J Mater Sci: Mater Med (2015) 26:67 DOI 10.1007/s10856-014-5368-0

BIOMATERIALS SYNTHESIS AND CHARACTERIZATION

Synthesis and characterization of cerium- and gallium-containing borate bioactive glass scaffolds for bone tissue engineering Aylin M. Deliormanlı

Received: 7 May 2014 / Accepted: 12 September 2014 Ó Springer Science+Business Media New York 2015

Abstract Bioactive glasses are widely used in biomedical applications due to their ability to bond to bone and even to soft tissues. In this study, borate based (13-93B3) bioactive glass powders containing up to 5 wt% Ce2O3 and Ga2O3 were prepared by the melt quench technique. Cerium (Ce?3) and gallium (Ga?3) were chosen because of their low toxicity associated with bacteriostatic properties. Bioactive glass scaffolds were fabricated using the polymer foam replication method. In vitro degradation and bioactivity of the scaffolds were evaluated in SBF under static conditions. Results revealed that the cerium- and galliumcontaining borate glasses have much lower degradation rates compared to the bare borate glass 13-93B3. In spite of the increased chemical durability, substituted glasses exhibited a good in vitro bioactive response except when the Ce2O3 content was 5 wt%. Taking into account the high in vitro hydroxyapatite forming ability, borate glass scaffolds containing Ce?3 and Ga?3 therapeutic ions are promising candidates for bone tissue engineering applications.

1 Introduction Bone regeneration is a natural phenomenon by which new bone is formed after an injury. Bioactive glasses are A. M. Deliormanlı (&) Department of Materials Science and Engineering, and Center for Bone and Tissue Repair and Regeneration, Missouri University of Science and Technology, Rolla, MO, USA e-mail: [email protected]; [email protected] A. M. Deliormanlı Department of Materials Engineering, Celal Bayar University, Manisa, Turkey

promising scaffold materials for bone regeneration because of their unique ability to convert to hydroxyapatite (HA) in vivo, and their ability to bond with bone and soft tissues [1–4]. 45S5 Bioglass developed by Hench was the first material that was found to form a chemical bond with bone [5, 6]. However, it is difficult to produce porous scaffolds for bone regeneration from 45S5 since it crystallizes during sintering [6]. Borate based bioactive glasses have also been developed for potential applications in tissue engineering [7, 8]. Especially, the borate bioactive glass designated 13-93B3, formed by replacing all the SiO2 in 13-93 with B2O3, has received special interest [4, 9]. Because of their higher dissolution, some borate bioactive glasses convert faster to a calcium phosphate completely than the silicate bioactive glasses [10, 11]. Borate bioactive glasses have been shown to support tissue infiltration in vivo, as well as cell proliferation and differentiation in vitro [12]. Recently, efforts have been made to incorporate elements in the bioactive glass matrix [13]. Substitution of bioactive glasses with different trace elements such as copper, zinc, manganese, iron, magnesium and silver to modify their biological and bioactive response has been studied by several research groups [5, 14–16]. It is known that gallium (Ga?3) ions are effective against bone desorption and for the treatment of osteoporosis and cancer related hypercalcemias [17–19]. It modulates osteoclastic bone resorption without affecting osteoblasts [20] and it is the second metal ion, after platinum, to be used in cancer treatment [21]. Certain gallium compounds are used as diagnostic and therapeutic agents especially in the areas of the metabolic bone, cancer and infectious diseases [22]. It is a drug already approved for clinical use which has exhibited properties such as the ability to lower bloodcalcium levels and inhibit bacterial growth [23–25]. Gallium has been shown to be effective against the organisms

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causing tuberculosis and malaria in human beings as well as Pseudomonas aeruginosa and in the treatment of Rhodococcus equi which causes pneumonia in foals [26, 27]. Similarly, cerium oxide nanoparticles, a redox-active rare earth nanoparticle, have gained considerable interest because of their potential therapeutic applications [28]. In particular, the oxygen buffering capacity of ceria is well established [28]. Schubert et al. [29] reported that cerium oxide nanoparticles act as neuroprotective agents. Cerium oxide nanoparticles limit the amount of oxygen required to kill the cells. Zhang et al. [30] proved the positive effect of cerium on the proliferation, differentiation and mineralization of primary osteoblasts. Additionally, materials incorporated with cerium have shown promising antibacterial activities [31]. The in vitro behavior of cerium, gallium and zinc substituted sol–gel bioactive glasses and mesoporous bioactive silicate based glasses was studied by Shruti et al. [32, 33]. While the in vitro bioactivities of sol–gel derived bioactive glasses containing xCe2O3 and xGa2O3 showed moderate responses, they were found to be lower than those of xCe2O3 and xGa2O3 substituted mesoporous bioactive glasses. Similarly, the previous work of Karakoti et al. [34] have shown that nanocomposite scaffolds of bioactive glass foams containing nanoceria additives were shown to enhance the production of collagen by human mesenchymal stem cells compared to bioactive glass scaffolds without nanoceria. Although, the therapeutic action of gallium and cerium ions are well known, their influence on in vitro bioactivity and dissolution behavior of borate based bioactive glass has not been studied yet. Therefore, the aim of this study was to investigate the degradation behavior and in vitro HA forming ability of the cerium and gallium-doped 13-93B3 borate glass scaffolds.

2 Experimental work 2.1 Synthesis of bioactive glass powders The glasses used in this study were prepared by melting reagent-grade CaCO3, Na2CO3, MgCO3, K2CO3, H3BO3, CaHPO42H2O, Ga2O3 or Ce(CH3CO2)3 (Fisher Scientific, St. Louis, MO) in a Platinum crucible in air for 1 h at 1,100 °C and quenching between stainless steel plates. The composition of the glass frits synthesized in the study is given in Table 1. Particles of borate bioactive glasses were prepared by grinding the as-prepared glass frits for 3 min in a SPEX swing mill (Model 8500, Metuchen, NJ), sieving to obtain particles of size \100 lm, followed by ball milling for 5 days using ethanol as the solvent and

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zirconia balls (3 mm) as the milling media. The slurries were dried at 60 °C and the powder was sieved through a 53 lm stainless steel sieve to eliminate the agglomerates resulting from the drying step. Particle size analysis (Microtrac 3501; Microtrac Inc, USA) showed a median diameter of 13.0–13.8 lm for ball-milled bioactive glass powders. 2.2 Preparation of scaffolds Porous scaffolds were prepared using a polymer foam replication technique. A slurry containing 40 vol% glass particles was prepared by dispersing the particles in unhydrous ethanol and 4 wt% ethyl cellulose (Acros, USA) was used as dispersant. The slurry was mixed for 10 min in a planetary centrifugal mixer (Thinky AR 310). A polymer foam with a pore architecture similar to that of dry human trabecular bone was immersed in the slurry to coat the walls of the foam with the slurry. The coated foam was dried and subjected to a controlled heat treatment to decompose the foam and sinter the glass particles into a dense network. Thermal de-binding of the constructs was performed in flowing O2, using a furnace. Typically the heating rate was 0.1 °C/min in the range 100–400 °C (when the decomposition of the polymeric phase was rapid) and 0.5–1.0 °C/min outside this range. Because of the better sinterability of the borate glass, the highest thermal de-binding temperature was 500 °C. Following binder burnout, the constructs were sintered in O2 for 1 h at 575 °C, using a heating rate of 5 °C/min. At this sintering temperature scaffolds were found to produce a dense network without crystallizing the glass. 2.3 Characterizations X-ray diffraction, XRD (Philips X’Pert) was used to determine the presence of any crystalline phase formation in the as-prepared powders; XRD was performed using Cu Ka radiation at a scanning rate of 0.01°/min in the range 3–90° 2h. Fourier transform infrared spectroscopy (FTIRATR, Agilent Cary 660) was utilized to analyze the identification of the glass forming structural units in the powders. The glass transition and crystallization properties of the as-prepared glass powders were analyzed using a differential thermal analyzer (Perkin Elmer SII 7300) combined with a thermogravimetric analyzer as a function of temperature. Samples were heated to 900 °C in N2 atmosphere at 10 °C/min. The microstructure of the fabricated bioactive glass scaffolds was examined using SEM (FEI Quanta, 250 FEG) at an accelerating voltage of 5 kV and a working distance of 12 mm. The porosity of the as-prepared glass scaffolds was measured using the Archimedes method.

J Mater Sci: Mater Med (2015) 26:67 Table 1 Composition of the glasses (in wt%) used in this work

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Glass

B2O3

CaO

Na2O

K2O

MgO

P2O5

Ce2O3

Ga2O3

13-93B3

56.6

18.5

5.5

11.1

4.6

3.7

–

–

1 % Ce–B3

56.034

18.315

5.445

10.989

4.554

3.663

1

–

3 % Ce–B3

54.902

17.945

5.335

10.767

4.462

3.589

3

–

5 % Ce–B3

53.77

17.575

5.225

10.545

4.37

3.515

5

–

1 % Ga–B3

56.034

5.445

18.315

10.989

4.554

3.663

–

1

5 % Ga–B3

53.77

5,225

17.575

10.545

4.37

3.515

–

5

2.3.1 In vitro degradation and bioactivity

2.3.2 Mechanical properties

The degradation and bioactivity of the scaffolds was investigated in vitro in a simulated body fluid (SBF) under static conditions. SBF was prepared in compliance with the protocol of Kokubo et al. [35], by dissolving reagent-grade chemicals of NaCl, NaHCO3, KCl, K2HPO43H2O, MgCl26H2O, CaCl2 and Na2SO4 (Sigma Aldrich, USA) in deionized water and buffering at a pH of 7.40 with tris(hydroxymethyl)aminomethane ((CH2 OH)3CNH2) and 1 M hydrochloric acid (Fisher Scientific Inc., USA) at 37 °C. When immersed in SBF, conversion of the bioactive glass to an HA-like material was resulted in a weight loss. The weight loss of the scaffolds and variation in pH (initial pH 7.4) of SBF was measured after a certain time period to determine the rate and the extent of the conversion. A ratio of 1 g of scaffold to 500 ml of SBF was used in the conversion experiments. Scaffolds of each glass were each immersed in a polyethylene bottle containing the SBF solution, and kept for varying time periods, without shaking, in an incubator at 37 °C. Three scaffolds were used for each immersion time. After removal from the SBF, the scaffolds were dried at 60 °C and weighed. Degradation degree (weight loss, %) of the scaffolds was estimated simply as;
ðW%Þ ¼ W0  Wf =W0 ;

Vickers microhardness testing was performed to measure the surface hardness of the substituted bioactive glass samples. For this purpose, bioactive glass powders (0.5 g) were die pressed without binder using a hydraulic press under 100 MPa pressure and subsequently were sintered at 575 °C with a 5 °C/min heating rate for 1 h. Bulk densities of the pellets were measured according to the Archimedes method. Hardness measurements were conducted using a Future Tech. FM 700 Vickers microhardness tester. Prior to measurements the surface of the samples were ground with SiC grinding papers of 600 through 1,200 grit size. No etching was applied to the polished surfaces of the samples as it may have created surface irregularities affecting the indentation and the visibility of the indentation under microscope. Indentations were conducted using a Vickers diamond pyramid at 100, 300 and 500 g load for 15 s. Fifteen different measurements were carried out on each sample and the results were averaged.

where, W0 is the initial mass of the scaffold and Wf is the final mass. Additionally, after the scaffolds were removed from the bottle, the SBF solution in the bottle was cooled to room temperature, and its pH was measured using a pH meter. SEM and XRD were used to analyze the structure of the reacted scaffolds, using the conditions described previously. Fourier transform infrared spectroscopy (FTIRATR, Agilent Cary 660) was also used to characterize HAlike layer formed on the surfaces of the glass. FTIR analysis was performed in the wavenumber range of 400–4,000 cm-1.

3 Results and discussion 3.1 Bioactive glass powders The particle size distribution and an SEM image of the borate 13-93B3 bioactive glass particles used in scaffold preparation are shown Fig. 1. The particles have an angular geometry and a wide distribution of sizes, with a median diameter (d50) of 13.2 lm. The Ce2O3- and Ga2O3substituted borate glass particles showed a particle size distribution and morphology similar to those of the borate bare glass particles (results not shown). XRD analysis (Fig. 2) showed no crystalline phase formation in the Ce2O3- and Ga2O3- containing bioactive glass powders even at highest substitution levels. Figure 3 shows the FTIR spectra of the as-prepared glass powders. The region 600–800 cm-1 is assigned to bending vibrations of various borate segments. The

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Fig. 1 Particle size distribution and the SEM image of the ballmilled 13-93B3 particles

(a) 920

Absorbance (a.u.)

708

920

708

920

708

1235 5%Ce-B3 1235 3%Ce-B3 1235

928

714

1%Ce-B3 1202 13-93B3

600

800

1000 1200 1400 1600 1800

Wave number (1/cm) Fig. 2 XRD diagram of the as-prepared bioactive glass powders

(b)

123

910

709

Absorbance (a.u.)

infrared bands at 800–1,200 cm-1 are assigned to B–O stretching of tetrahedral BO4 units. The IR bands at 1,200–1,500 cm-1 are due to B–O stretching vibrations of BO3 units. The absence of band at 806 cm-1 in the spectra reveals the non-existence of boroxol rings. The band at 714 cm-1 is due to B–O–B bending vibrations of BO3 and BO4 groups. This band is shifting lower wave numbers with increasing cerium or gallium content in the samples. Pure B2O3 contains only three coordinated boron atoms and if an alkali oxide is added some of these units transform into four coordinated tetrahedral. Therefore, in 13-93B3 glass alkali oxides such as sodium, potassium and magnesium cause the formation of BO4 groups in the structure. In Fig. 3 the band at 928 cm-1 is assigned to B–O bond stretching of BO4 group vibration. A shift in this band was observed with an increase in cerium or gallium concentration. This may be due to a formation of new BO4 or BO3 units in the borate network which means that addition of cerium or gallium modifies the glass network. Hence Ce2O3 and Ga2O3 act as a network modifier in borate glass.

1230 5%Ga-B3

914

712

1230 714

928

1%Ga-B3 1202 13-93B3

600

800

1000 1200 1400 1600 1800

Wave number (cm -1) Fig. 3 FTIR spectra of the as-prepared a Ce2O3; b Ga2O3—containing bioactive glass powders

3.1.1 Thermal analysis Figure 4 shows the DTA curves of the bioactive glass powders prepared in the study. The differential thermal analysis enabled the observation of a possible formation of

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possess hierarchical porosity which is necessary for physiological functions. Ideally the template must consist of an interconnected porous structure with *80–90 % porosity. Pore sizes greater than 100 lm enable cell seeding, tissue ingrowths and vascularisation [33]. 3.2.1 In vitro degradation and bioactivity

Fig. 4 DTA curves of the as-prepared bioactive glass powders. a 1393B3; b 5 % Ce–B3 and c 5 % Ga–B3

the crystalline phase in the samples. Accordingly, incorporation of Ce?3 and Ga?3 caused an increase both in the glass transition and the crystallization temperature of the 13-93B3 bioactive glass. The endothermic and exothermic peaks shifted towards a higher temperature as Ce?3 and Ga?3 incorporated into the network composition. According to the DTA curves shown in Fig. 4 crystallization of 5 % Ga-B3 glass occurs at 720 °C whereas crystallization takes place at 666 °C for the 13-93B3 glass. The Ce?3 and Ga?3-containing glasses have a wider thermal processing window compared to the bare 13-93B3. The changes in the glass transition temperature could be explained based on the nature of the chemical bonds in the structure of the glasses. It is known that the sintering window can be widened by introducing a variety of network modifiers e.g. K2O, MgO, B2O3 and Al2O3 which increases the activation energy for crystallization [6]. It was reported that replacing 0.1 wt% of Na2O with ZnO increases the sintering window by 5 °C. Magnesium is also effective at widening the sintering window [6]. In this study, the increase observed in sintering window of the substituted samples was presumably due to the replacement of the network modifier ions such as Na?1 and Mg?2 in 13-93B3 with Ce?3 and Ga?3. 3.2 Bioactive glass scaffolds SEM images of the cross sections of the scaffolds showed that the bare 13-93B3 and Ce?3- and Ga?3- containing borate scaffolds had a similar microstructure (Fig. 5a–c). The scaffolds consisted of a dense glass network and interconnected cellular pores. The three groups of scaffolds had approximately the same porosity; within the range of 78–83 % and pores of size 100–500 lm. Natural bones

When immersed in an aqueous phosphate solution, such as the body fluid, bioactive glasses convert to an amorphous calcium phosphate or HA-like material, which is responsible for their strong bonding with surrounding tissue [2–4, 7, 12]. In this study, the bioactivity of the scaffolds was evaluated in vitro in SBF. Figure 6 shows data for the weight loss of the scaffolds containing cerium and gallium at different concentrations. According to Fig. 6a, weight loss of bare borate 13-93B3 scaffold was larger than those for the gallium-containing scaffolds. After immersion for 15 days, the weight loss reached to 49 % for the 13-93B3 glass. On the other hand, a significant decrease was observed in degradation amount of the scaffolds as a function of gallium substitution level. The weight loss of the 5 % Ga-B3 glass was only 19 % under the same conditions. Assuming the glass is converted completely to a stoichiometric HA, the theoretical weight loss is calculated to be 67 %. The lower value of the measured weight loss compared to the theoretical value may indicate that the gallium-containing borate scaffolds were not completely converted to HA after the 30-days immersion. The degradation and conversion of the borate bioactive glass scaffolds to HA in a SBF occurs by the dissolution of components such as Na2O, K2O, and B2O3 into the solution to form Na?, K?, (BO3)-3, coupled with the reaction of 4 Ca?2 ions from the glass with POfrom the solution to 3 form a HA layer on the glass [11, 36]. Gallium-containing 13-93B3 bioactive glass scaffolds showed a higher chemical durability compared to the bare borate glass scaffolds. The improvement of the durability can be correlated with the stabilizing role of Ga?3 in the glass network as a network modifier. A previous study of Franchini et al. [22] reported a similar result regarding the degradation behavior of gallium-containing phosphosilicate glasses in SBF. In this study, especially at high substitution level (5 wt%) Ga2O3 presumably modified the glass network by increasing network connectivity and reduced the degradation rate in SBF. A similar dissolution behavior was observed in cerium-containing glasses (Fig. 6b). When the experiment was stopped after 30 days, the weight loss of the 3 % Ce–B3 and 5 % Ce–B3 glasses were *42 and %31 respectively, indicating the in-complete conversion of the glass. Weight loss of the bare 13-93B3 glass scaffolds were % 58 under the same conditions.

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Fig. 5 SEM images of the sintered bioactive glass scaffolds. a 13-93B3; b 5 % Ce–B3 and c 5 % Ga–B3. Magnification 9100

(a)

(a) 70

8,6

pH

Weight loss (%)

50 40

8,4

30

8,2

20

8,0

10

7,8 0

0 0

5

10

15

20

25

10

15

20

30

35

13-93B3 1%Ce-B3 3%Ce-B3 5%Ce-B3

9,0 B3 1%Ce-B3 3%Ce-B3 5%Ce-B3

8,8

pH

50

25

(b) 9,2

70 60

Weight loss (%)

5

Time (days)

30

Time (days)

(b)

13-93B3 1%Ga-B3 5%Ga-B3

8,8

1%Ga-B3 5%Ga-B3 B3

60

9,0

8,6

40

8,4

30

8,2

20

8,0

10

7,8 0

10

20

30

Time (days)

0 0

5

10

15

20

25

30

Time (days) Fig. 6 Weight loss of the a Ga2O3; b Ce2O3—containing bioactive glass scaffolds in SBF solution as a function of immersion time

The pH of the SBF (initial value = 7.4) increased with time upon immersion of the bioactive glass scaffolds (Fig. 7). After immersion of the scaffolds for 7 days, the pH of the SBF increased to 8.11 and 8.35 for the scaffolds containing 5 % Ga2O3 and 5 % Ce2O3 respectively. Under the same conditions the pH of the SBF was 8.68 for 13-93B3 scaffolds and it

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Fig. 7 The pH of SBF solutions a Ga2O3; b Ce2O3—substituted 13-93 B3 scaffolds

continued to increase as a function of immersion time. On the other hand, a decrease was noticed on the pH of the SBF of cerium and gallium-containing borate glasses. The change in pH of the SBF was attributed to the dissolution of boron (presumably as borate ions) and the network modifiers (such as Na? and K?) during the degradation of the glass, coupled with the consumption of phosphate ions from the solution in the formation of the HA-like product.

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SEM images in Figs. 8 and 9 show the surfaces of the cerium- and gallium- containing borate scaffolds respectively after immersion in SBF for 30 days. When compared to the smooth surface of the sintered scaffolds, the surface of the reacted scaffolds had a fine particulate structure. Converted layers on the scaffolds were composed of rounded particles, which is consistent with the information reported previously for 13-93B3 glass [11, 36]. However, some differences were observed on the morphology and the particle size of the converted material depending on the cerium and gallium concentration. The surface of the scaffolds containing 1 wt% Ce2O3 was completely covered with a spherical second phase material presumably an amorphous calcium phosphate or HA. At 5 % Ce2O3 a decrease was observed in the quantity of the converted layer and the shape of the converted layer material was not spherical in all regions. The morphology of this layer was different in respect to other glasses: it was possible to observe the presence of new aggregates. EDX analysis showed that this area was essentially rich in cerium and also contains calcium and phosphorous. On the other hand, a different in vitro reactivity behavior was observed in gallium containing scaffolds. As the gallium concentration increased to 5 wt%, a calcium phosphate based material was still observed on the surface but the particle size of this material was much lower compared to the borate glass scaffolds containing 1 wt% Ga2O3. Figure 10 demonstrates the FTIR spectra of the powders obtained from the surface (converted region) of the Ce?3

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and Ga?3 substituted borate scaffolds treated in SBF (2 mg/ ml) for 20 days. Accordingly, FTIR spectra of the scaffolds showed resonances at 1,000–1,100 cm-1 and at 570 cm-1 corresponding to calcium phosphate [9, 37, 38]. The resonances at 1390 cm-1 were attributed to C–O in the (CO3)2group. A crystalline Ca–P layer, as indicated by the divided P–O bending vibration band between 500 and 600 cm-1, formed after 20 days for borate scaffolds immersed in SBF containing 1 and 3 wt% Ce2O3. A crystalline HA formation was also observed on the bare 13-93B3 scaffold after 20 days of SBF immersion. A C–O stretching vibration band appeared between 890 and 800 cm-1 indicating the formation of carbonated calcium phosphate [38, 39]. In 5 % Ce– B3 scaffolds divided P–O bending vibration band between 500 and 600 cm-1 was barely observed. Instead the converted layer might be an amorphous calcium phosphate (ACP) material, presumably a precursor to the formation of the crystalline HA phase. It is known that ACP is formed on the surface of bioactive glass at the initial stage of conversion in SBF. Ca2? deficient HA was then formed by crystallization of ACP [40]. Additionally, a decrease was observed in peak intensities as the cerium concentration increased in borate glass scaffolds. FTIR spectra of the gallium containing bioactive glass scaffolds are demonstrated in Fig. 10b. Based on these spectra it is possible to say that gallium-containing scaffolds converted to crystalline HA phase at both concentrations and substituting the glass with Ga2O3 did not degrade the HA formation ability of the borate glass scaffolds.

Fig. 8 SEM images of the surface of cerium- containing bioactive glass scaffolds immersed in SBF for 30 days. a, d 1 % Ce–B3; b, e 3 % Ce–B3 and c, f 5 % Ce–B3

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Fig. 9 SEM images of the surface of gallium-containing bioactive glass scaffolds immersed in SBF for 30 days. a, c 1 % Ga–B3; b, d 5 % Ga–B3

To eliminate the possible effects of processing conditions (such as sintering temperature etc.) on the bioactivity of the scaffolds, in vitro response was also tested on as-prepared cerium and gallium-containing borate bioactive glass powders. For this purpose, as-prepared powders were treated in SBF up to 40 days under the same conditions with the scaffolds and HA formation on the surface of the treated powders were analyzed using FTIR-ATR as described previously. Figure 11 shows the FTIR spectrum of the ceriumcontaining borate powders at different cerium concentrations and immersion times in SBF. FTIR spectra of the unconverted glasses are also shown for comparison purposes. Accordingly, A crystalline Ca–P layer (HA), as indicated by the divided P–O bending vibration band between 500 and 600 cm-1, formed after 15 days for borate scaffolds containing 1wt % cerium. The same vibration band between 500 and 600 cm-1 appeared after 20 days for 3 % Ce–B3 glass and after 30 days for 5 % Ce–B3 bioactive glass composition, respectively. Therefore, a decrease was observed in vitro bioactivity of the borate glass as the cerium concentration was increased. This is consistent with the results obtained from SEM analysis since Fig. 8 revealed a

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decrease in the amount of converted layer on the surface of 5 % Ce2O3-containing glass scaffolds. Additionally, 5 % Ce2O3-containing scaffolds after treatment in SBF for 20 days showed a similar FTIR spectrum with the related glass powders. The recent study of Goh et al. [31]. on ceriacontaining on SiO2–CaO–P2O5 bioactive glasses also showed that the low content of ceria favors the formation of stoichiometric HA. Increasing the cerium content resulted in the formation of calcium rich HA layer over 14 days based on the EDS analysis. This was attributed to the release of cerium ions which in turn competes with Ca ions for the formation of phosphates. Similarly, previous work of Shruti et al. [32] showed that substituting the glass with cerium caused a decrease in bioactivity. In this study, it was shown that the converted layer on the 5 % cerium-containing glass was amorphous calcium phosphate after 20 days of immersion in SBF and crystalline HA formation could be obtained after 30 days. Therefore, substitution of the glass with cerium at high amount decreased the reactivity and increased the chemical durability of the glasses. Results demonstrated that at both substitution level (1 and 5 wt% Ga2O3) divided P–O bending vibration band

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(a)

67

(a) 5%Ce-B3

Absorbance

Absorbance

40 d

3%Ce-B3 1%Ce-B3

30 d 20 d 15 d 7d

13-93B3

500

1000

1500

2000

un-treated

500

2500

1000

1500

2000

2500

Wavenumber (1/cm)

Wavenumber (1/cm)

(b)

(b)

40 d

Absorbance

Absorbance

5%Ga-B3

1%Ga-B3

30 d 20 d 15 d 7d un-treated

13-93B3 500

500

1000

1500

2000

2500

1000

1500

2000

2500

Wavenumber (1/cm)

Wavenumber (1/cm)

between 500 and 600 cm-1 indicating the crystalline HA formation, appeared after 15 days of immersion (Fig. 12). This is consistent with the information obtained from FTIR spectrum of the gallium-containing scaffolds treated in SBF under the same conditions with the powders. Results showed that addition of gallium to the borate glass composition increased the chemical durability of the glass in SBF and lower degradation rates were obtained compared to the bare 13-93B3 glass. FTIR analysis proofed the second phase material on the surface of the scaffolds and the SBF treated powders were crystalline HA. Therefore, it is possible to conclude that substitution of gallium in borate glass increased the chemical durability of the glass but did not cause any negative effect on the in vitro bioactivity. The previous work of Shruti et al. [32]. on 3.5 % Ga2O3 containing SiO2–CaO–P2O5 bioactive glass showed no surface layer formation until 7 days of SBF immersion. However, after 15 days the surface morphology was found

(c)

40 d

Absorbance

Fig. 10 FTIR spectra of the bioactive glass scaffolds containing a Ce2O3 and b Ga2O3 after immersion in SBF for 20 days

30 d 20 d 15 d 7d un-treated

500

1000

1500

2000

2500

Wavenumber (1/cm) Fig. 11 FTIR spectra of the cerium-containing bioactive glass powders after immersion in SBF at different time intervals. a 1 % Ce–B3; b 3 % Ce–B3; c 5 % Ce–B3

to be quite similar to the bare bioactive glass with high amount of Ca and P confirmed by the EDS results. In this study, in vitro bioactive response of the samples were not tested after shorter times of soaking such as 1 day since

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Fig. 12 FTIR spectra of the gallium-containing bioactive glass powders after immersion in SBF at different time intervals. a 1 % Ga–B3; b 5 % Ga–B3

in a previous study it was reported that borate based bioactive glass 13-93B3 has a higher degradation rate but lower HA formation ability at shorter times compared to the silicate based bioactive glasses such as 13-93 and 45S5 [42]. But in spite of the slow degradation of 5 wt% Ga2O3 -containing glass, formation of HA layer after 15 days as in the case of bare 13-93B3 glass, revealed that the bioactivity of the glass is still maintained. Lower rate of dissolution was presumably due to the participation of gallium of the glass network along with borate and its equal electric charge with (BO3)-3. These factors can make it difficult to release gallium ions from the glass network as compared to other ions such as sodium. The recent study of the Keenan et al. [23] also showed that inclusion of the gallium increased the network connectivity of the silicate based bioactive glasses; glasses containing Ga?3 maintained their amorphous nature and possess the ability to release ions when submerged in water. Similarly, previous

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work of Pickup and co-workers revealed that substitution of Na2O with Ga2O3 in Na2O–CaO–P2O5 glasses also increases the stability of the structure via the formation of GaO6 octahedra which block the migration of the Na? ions. The work of Leonelli [41] describing the in vitro bioactivity of the cerium-containing 45S5 stated that the high cerium content (13.5 wt%) improves the chemical durability of glasses so the reactivity is negatively affected. The apatite formation is prevented both by glass durability and by cerium ability to interact with phosphate giving rise to an amorphous phase. If the results of the study is compared with previously cited cerium- and gallium- containing silicate based glasses, it would be possible to say that cerium- and galliumcontaining borate glasses showed much lower degradation rates compared to the bare borate glass 13-93B3. However, bioactive response of the glass still maintained in glasses substituted with gallium. On the other hand, a decrease was observed on the HA forming ability of the cerium containing borate glasses at high cerium concentrations. It is also important to note that in this study the behavior of borate based 13-93B3 glasses were evaluated and it is known that crystalline HA formation rates of silicate and borate based glasses are different. The previous study of Liu et al. [42]. showed that 13-93B3 glass degraded almost completely and converted to a calcium phosphate material within 7–14 days in SBF. An amorphous calcium phosphate (ACP) product that formed on the 13-93B3 crystallized at a slower rate to crystalline HA when compared to the ACP that formed on the 45S5 glass. A similar degradation and bioactive response were observed in vitro for 13-93B3 glass in another work [43]. In the current study in vitro biological evaluation of the scaffolds was not examined due to a concern associated with the toxicity of boron released into the solution as borate ions, (BO3)-3 from the glass structure. Previous studies showed that in conventional ‘‘static’’ in vitro cell culture conditions, some borate glasses were toxic to cells, but the toxicity was diminished in ‘‘dynamic’’ culture conditions [4, 44]. Scaffolds of 13-93B3 borate bioactive glass were found to be toxic to murine MLO-A5 osteogeneic cells in vitro [45]. However, the same scaffolds did not show toxicity to cells in vivo and supported new tissue infiltration when implanted subcutaneously in rats [45]. Similarly, borate glass pellets implanted in rabbit tibiae produced boron concentrations in the blood far below the toxic level [4, 46]. Therefore, to examine the biological response of the cells on 13-93B3 bioactive glass scaffolds in vivo studies are preferred compared to the in vitro cell culture experiments. In our previous study we have used rat subcutaneous implantation model to determine the effect of pore size on the blood vessel formation of 13-93B3 scaffolds [47]. In this context, our next objective will be to

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study the toxicity of these new glass compositions using in vivo experiments to verify that the substitutions performed induce the desired biological properties. 3.2.2 Mechanical properties The Vickers microhardness test results of the Ce2O3 and Ga2O3 -containing borate glass pellets are shown in Fig. 13a. Accordingly, Vickers hardness values (Hv) of the substituted samples were higher compared to the bare 13-93B3 samples. As the indentation load increased, an increase was observed in hardness values of all glass samples. The variation of Hv with applied indentation load can be considered as the reflection of the reverse indentation size effect. In this type of behavior the material undergoes a relaxation which involves a release of the indentation stress along the surface away from the indentation site, which may be because of the crack formation, dislocation activity and/or elastic deformation of the indenter. For brittle materials, it is a common observation

Vickers Microhardness (GPa)

(a)

8,0 7,5 7,0

100 gr 300 gr 500 gr Regr

6,5 6,0 5,5 5,0 4,5 4,0 13-93B3

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2,56 Cerium Gallium Regr

2,54

Density (g/cm 3)

5%Ga-B3

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that cracking occurs during indenter loading half cycle. Owing to such cracking, a fraction of energy is spent in crack propagation and small indentation size results. Thus the indentation tests yield an apparently high hardness value [48]. Results also showed that the densities of the samples containing 5 wt% Ce2O3 (q = 2.553 g/cm3) and 5 wt% Ga2O3 (q = 2.535 g/cm3) were higher compared to the bare 13-93B3 (q = 2.481 g/cm3) bioactive glass (Fig. 13b). This may be attributed to the much higher densities of the Ce2O3 (6.770 g/cm3) and Ga2O3 (5.91 g/ cm3) compared to the B2O3 (2.37 g/cm3). Additionally, the observed increase in density and the Hv may be related to the formation of new BO4 groups in the borate glass network by the inclusion of Ce2O3 and Ga2O3. Since the tetrahedral BO4 groups are more strongly bonded than the triangular BO3 groups, a compact structure is expected leading to a higher density in the cerium and galliumsubstituted glasses. The recent work of Srivastava and Pyara revealed [49] an increase in Vickers microhardness of the CuO-containing 45S5 bioactive glass pellets as a function of CuO concentration. At the highest substitution level (4 % CuO) hardness value of the glass was 6.13 GPa and it was 5.75 GPa for the bare 45S5 glass at low indentation loads. A similar trend was observed when the 45S5 glass was substituted with Fe2O3. The increase of Fe2O3 in the base bioactive glass (45S5) also caused an increase in its microhardness and flexural strength [49]. The results obtained in this work, is in agreement with the previously reported studies of Srivastava et al. [49, 50] and slightly higher Hv obtained in the previous work may be related to the differences in the applied loads in the measurement or the sample preparation procedure. It is also known that silicate based glass such as 45S5 or 13-93 has higher mechanical strength compared to the 13-93B3 borate based bioactive glass [11]. This may be the possible reason for the higher Hv reported previously for the silicate based bioactive glasses.

2,52

4 Conclusions

2,50

In this study, borate based 13-93B3 bioactive glass scaffolds were made using the polymer foam replication technique. Effects of Ce2O3 and Ga2O3- substitutions on the in vitro degradation and bioactivity of borate glass scaffolds were studied. Results showed that cerium-containing scaffolds degraded at a slower rate compared to the bare borate glass scaffolds. Based on the FTIR-ATR analysis and SEM observations a decrease was obtained in vitro bioactivity of the cerium-containing borate glass scaffolds and powders at high substitution levels (at 3 and 5 %). The converted layer

2,48 2,46 0

1

2

3

4

5

6

Concentration (wt%) Fig. 13 a Vickers microhardness values (GPa); b density of the bioactive glass specimens prepared in the study

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on the 5 % cerium-containing glass was amorphous calcium phosphate after 20 days of immersion in SBF and crystalline HA formation could be obtained after 30 days. Incorporation of Ce2O3 did not alter the HA forming ability only at low concentrations. Similarly, gallium-containing borate glasses showed much lower degradation rates depending on the gallium concentration compared to the bare borate glass 13-93B3. However, a decrease was not observed in their bioactive response when they immersed in SBF. Lower rate of dissolution was attributed to the participation of gallium of the glass network along with borate. Results revealed that at high concentrations gallium and cerium ions possibly acted as a network modifier in borate glass causing an improved chemical durability. It was concluded that cerium and gallium-containing borate based bioactive glasses are suitable candidate materials for biomedical applications. Acknowledgments Support of Prof. Mohamed N.Rahaman (Missouri S&T, Rolla, USA) is greatly appreciated. The author would like to thank Xin Liu for technical assistance and Mehmet Yıldırım for performing microhardness measurements. The financial support for this research was provided by the Scientific and Technical Research Council of Turkey in the form of a TUBITAK 2219 fellowship and 1001 grant program for scientific and technological research projects; Grant No: 111M766.

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