

900 8C 900 8C 900 8C 900 8C 900 8C 800 8C 700 8C

0.40 0.37 0.34 0.33 0.28 0.34 0.34

9.2 8.8 9.1 8.6 8.5 9.1 9.1

9.1 9.0 8.9 9.2 9.4 7.3 5.2

84 85 86 86 88 88 91

57 59 60 66 76 67 71

1.36 2.78 1.95 1.21 0.92 1.54 0.87

146 219 204 78 57 148 92

(0.004) (0.006) (0.007) (0.004) (0.005) (0.007) (0.007)

(0.3) (0.3) (0.4) (0.2) (0.4) (0.4) (0.4)

(0.2) (0.1) (0.3) (0.3) (0.2) (0.3) (0.4)

(1) (1) (1.5) (2) (2) (1) (1.7)

(2) (1.5) (2) (1) (2) (1.4) (1.8)

(0.09) (0.11) (0.16) (0.05) (0.09) (0.08) (0.12)

(15) (9) (6) (7) (11) (12) (5)

S.D., standard deviation.

Moreover, decreasing particle size through increasing the specific surface area facing the body fluid improves the osteoblast functions, which are proliferation, alkaline phosphatase synthesis, and calcium containing mineral deposition [19]. Fig. 7 shows a compressive stress/strain graph of the typical behavior of foams undergoing compressive testing in this study. Fig. 7 shows that the foams failed as typical porous brittle materials [33]. Region 1 of the diagram is the linear elastic region, which ends when the foam fails (maximum compressive strength). Failure is by the cracking of foam struts. Region 2 shows failure of continuous rows of pores until all struts have failed and the foam is compacted and a little collapsed. Region 3 indicates compression of the compacted piece. The in vitro studies were successful in confirming the high ability for apatite formation on the surface of the composite foams, which is a measure of the considerable bioactivity of the material [27]. Fig. 8 shows SEM micrographs of the composite foams with different compositions after 28 days of immersion in SBF. It is apparent that the bone-like apatite on the surface and inner wall of the pores are formed as the result of contacting the composite

foams with SBF. On the other hand, bone-like apatite formation leads to changing the morphology and size of the pores in the composite foams. The growth of apatite crystals into the pores leads to filling the small pores and decreasing the size of large pores and so the strength of the foams will increase too. High magnifications show that the formed bone-like apatite on the surface of the foams has tiny spherical particles with sizes in about 200–500 nm in diameter. It was confirmed that an apatite layer was formed within few days of immersion in SBF. Furthermore, by increasing the amount of BG in the composite foams, the amount of bone-like apatite formed in dependence of immersion time in SBF increased and more pores were filled. It is in good agreement with the results of other researchers that reported the thickness of the surface HA layer on BG-containing composites increased manifold in comparison to HA-containing composites after the same time of in vitro degradation in SBF [37]. It is well known that BG has a higher index of bioactivity than HA, which makes it a more appropriate material for applications in bone regeneration. Hench et al. [37] have shown that there is much more bone formed in 1 week in the presence of Bioglass than is formed

H. Ghomi et al. / Materials Research Bulletin 47 (2012) 3523–3532

Fig. 6. Effect of foams composition and sintering temperature on the compressive strength and elastic modulus of the composite foams.

when HA or other calcium phosphate ceramics are placed in the same type of defect. Furthermore, BG, being a class A bioactive material (in comparison to HA, which is class B), has shown strong bond also to soft tissues [37,38]. The application potential of the composite foams fabricated here, therefore, could include both hard and soft tissue regeneration and repair. For the foams with 25, 50, 75, and 100 wt.% BG, the functional groups of the white particles nucleated on the surface of the foams were investigated using FTIR, as shown in Fig. 9. Samples for FTIR analysis were taken from the surface of the foams to ensure detection of any formation of the surface layer and correlation between samples. This material was then ground in KBr with a sample to KBr dilution ratio of 1:100 to avoid spectral reflectance.

Fig. 7. A typical compressive stress–strain curve of the composite foam with 25 wt.% BG sintered at 900 8C for 4 h.

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The foam with 100 wt.% BG was selected because it was completely free of HA, so the formation of HA, because of being soaked in SBF, could be indicated clearly. FTIR confirmed that apatite formation took place on foams surface after being immersed in SBF and white particles were bone-like apatite. It is quite obvious that phosphate groups have been formed because of soaking in SBF for 28 days. Fig. 9 shows that bands corresponding to crystalline HCA layer formation (P–O bands at 571 and 602 cm1); appear in the FTIR spectra after 4 weeks of immersion in SBF solution. These P–O bands in the samples with 25, 50, and 75 wt.% BG due to the presence of HA, are more specific, but in the sample with 100 wt.% BG a broad band at 570–602 cm1 was present after soaking in SBF for 28 days, implying the formation of an amorphous or thin calcium phosphate layer. Bands for the phosphate groups (471 cm1: O–P–O bending n2; 571 cm1: O–P–O antisymmetric bending n4; 602 cm1: O–P–O bending n4; 961 cm1: P–O stretching n1; 1045 cm1 and 1091 cm1: P–O stretching n3) and carbonate groups (878 cm1: C–O stretching n4; 1420 cm1: C–O asymmetric stretching n3; 1472 cm1: C–O stretching n3) are consistent with the spectra for apatite compositions [39,40]. The carbonate peaks revealed that the HA contained some CO32 groups in PO43 sites of apatite lattice. This kind of apatite is more similar to biological apatite and could be more appropriate for bone replacement materials. The peaks at 731, 2981 and 3571 cm1 correspond to hydroxyl groups [40]. Two bands relating to vibration of the adsorbed water in apatite lattice could be noticed. One observed as a broad peak at 3420 cm1 and another widened band at 1630 cm1 [40]. The FTIR data in Fig. 9 shows that in addition to the absorbances for bone-like apatite (bands of phosphate and carbonate groups), strong bands due to silicate groups of BG (523 cm1: Si–O–Si bending n4; 743 and 844 cm1: Si–O symmetric stretching n3) are present for the foams with different compositions [39,41,42]. Fig. 10(a) illustrates the FTIR spectra of the foam with 100 wt.% BG which was completely free of HA, before and after soaking in SBF for 28 days. Before soaking in SBF, the sample exhibited bending and stretching vibrations of Si–O–Si bonds at 523, 743, and 844 cm1. After soaking in SBF for 28 day, bands for the phosphate groups (471, 571, 602, 961, 1045 and 1091 cm1) and carbonate groups (878 and 1420 cm1) were observed, indicating the formation of crystalline hydroxycarbonate apatite. We also confirmed the formation of apatite layer on the surface of the foams by XRD analysis. Fig. 10(b) shows the XRD pattern of the newly formed phase on the surface of the foam with 100 wt.% BG after soaking in the SBF solution for 28 days. After soaking in the SBF, the characteristic peaks for apatite [28] were observed, showing that the new material observed on the foams is crystalline bone-like apatite. Fig. 11 shows dissolution profiles and graphs of pH trends for the composite foams having different compositions as a function of soaking time. This figure also shows the change in pH with time for the SBF without any immersed foam in it, as a control. Dissolution profiles are the ionic concentration of calcium and phosphorous in SBF as a function of soaking time. These profiles illustrate the dissolution or bioresorbability behavior of synthesized composite foams with different degrees of BG. As shown in Fig. 11, the pH of the SBF without any immersed foam is approximately stable throughout the experimental period. The pH trends of the composite foams in terms of soaking time clearly showed that increasing the amount of BG increased the pH value during immersion in the SBF. The graphs of the pH trends include two regions. The Region 1 of the diagram is the linear increasing region, which ends after 5–7 days of immersion time. Region 2 shows the gradual decrease of pH, but the pH value does not reach the initial value. The pH value depends on the solubility of the foam, wherein the pH is increased as the solubility increases.

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Fig. 8. SEM micrographs of the composite foams with different compositions after 28 days of immersion in SBF at different magnifications.

There are eleven reaction stages involved in the process of complete bonding of BG to bone. The first five stages are chemical and result in the formation of a crystalline HCA layer on the surface of the BG and the other stages are the biological response [1,43]. Rapid exchange of Na+ and Ca2+ with H+ or H3O+ from solution (SBF) in the first stage increases the hydroxyl concentration and pH of the solution. Formation of Si–OH (silanols) at the glass solution interface and condensation of the Si–OH groups are then expected to occur, leaving a silica-rich layer on the surface, depleted in alkalis and alkali-earth cations (stages 2–3). The formed Si–OH groups induce apatite nucleation, and the released Ca2+ and Na+ ions accelerate apatite nucleation. Calcium and phosphate ions then migrate from the surrounding fluid to the surface through the

silica-rich layer, forming a CaO–P2O5-rich film on top of the silicarich layer (stage 4). The CaO–P2O5 film crystallizes as it incorporates OH and CO32 anions from SBF to form a mixed HCA layer and so the pH of the SBF is decreased gradually due to consuming OH ions in the formation of HCA layer [1,43]. Fig. 11 also shows the effect of foams composition on Ca and P ion concentrations at different soaking times. SBF contains approximately 80 ppm of calcium ions. The release profiles of calcium from composite foams sintered at 900 8C show that the calcium ion concentration increases by increasing the soaking time and the amount of BG. Increasing calcium ion concentration by increasing the soaking time, implies high bioactivity and resorbability of BG. It shows that the rate of BG dissolution is more than

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Fig. 9. FTIR spectra of the composite foams with different compositions after 28 days of immersion in SBF.

that of calcium ions deposited on the foams surface. Increasing calcium ion concentration by increasing the amount of BG is related to more bioactivity and resorbability of BG in comparison to HA. SBF contains 30 ppm of phosphorous ions. The release profiles of phosphorous from composite foams sintered at 900 8C show that there is a decrease in concentration of phosphorous in SBF by increasing the soaking time due to formation of the HCA layer on the foams surface and slow dissolution of phosphorus from the foams in SBF as reported by other researchers [33,40,43]. Furthermore, decreasing phosphorous by increasing the amount of BG shows the enhanced ability of formation of the HCA layer. Porous bioceramic implants or devices have been paid much attention. In fact, porous bioceramics have been used for bone defect filling, implant fixation via bone ingrowth (i.e. biological fixation), bone regeneration via tissue engineering, drug delivery, cell loading, and ocular implant [1–3]. Porous bioceramics have a high surface area that leads to excellent osteoconductivity and

Fig. 10. (a) FTIR spectra of the foam with 100 wt.% BG before and after immersion in SBF for 28 days and (b) XRD pattern of the newly formed phase on the surface of the foam with 100 wt.% BG after immersion in SBF for 28 days.

biodegradability providing fast bone ingrowth [1]. Their high surface area enhances cell adhesion to the scaffold and so promotes bone tissue regeneration [1]. It was realized that appropriate scaffold for tissue engineering applications should have appropriate

Fig. 11. Dissolution profiles and graphs of pH trends for the composite foams having different composition as a function of soaking time.

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pore size, pore interconnectivity, biocompatibility, osteoconductivity, mechanical strength, and biodegradability [1,33]. The manufactured nanocomposite foams, because of their sufficient pore size, compressive strength, and interconnectivity between pores, could be a good candidate to use in tissue engineering. Their high porosity level and nanosized structure increase the specific surface and so the bioreactions will significantly increase too. On the other hand, by changing the ratio of HA/BG, they can reach the appropriate bioactivity and biodegradability level needed for different applications. Further studies will be focused on the cell culture and in vivo tests on the prepared nanocomposite foams. 4. Conclusions This work successfully synthesized highly porous, mechanically competent, bioactive, and bioresorbable nanocomposite foams on addition of 63S BG to pure HA. The high specific surface area of the prepared foams (40.65 m2 g1), due to the high porosity level (84– 88%) and nanosized structure (24–38 nm), increases the rate of bioresorbability and accelerates the deposition process of bonelike apatite, which aids bone repair and fixation. The compressive strength measurements showed that the foams compressive strength has increased with decreasing the particle size to nano scale and production of BG reinforced HA. The resulting composite foams have similar chemical composition to the mineral phase of bone and by changing the ratio of HA/BG can reach the appropriate bioactivity and biodegradability level needed for different applications. Considering the results obtained, it seems that manufactured foams could be a good candidate for tissue engineering applications, such as drug delivery and cell loading, but cell culture and in vivo tests are needed for more assurance. Acknowledgement The authors are grateful for support of this research by Isfahan University of Technology. References [1] V.J. Shirtliff, L.L. Hench, J. Mater. Sci. 38 (2003) 4697–4707. [2] S. Oh, N. Oh, M. Appleford, J.L. Ong, Am. J. Biochem. Biotechnol. 2 (2006) 49–56.

[3] J.R. Jones, P.D. Lee, L.L. Hench, Math. Phys. Eng. Sci. A 364 (2006) 263–281. [4] J.X. Lu, B. Flautre, K. Anselme, P. Hardouin, A. Gallur, M. Descamps, B. Thierry, J. Mater. Sci. Mater. Med. 10 (1999) 111–120. [5] T.M. Freyman, I.V. Yannas, L. Gibson, Prog. Mater. Sci. 46 (2001) 273–282. [6] M.y. Darmawati, B. Oana, R.B. Aldo, J. Mater. Sci. 43 (2008) 4433–4442. [7] L.L. Hench, J.K. West, Life Chem. Rep. 13 (1996) 187–241. [8] X.L. Tang, X.F. Xiao, R.F. Liu, Mater. Lett. 59 (2005) 3841–3846. [9] M. Vallet-Regi, C.V. Ragel, A. Salinas, Eur. J. Inorg. Chem. (2003) 1029–1042. [10] P. Valerio, M.M. Pereira, A.M. Goes, M.F. Leite, Biomaterials 25 (2004) 2941–2948. [11] L.L. Hench, I.D. Xynos, J.M. Polak, J. Biomater. Sci. Polym. Ed. 15 (2004) 543–562. [12] P. Sepulveda, F.S. Ortega, M.D.M. Innocentini, V.C. Pandolfelli, J. Am. Ceram. Soc. 83 (2000) 3021–3024. [13] M. Potoczek, Mater. Lett. 62 (2008) 1055–1057. [14] Z.Y. Wu, R.G. Hill, S. Yue, D. Nightingale, P.D. Lee, J.R. Jones, Acta Biomater. 7 (2011) 1807–1816. [15] E. Adolfsson, J. Am. Ceram. Soc. 89 (2006) 1897–1902. [16] P. Sepulveda, J.G.P. Binner, S.O. Rogero, O.Z. Higa, J.C. Bressiani, J. Biomed. Mater. Res. 50 (2000) 27–34. [17] B. Chen, Z. Zhang, J. Zhang, M. Dong, D. Jiang, Mater. Sci. Eng. A 198 (2006) 435–436. [18] M. Potoczek, A. Zima, Z. Paszkiewicz, A. Slosarczyk, Ceram. Int. 35 (2009) 2249–2254. [19] T.J. Webster, C. Ergun, R.H. Doremus, R.W. Siegel, R. Bizios, Biomaterials 21 (2000) 1803–1810. [20] H. Ghomi, M.H. Fathi, H. Edris, Ceram. Int. 37 (2011) 1819–1824. [21] H. Ghomi, M.H. Fathi, H. Edris, J. Sol–Gel Sci. Technol. 58 (2011) 642–650. [22] M.H. Fathi, A. Hanifi, Mater. Lett. 61 (2007) 3978–3983. [23] V. Mortazavi, M. Mehdikhani Nahrkhalaji, M.H. Fathi, S.B. Mousavi, B. Nasr Esfahani, J. Biomed. Mater. Res. A 94 (2010) 160–168. [24] J. Will, R. Melcher, C. Treul, N. Travitzky, U. Kneser, E. Polykandriotis, R. Horch, P. Greil, J. Mater. Sci. Mater. Med. 19 (2008) 2781–2790. [25] B.D. Cullity, Elements of X-ray Diffraction, 2nd ed., Addison-Wesley, 1978. [26] D.R. Askeland, The Science and Engineering of Materials, 2nd ed., PWS Pub. Co., 1989. [27] T. Kokubo, H. Takadama, Biomaterials 27 (2006) 2907–2915. [28] XRD JCPDS data file No. 09-0432. [29] XRD JCPDS data file No. 20-0237. [30] R.C. Atwood, J.R. Jones, P.D. Lee, L.L. Hench, Scr. Mater. 51 (2004) 1029–1033. [31] S. Kapoor, U. Batra, Int. J. Chem. Biomol. Eng. 3 (2010) 24–28. [32] L.L. Hench, J. Am. Ceram. Soc. 81 (1998) 1705–1728. [33] J.R. Jones, L.M. Ehrenfried, L.L. Hench, Biomaterials 27 (2006) 964–973. [34] Implants for Surgery-Hydroxyapatite – Part 1: Ceramic Hydroxyapatite. BS ISO 13779-1:2000. [35] D.R. Carter, W.C. Hayes, Science 194 (1976) 1174–1176. [36] H. Yoshikawa, A. Myiui, J. Artif. Organs 8 (2005) 131–136. [37] J.A. Roether, A.R. Boccaccini, L.L. Hench, V. Maquet, S. Gautier, R. Jerome, Biomaterials 23 (2002) 3871–3878. [38] W. Cao, L.L. Hench, Ceram. Int. 22 (2006) 493–507. [39] C. Soundrapandian, S. Datta, B. Kundu, D. Basu, B. Sa, Am. Assoc. Pharm. Sci. 11 (2010) 1675–1683. [40] M.H. Fathi, E. Mohammadi Zahrani, J. Cryst. Growth 311 (2009) 1392–1403. [41] J. Qian, Y. Kang, Z. Wei, W. Zhang, Mater. Sci. Eng. C 29 (2009) 1361–1364. [42] M.H. Fathi, M. Kharaziha, J. Alloys Compd. 472 (2008) 540–545. [43] J.R. Jones, E. Gentleman, J. Polak, Elements 3 (2007) 393–399.