polycaprolactone composites

Materials Science and Engineering C 46 (2015) 1–9

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Materials Science and Engineering C journal homepage: www.elsevier.com/locate/msec

In vitro bioactivity and mechanical properties of bioactive glass nanoparticles/polycaprolactone composites Lijun Ji a,⁎, Wenjun Wang a, Duo Jin b, Songtao Zhou a, Xiaoli Song a a b

College of Chemistry and Chemical Engineering, Yangzhou University, Yangzhou 225002, China College of Veterinary Medicine, Yangzhou University, Yangzhou 225002, China

a r t i c l e

i n f o

Article history: Received 8 April 2014 Received in revised form 14 August 2014 Accepted 30 September 2014 Available online 5 October 2014 Keywords: Bioactive glass nanoparticles Polycaprolactone composites Bioactivity Mechanical properties

a b s t r a c t Nanoparticles of bioactive glass (NBG) with a diameter of 50–90 nm were synthesized using the Stöber method. NBG/PCL composites with different NBG contents (0 wt.%, 10 wt.%, 20 wt.%, 30 wt.% and 40 wt.%) were prepared by a melt blending and thermal injection moulding technique, and characterized with XRD, FTIR, and SEM to study the effect of NBG on the mechanical properties and in vitro bioactivity of the NBG/PCL composites. In spite of the high addition up to 40 wt.%, the NBG could be dispersed homogeneously in the PCL matrix. The elastic modulus of the NBG/PCL composites was improved remarkably from 198 ± 13 MPa to 851 ± 43 MPa, meanwhile the tensile strength was retained in the range of 19–21.5 MPa. The hydrophilic property and degradation behavior of the NBG/PCL composites were also improved with the addition of the NBG. Moreover, the composites with high NBG content showed outstanding in vitro bioactivity after being immersed in simulated body fluid, which could be attributed to the excellent bioactivity of the synthesized NBG. © 2014 Published by Elsevier B.V.

1. Introduction As an important part of tissue engineering research, bone tissue engineering, especially in the aspect of tissue repair and regeneration has attracted many researchers. Natural hard tissues are composite materials, consisting of organic matrixes such as collagen, together with glycoprotein and ceramics similar to nanocrystalline hydroxyapatite (HA) [1]. For hard tissue applications, biodegradable polymers, such as poly(L-lactic acid) [2–4], polycaprolactone (PCL) [5–7], poly(L-lactide/ ε-caprolactone) copolymers [8], poly(lactide-co-glycolide) [9,10], poly(3-hydroxybutyrate) [11,12] and poly(3-hydroxybutyrate-co-3hydroxyvalerate) [13], have been widely used because of their favorable biocompatibility and degradability. Among these polymers, PCL, a semicrystalline linear aliphatic polyester with a high degree of crystallinity and hydrophobicity, approved by the Food and Drug Administration (FDA), has been extensively used for tissue regeneration owing to its cost-effectiveness, high toughness, excellent biocompatibility and biodegradability [1,14]. Moreover, the relatively high mechanical strength and low degradation rate of PCL provided PCL composites various advantages for application in bone tissue engineering with long term implantation period [15]. PCL nanofibers containing bioactive glass (BG) nanoparticles and simvastatin drug were prepared by electrospinning [16]. The in vitro bioactivity and drug release behavior of the PCL/BG nanofibers were studied.

⁎ Corresponding author. E-mail address: [email protected] (L. Ji).

http://dx.doi.org/10.1016/j.msec.2014.09.041 0928-4931/© 2014 Published by Elsevier B.V.

From a material viewpoint, natural bones are composites composed of biopolymers and inorganic nanocrystals. There is about 69 wt.% HA in human bones. This fact implies that polymer based composites containing bioactive inorganic phases could be promising materials for bone regenerative matrixes as these materials could provide sufficient mechanical strength accompanied with excellent osteoconductivity and bioactivity [17–19]. It has been proved that addition of silica nanoparticles into poly(vinyl alcohol) can significantly improve the mechanical properties of the poly(vinyl alcohol) film [20]. Among the bioactive inorganic materials, bioactive glasses (BGs) have attracted much attention due to their excellent ability to chemically bond with living hard tissue through the formation of a bone mineral-like HA phase on the material surface that ultimately induces direct bonding with native bone tissue, which is their so-called bioactivity [1,21,22]. Many groups prepared macroporous structures for tissue engineering by mixing BG particles and biopolymers [23–27]. These BG particles were usually prepared by mechanically grinding bulk BGs synthesized by a melt-quenching or sol–gel process, thus had a wide range of particle size in micrometer scale. These big particles formed a dispersed phase in a biopolymer matrix and could lead to nonuniform surface bioactivity and mechanical properties [28]. In most cases, the content of BG particles was not more than 20 wt.% [11,17,29,30]. Enhancement of BG content could lead to an evident decrease in mechanical properties of BG/biopolymer composites due to the big defect caused by the interface separation between micrometer scale BG particles and biopolymer matrix. Cells cultured on a BG/biopolymer composite tended to nonuniformly attach and proliferate on the area where BGs gathered because particle aggregation could promote the roughness and wettability of the composite surface,

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Fig. 1. A. A typical TEM image of NBG; B. the particle size distribution of NBG measured by DLS.

which had positive impacts on the adherence of cells [12,30,31]. However, considering that natural bones have a high content of HA crystals (about 69 wt.%), there is a great potential to improve the integrated properties of BG/biopolymer composites by enhancing its BG content. A possible solution is using nanoscale BG particles as tiny as HA crystals in natural bones. BG nanoparticles uniformly dispersed in a biopolymer matrix could be helpful for the formation of uniform properties [28]. The high surface area of nanoparticles could enhance the surface interaction between a BG and biopolymer matrix, and decrease the effect of phase separation. Herein, the authors synthesized nanoparticles of bioactive glass (NBG) and investigated the effect of NBG content on the mechanical properties and in vitro bioactivity of PCL composites. The study aims at improving the bioactivity and mechanical properties of the NBG/ PCL composite by enhancing the content of NBG in the PCL matrix. 2. Materials and method

PCL were blended at 100 °C by a HAAKE Polydrive Mixer (Thermo Fisher Scientific, China) to prepare composites with different NBG contents (10 wt.%, 20 wt.%, 30 wt.%, 40 wt.%). Finally, the NBG/PCL composites were processed into different shapes at 30–120 °C under a pressure of 700 MPa by a HAAKE MiniJet II (Thermo Fisher Scientific, China) and kept for further characterization. 2.4. Material characterization A transmission electron microscope (TEM) (TECNAI-12, Philips, Holland) and a field emission scanning electron microscope (FE-SEM) (S-4800II, Japan) were used to characterize the surface morphology and microstructure of the NBG and NBG/PCL composites. The samples were sputtered with Au before SEM observation. The size distribution of the NBG was determined by dynamic light scattering (DLS) (Mastersizer 2000, Malvern, England). The chemical composition of the NBG and NBG/ PCL composites was examined by Fourier transform infrared spectroscopy (FTIR, Tensor27, Bruke, Germany).

2.1. Materials Tetraethyl orthosilicate (TEOS), ethyl alcohol (EtOH), calcium nitrate tetrahydrate (CaNT), and ammonia hydroxide (28 wt.%), purchased from Sinopharm Chemical Reagent Co., Ltd., were all analytical grade and used directly without further purification. Poly(ε-caprolactone) (PCL, average Mn = 80,000) was purchased from Solvay (Shanghai) Co., Ltd. Deionized water was obtained from a Millipore water purification system. 2.2. Synthesis of NBG NBG with a mole composition of 75% SiO2 and 25% CaO (75S25C) were synthesized through the alkali-mediated Stöber method [32]. For a typical synthesis, water (10.95 g), EtOH (187.06 g), and ammonia hydroxide (2.88 g) were mixed together in a conical flask with magnetic stirring at room temperature for 15 min; and then CaNT (3.07 g) was added into the conical flask and dissolved under stirring. Additional amount of ammonia hydroxide was added to adjust the pH of the solution to be around 9.0. TEOS (8.08 g) was added into the previous solution and the solution was continuously stirred for 3 h. After standing the solution for about 6 h and removing the supernatant, NBG were obtained. The NBG were washed with ethanol several times and dried at 60 °C, and eventually calcined at 600 °C in air for 3 h with a heating rate of 5 °C/min.

2.5. Mechanical properties The NBG/PCL composites were prepared to be dumbbell shape samples of 75 × 12.5 × 2 mm (ISO 527-2-A5). Tensile strength, elastic modulus and elongation at break were measured by a universal mechanical testing machine (Instron 3367, USA) at a drawing speed of 50 mm/min at room temperature of 25 ± 1 °C and relative humidity of 60–65%. The average and standard deviation results of each composite were acquired by measuring five specimens. 2.6. In vitro bioactivity The in vitro bioactivity of each composite was assessed by immerging samples in standard simulated body fluid (SBF) at 37 °C. The inorganic ion concentrations of SBF were 142.0 mM Na+, 5 mM K+, 1.5 mM Mg2 +, 2.5 mM Ca2 +, 147.8 mM Cl−, 4.2 mM HCO− 3 , 1.0 mM

2.3. Preparation of NBG/PCL composites The NBG/PCL composites were produced through a melt blending and thermal injection moulding technique. The synthesized NBG and

Fig. 2. Elemental distribution analysis of the synthesized NBG particles.

L. Ji et al. / Materials Science and Engineering C 46 (2015) 1–9

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Fig. 3. A. FTIR spectra and B. XRD patterns of the synthesized NBG before and after soaking in SBF for different time (0, 1, 3, 7, 14 days). 2− HPO2− 4 , and 0.5 mM SO4 , which were similar to those of human blood plasma. The SBF was adjusted to a pH of 7.40 at 37 °C, kept at 5–10 °C in a refrigerator, and used within 30 days after preparation [33]. Planchet

samples of each composite with 20 × 2 mm in size were placed in plastic centrifuge tubes (50 mL) with SBF inside, ensuring samples being submerged. The plastic centrifuge tubes were located in a shaker with a

Fig. 4. The photographs of: A. the dumbbell shape samples for tensile testing; and B. the planchet samples for in vitro bioactivity test. The SEM images of the cross sections of: C. pure PCL and D. NBG/PCL (40 wt.%) composites after wetting-off in liquid nitrogen; E. pure PCL and F. NBG/PCL (40 wt.%) composites after tensile testing.

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Fig. 5. A) is the stress–strain curve PCL matrix; B) idem but the composites; C) and D) the effect of NGB content on the mechanical properties is showed.

setting temperature of 37 °C and a shaking speed of 170 r/min. The SBF solution was refreshed every two days during the immersion period. The samples were collected after 1, 3, 7, and 14 days of incubation. The collected samples were rinsed in distilled water five times to ensure that the surface had no SBF, and then dried using a filter paper and stored in desiccators. The formation and growth of apatite layers on the composite surface were characterized by FE-SEM, energy-dispersive X-ray spectroscopy (EDS), FTIR, and X-ray diffraction (XRD, D8 Advance, Bruke AXS, Germany).

2.7. Hydrophilicity and degradation The hydrophilicity studies were executed with an OCA20 plus videobased optical contact angle meter (Dataphysics Instruments GmbH, Filderstadt, Germany). A quantitative volume of distilled water (3 μL) was placed carefully onto the surface of a sample by a syringe, forming a drop at room temperature. After 30 s, the photo of the drop was taken. The contact angle of the sample was obtained by analyzing the recorded image of the drop. The weight loss (WL%) of the composites were

Fig. 6. The contact angles obtained under deionized water on A. a pure PCL and B. a PCL composite containing 40 wt.% NBG.

L. Ji et al. / Materials Science and Engineering C 46 (2015) 1–9 Table 1 Contact angles of the NBG/PCL composites with alterable NBG contents. NBG content [wt.%] Contact angle [°]

0 81 ± 2

10 72 ± 3

20 70 ± 5

30 64 ± 3

40 56 ± 4

calculated by measuring the weights (m1, m2) of the original sample and the after-soaking sample, respectively. Gel permeation chromatography (PL-GPC50, USA, Agilent Technologies Inc.) was used to measure molecular weight of the polymer matrix in composites after degradation in SBF for different time. Dimethyl formamide (DMF) was the solvent and used as the mobile phase at a flow rate of 1.00 mL/min and the measurement was conducted at 40 °C. Calibration of the GPC was accomplished against narrow polystyrene standards. 3. Results and discussion 3.1. The morphology and composition of the NBG As shown in Fig. 1, the NBG with good dispersibility were successfully synthesized by the alkali-mediated Stöber method. DLS confirmed that the diameter of the NBG was in the range of around 50–90 nm with an average of 70 nm (Fig. 1B). The EDS result indicated that the mean value of Si/Ca ratio of the NBG was 2.91 ± 0.14, corresponding to a mole composition of 75% SiO2 and 25% CaO (75S25C) (Fig. 2). 3.2. The in vitro bioactivity test of the NBG FTIR and XRD results confirmed the formation of carbonated hydroxyapatite (HCA) on the surface of NBG samples after soaking in SBF. The spectra in Fig. 3A show the characteristic absorption bands of Si\O\Si at 1080 cm−1 (stretch vibration), Si\O at 794 cm−1 (asymmetrical stretching vibration), and 471 cm− 1 (bending vibration). After 1 day of soaking in SBF, double bands at 567 and 604 cm−1 corresponding to P\O bending vibrations from crystalline HCA appeared and the intensity versus that of Si\O\Si increased with the extension of soaking time. The bands of the phosphate group at 958, 604, and 567 cm− 1, together with the bands of the carbonate group at 1461, 1420, and 874 cm− 1, were consistent with the reported spectra of HCA [34]. As shown in Fig. 3B, the broad XRD peak of curve 0D at 2θ = 24° confirmed the amorphous structure of the NBG. After soaking the NBG in the SBF for 1 to 14 days, new strong peaks at 2θ = 26° (002), 32–35° (211, 112, 300, 202), 39° (310), 44° (113), 46° (222), 49° (213)

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Table 2 The change of the weight average (Mw) and number average (Mn) molecular weights of the PCL matrix in the composites during degradation in SBF at 37 °C for different time. Degradation time (days)

Mw

Mn

Polydispersity index (PDI)

0 1 3 7 14

151,903 152,943 155,059 154,599 153,270

79,949 79,658 78,710 76,915 74,403

1.90 1.92 1.97 2.01 2.06

and 53° (004) corresponding to HCA crystal were observed (JCPDS 090432). These results suggested the formation of a crystalline HCA layer on the NBG after immersing in SBF, thus demonstrated the bioactivity of the NBG. 3.3. The microstructure of the NBG/PCL composites Fig. 4A and B showed the photographs of the dumbbell shape samples for tensile testing and planchet samples for in vitro bioactivity test, respectively. They had uniform color, implying good dispersion of NBG in PCL matrix. Both the pure PCL and NBG/PCL composite indicated a smooth surface after wetting-off in liquid nitrogen due to brittle fracture at low temperature (Fig. 4C and D). The good dispersion of NBG in PCL matrix was well confirmed by Fig. 4D, which was consistent with Fig. 4A and B. 3.4. Mechanical properties of the NBG/PCL composites As shown in Fig. 5A, the pure PCL dumbbell sample indicated a typical characteristic of ductile fracture with stress yield, strain softening and a high elongation at break as often seen in ductile thermoplastic polymers. The cross section of the pure PCL sample after tensile testing indicated apparent craze damage, confirming the ductile fracture in the stretching process (Fig. 4E). The addition and content increase of the NBG in the PCL matrix led to a significant decrease in the fracture strain, as shown in Fig. 5B. Only elastic deformation but no stress yield and strain softening took place in the tensile experiment for the NBG/PCL composites, suggesting a brittle fracture. The NBG/PCL (40 wt.%) sample after tensile testing showed a smooth surface similar to the surface of the NBG/PCL sample after wetting-off and no crazing was observed, confirming a brittle fracture (Fig. 4F). The possible reason could be

Fig. 7. A. The pH values of the SBF as a function of soaking time for the NBG/PCL composites, and B. the weight loss of the PCL and NBG/PCL (40 wt.%) composites after soaking in SBF for different time.

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that the NBG was not just a sparsely dispersed phase, but may become partially continuous phase in some area due to the increase of the NBG content [26]. Thus the mechanical property of the NBG/PCL composite was more like an inorganic material. Another reason could be that the poor compatibility between the polymer matrix and NBG fillers weaken the interaction among the polymer molecules and created lots of flaws, thus the composite became brittle. The relationship of NBG contents versus the tensile strength, elastic modulus and elongation at break of the NBG/PCL composites was summarized in Fig. 5C and D, respectively. A decrease in tensile strength and elongation at break resulted from the stress concentration spots due to the addition of NBG to the PCL matrix, as is often seen in inorganic substance reinforced polymer composites [35,36]. Considering that NBG was a rigid inorganic material, the NBG could rupture the PCL polymer as supporting points without deformation ability when the NBG/PCL composites were stretched and decrease the elongation at break. This could be the primary reason for the apparent decrease of elongation at break when the content of the NBG increased from 0% to 10% (Fig. 5D). The tensile strength had almost no decrease when the content of NBG increased from 10 wt.% to 20 wt.%, and had a slight decrease of less than 3 MPa from 20 wt.% to 40 wt.%, which could be attributed to the agglomeration of NBG produced by the further increase of NBG content. This result suggested that changing the NBG content from 10 wt.% to 40 wt.% had little influence on the tensile strength of the composites. Nevertheless, these NBG/PCL composites still exhibited high tensile strengths within the range of 19–21.5 MPa. There was a slight decrease of tensile strength from 19 MPa to 16 MPa after the NBG/PCL composites containing 40 wt.% NBG were soaked in SBF for 14 days, comparing with the before-soaking composite, as shown in Fig. 5C. The decrease could be attributed to the holes produced by the degradation of the NBG and the matrix, as shown in Fig. 10E–F. However, these holes were all in nanoscale benefiting from the nanoscale NBG particles, which ensured that there were no big defects emerging on the samples following the degradation of NBG. Thus the NBG/PCL composites even containing high amount of NBG could avoid significant decrease of mechanical property in the degradation process. In another aspect, the elastic modulus was improved remarkably from 198 ± 13 MPa to 851 ± 43 MPa (Fig. 5D) due to the addition of rigid NBG, confirming that NBG/PCL composite was a more rigid material than PCL. While the elastic modulus further increased to 922 ± 34 MPa after soaking the samples in SBF, mostly because of the precipitation of apatite crystals on the NBG/PCL composites. These results suggested that devices made from nanoscale NBG/PCL composite would not decrease significantly in tensile strength and become even tougher in a short term of degradation in SBF.

Fig. 8. The FTIR spectra of: (a) NBG, (b) PCL, (c) the powder scratched from the surface of the NBG/PCL composite after soaking in SBF for 7 days and (d) 14 days.

and accelerate the degradation of PCL. Thus the pH value after the third day became lower. The pH values showed a sudden decrease from day 7 to day 8. The reason could be that the exposed NBG completely degraded on day 7 and the degradation rate of PCL reached the maximum. After day 8, the PCL exposed after NBG degradation could be covered by HCA and new NBG could be exposed following the PCL degradation, thus the pH value increased from day 8 to day 10. The pH after day 10 tended to decrease very slowly and the final pH value was in the range of 7.2 to 7.4, close to the normal value of the human body fluid. As shown in Fig. 7B, the pure PCL matrix showed much weaker degradation in the in vitro environment comparing to the composites exhibiting a weight loss of 2.2% after soaking in the SBF for 14 days. It could be concluded that the weight loss of the composites was mainly due to the degradation of the NBG. The decrease of the weight loss from day 8 to day 14 could be due to the formation of apatite crystals on the NBG/PCL surface. The result of GPC indicated that the average molecular weight of the PCL in the composites had a slight decrease following the extension of the degradation time (Table 2). The polydispersity index increased a little

3.5. Hydrophilicity and degradation behavior of the NBG/PCL composites As the NBG content increased, the water contact angle of the NBG/ PCL composites decreased progressively due to the presence of the hydrophilic NBG dispersed in the composites (Fig. 6 and Table 1). Since the NBG/PCL composite with a NBG content of 40 wt.% showed similar mechanical properties to the NBG/PCL composites with low NBG contents and they were expected to possess the best bioactivity, the tests for bioactivity and degradation property were focused on the samples containing 40 wt.% NBG. The degradation process of the NBG/PCL composites could be concluded by the pH change of the SBF, the weight loss of the composites, and the change of the molecular weight of the polymer matrix PCL in the composites after soaking in SBF at 37 °C for different time. It is well known that PCL releases an acidic degradation product and NBG produce alkaline products in a SBF solution. As shown in Fig. 7A, the pH value of the SBF solution with NBG/PCL (40 wt.%) composites immersed had no change in the first day, because the alkaline degradation products from the NBG neutralized the acidic degradation product of the PCL. As confirmed by the above in vitro test, the NBG degraded quickly in SBF. The quick consumption of NBG could leave a rough PCL surface with a larger area than a smooth surface,

Fig. 9. XRD patterns of the NBG/PCL composites soaking in SBF for different time (0, 1, 3, 7, 14 days) and the pure PCL soaking in SBF for 14 days.

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after the NBG/PCL composite degraded for 14 days because the distribution of the molecular weight became broad. 3.6. The bioactivity of the NBG/PCL composites The in vitro bioactivity of the NBG/PCL composites with a NBG content of 40 wt.% was studied by soaking them in an SBF solution at 37 °C. FTIR confirmed the formation of HA on the composite surface. The absorption bands at 1723, 1192, 1468–1244, 1108, 1047, 963 and 733 cm−1 could be assigned to the characteristic absorption peaks of at 564, PCL (Fig. 8(b)). As shown in Fig. 8(c)–(d), the bands of PO3− 4 605, 958, 1034 and 1098 cm−1, and CO23 − vibrations at 1460, 1420 and 874 cm−1, suggested the existence of HCA. The double peaks at 1460 and 1420 cm−1 indicated that B-type HCA was formed because

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CO23 − replaced sectional PO34 −, and the peak at 1550 cm− 1 could be assigned to the A-type replacement [37]. In addition, the peak intensity of the HCA after 14 days was stronger than those after 7 days. The peak at 1723 cm−1 in Fig. 8(c)–(d) could be attributed to the vibration band of the carbonyl group of PCL, which was scratched from the surface of the NBG/PCL composite together with the HCA particles during the preparation of the FTIR sample. As shown in Fig. 9, the spectrum of the NBG/PCL composite (40 wt.%) before soaking in SBF indicates two main peaks at 2θ = 21° and 23°, corresponding to the peaks of pure PCL [38]. After soaking the NBG/PCL composite in SBF for 7 days, the crystalline peaks at 2θ = 25.9°, 31.8°, 40.5° and 43.8° were detected, corresponding to the crystal faces (002), (211), (221) and (113) of HA crystalline phase (JCPDS 09-0432) (Fig. 9(d)). The intensity of the peaks increased after

Fig. 10. SEM images of the surface of the NBG/PCL composites soaking in SBF for different days: A. 0 day; B. 1 day; C and D. 3 days; E and F. 7 days; G and H. 14 days.

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Fig. 11. EDS patterns of NBG/PCL (40 wt.%) composites after soaking in the SBF for (a) 1, (b) 3, (c) 7 and (d) 14 days.

soaking in SBF for 14 days (Fig. 9(e)). The XRD spectra of pure PCL before and after soaking in SBF for 14 days had almost no change. This result confirmed that the massive formation of HCA on the NBG/PCL surface should be attributed to the good bioactivity of the NBG dispersed in the PCL matrix. The apatite-forming bioactivity of the NBG/PCL composites was also estimated by observing the morphology change on the surface of the composites, as shown in Fig. 10. After soaking the NBG/PCL composites in the SBF for 1 day, no obvious change was observed on the NBG particles. After 3 days, some rod-like HCA emerged and formed a porous structure covering part of the composite surface. The rod-like HCA agglomerated to be a spherical morphology with a size of 3–5 μm in diameter, as shown in Fig. 10D. After 7 days, more rod-like HCA emerged and covered most of the composite surface. Many holes with similar size were observed on the rod-like HCA layer (Fig. 10E and F). These holes could have formed due to the degradation of the NBG embedded in the PCL matrix. The rough surface consisting of these holes could have enlarged the exposed PCL surface area, accelerated the degradation of PCL and lowered the pH value quickly between day 7 and day 8, which was confirmed by the above pH test (Fig. 6). After 14 days, most of the holes disappeared and could be covered by the rod-like HCA crystals (Fig. 10G). Some broken HCA layers appeared on the surface (Fig. 10H). The relative concentrations (in mol.%) of silicon, calcium and phosphorus were determined by EDS. With the extension of immersing time in the SBF, an increase in calcium and phosphorus peak intensity and a decrease in silicon peak intensity were observed (Fig. 11), and the Ca/P ratios were between 1.62 and 1.72, similar to the ratios reported in the bone apatite [39], confirming the formation of a HCA layer [40].

4. Conclusions The successful synthesis of BG nanoparticles with a diameter lower than 100 nm made it much easier to prepare NBG/PCL composites with high NBG contents via a melt blending and thermal injection moulding technique. The NBG were uniformly dispersed in the PCL matrix even with an addition content of 40 wt.% due to the melt blending process, which led to an obvious improvement in the elastic modulus of the composites and meanwhile the tensile strength did not show an obvious decrease. Moreover, the NBG/PCL composites presented outstanding apatite-forming bioactivity after soaking in SBF due to the high content of NBG. These findings confirmed that NBG of small size were a promising compound for synthesizing biopolymer composites containing a high content of BG. With a further decrease of NBG size, it is possible to synthesize biopolymer/NBG composites with a NBG content close to the inorganic component content in natural bones. These

NBG/PCL composites with excellent mechanical property and bioactivity could be potential candidates for bone tissue regeneration and remodeling. Acknowledgments LJ and XS were supported to work by the National Natural Science Foundation of China (No. 51273171 and No. 21201149), the Natural Science Foundation of Jiangsu Province (No. BK20131226), and a project funded by the Priority Academic Program Development of Jiangsu Higher Education Institutions. References [1] H.-s. Yun, S.-e. Kim, E.K. Park, Bioactive glass–poly(ε-caprolactone) composite scaffolds with 3 dimensionally hierarchical pore networks, Mater. Sci. Eng. C 31 (2011) 198–205. [2] J.J. Blaker, J.E. Gough, V. Maquet, I. Notingher, A.R. Boccaccini, In vitro evaluation of novel bioactive composites based on Bioglass-filled polylactide foams for bone tissue engineering scaffolds, J. Biomed. Mater. Res. A 67 (2003) 1401–1411. [3] K. Zhang, Y. Wang, M.A. Hillmyer, L.F. Francis, Processing and properties of porous poly(L-lactide)/bioactive glass composites, Biomaterials 25 (2004) 2489–2500. [4] Z. Hong, R.L. Reis, J.F. Mano, Preparation and in vitro characterization of scaffolds of poly(L-lactic acid) containing bioactive glass ceramic nanoparticles, Acta Biomater. 4 (2008) 1297–1306. [5] G. Chouzouri, M. Xanthos, In vitro bioactivity and degradation of polycaprolactone composites containing silicate fillers, Acta Biomater. 3 (2007) 745–756. [6] H.-W. Kim, J.C. Knowles, H.-E. Kim, Hydroxyapatite/poly(ε-caprolactone) composite coatings on hydroxyapatite porous bone scaffold for drug delivery, Biomaterials 25 (2004) 1279–1287. [7] D. Choi, K.G. Marra, P.N. Kumta, Chemical synthesis of hydroxyapatite/poly(εcaprolactone) composites, Mater. Res. Bull. 39 (2004) 417–432. [8] T. Jaakkola, J. Rich, T. Tirri, T. Närhi, M. Jokinen, J. Seppälä, A. Yli-Urpo, In vitro Ca–P precipitation on biodegradable thermoplastic composite of poly(ε-caprolactoneco-DL-lactide) and bioactive glass (S53P4), Biomaterials 25 (2004) 575–581. [9] M.V. Jose, V. Thomas, K.T. Johnson, D.R. Dean, E. Nyairo, Aligned PLGA/HA nanofibrous nanocomposite scaffolds for bone tissue engineering, Acta Biomater. 5 (2009) 305–315. [10] J. Schnieders, U. Gbureck, R. Thull, T. Kissel, Controlled release of gentamicin from calcium phosphate-poly(lactic acid-co-glycolic acid) composite bone cement, Biomaterials 27 (2006) 4239–4249. [11] S.K. Misra, T. Ansari, D. Mohn, S.P. Valappil, T.J. Brunner, W.J. Stark, I. Roy, J.C. Knowles, P.D. Sibbons, E.V. Jones, A.R. Boccaccini, V. Salih, Effect of nanoparticulate bioactive glass particles on bioactivity and cytocompatibility of poly(3-hydroxybutyrate) composites, J. R. Soc. Interface 7 (2010) 453–465. [12] S.K. Misra, D. Mohn, T.J. Brunner, W.J. Stark, S.E. Philip, I. Roy, V. Salih, J.C. Knowles, A.R. Boccaccini, Comparison of nanoscale and microscale bioactive glass on the properties of P(3HB)/Bioglass composites, Biomaterials 29 (2008) 1750–1761. [13] O. Bretcanu, S.K. Misra, D.M. Yunos, A.R. Boccaccini, I. Roy, T. Kowalczyk, S. Blonski, T.A. Kowalewski, Electrospun nanofibrous biodegradable polyester coatings on Bioglass®-based glass–ceramics for tissue engineering, Mater. Chem. Phys. 118 (2009) 420–426. [14] V.R. Sinha, K. Bansal, R. Kaushik, R. Kumria, A. Trehan, Poly-epsilon-caprolactone microspheres and nanospheres: an overview, Int. J. Pharm. 278 (2004) 1–23. [15] H. Sun, L. Mei, C. Song, X. Cui, P. Wang, The in vivo degradation, absorption and excretion of PCL-based implant, Biomaterials 27 (2006) 1735–1740.

L. Ji et al. / Materials Science and Engineering C 46 (2015) 1–9 [16] M. Kouhi, M. Morshed, J. Varshosaz, M.H. Fathi, Poly(e-caprolactone) incorporated bioactive glass nanoparticles and simvastatin nanocomposite nanofibers: preparation, characterization and in vitro drug release for bone regeneration applications, Chem. Eng. J. 228 (2013) 1057–1065. [17] E. Tamjid, R. Bagheri, M. Vossoughi, A. Simchi, Effect of particle size on the in vitro bioactivity, hydrophilicity and mechanical properties of bioactive glass-reinforced polycaprolactone composites, Mater. Sci. Eng. C 31 (2011) 1526–1533. [18] K. Rezwan, Q.Z. Chen, J.J. Blaker, A.R. Boccaccini, Biodegradable and bioactive porous polymer/inorganic composite scaffolds for bone tissue engineering, Biomaterials 27 (2006) 3413–3431. [19] A.A. Rocha de Oliveira, S.M. de Carvalho, M.F. Leite, Development of biodegradable polyurethane and bioactive glass nanoparticles scaffolds for bone tissue engineering applications, J. Biomed. Mater. Res. 100B (2012) 1387–1396. [20] M. Bhattacharya, S. Chaudhry, High-performance silica nanoparticle reinforced poly(vinyl alcohol) as templates for bioactive nanocomposites, Mater. Sci. Eng. C 33 (2013) 2601–2610. [21] T. Kokubo, H. Kushitani, S. Sakka, T. Kitsugi, T. Yamamuro, Solutions able to reproduce in vivo surface-structure changes in bioactive glass, J. Biomed. Mater. Res. 24 (1990) 721–734. [22] L.L. Hench, The story of Bioglass, J. Mater. Sci. Mater. Med. 17 (2006) 967–978. [23] P. Fabbri, V. Cannillo, A. Sola, A. Dorigato, F. Chiellini, Highly porous polycaprolactone45S5 Bioglass® scaffolds for bone tissue engineering, Compos. Sci. Technol. 70 (2010) 1869–1878. [24] Y. Minaberry, M.A. Jobbágy, Macroporous Bioglass scaffolds prepared by coupling sol–gel with freeze drying, Chem. Mater. 23 (2011) 2327–2332. [25] A. Garcia, I. Izquierdo-Barba, M. Colilla, C.L. de Laorden, M. Vallet-Regi, Preparation of 3D scaffolds in the SiO2–P2O5 system with tailored hierarchical meso-macroporosity, Acta Biomater. 7 (2011) 1265–1273. [26] L. Ji, Y. Si, A. Li, W. Wang, D. Qiu, A. Zhu, Progress of three-dimensional macroporous bioactive glass for bone regeneration, Front. Chem. Sci. Eng. 6 (2012) 470–483. [27] M. Yazdimamaghani, D. Vashaee, S. Assefa, Hybrid macroporous Gelatin/BioactiveGlass/Nanosilver scaffolds with controlled degradation behavior and antimicrobial activity for bone tissue engineering, J. Biomed. Nanotechnol. 10 (2014) 911–931.

9

[28] E.M. Valliant, J.R. Jones, Softening bioactive glass for bone regeneration: sol–gel hybrid materials, Soft Matter 7 (2011) 5083. [29] A. Larrañaga, J.-R. Sarasua, Effect of bioactive glass particles on the thermal degradation behaviour of medical polyesters, Polym. Degrad. Stab. 98 (2013) 751–758. [30] A. Liu, Z. Hong, X. Zhuang, X. Chen, Y. Cui, Y. Liu, X. Jing, Surface modification of bioactive glass nanoparticles and the mechanical and biological properties of poly(L-lactide) composites, Acta Biomater. 4 (2008) 1005–1015. [31] D.D. Deligianni, N. Katsala, S. Ladas, D. Sotiropoulou, J. Amedee, Y.F. Missirlis, Effect of surface roughness of the titanium alloy Ti–6Al–4V on human bone marrow cell response and on protein adsorption, Biomaterials 22 (2001) 1241–1251. [32] W. Stöber, A. Fink, Controlled growth of monodisperse silica spheres in the micron size range, J. Colloid Interface Sci. 26 (1968) 62–69. [33] T. Kokubo, H. Takadama, How useful is SBF in predicting in vivo bone bioactivity, Biomaterials 27 (2006) 2907–2915. [34] I. Notingher, J.R. Jones, S. Verrier, I. Bisson, P. Embanga, P. Edwards, J.M. Polak, L.L. Hench, Application of FTIR and Raman spectroscopy to characterisation of bioactive materials and living cells, Spectroscopy 17 (2003) 275–288. [35] W.-Y. Choi, H.-E. Kim, S.-Y. Oh, Y.-H. Koh, Synthesis of poly(ε-caprolactone)/ hydroxyapatite nanocomposites using in-situ co-precipitation, Mater. Sci. Eng. C 30 (2010) 777–780. [36] H. Zou, S. Wu, J. Shen, Polymer/silica nanocomposites preparation, characterization, properties, and applications, Chem. Rev. 108 (2008) 3893–3957. [37] M. Vignoles, G. Bonel, D.W. Holcomb, R.A. Young, Influence of preparation conditions on the composition of type B carbonated hydroxyapatite and on the localization of the carbonate ions, Calcif. Tissue Int. 43 (1988) 33–40. [38] C.C. Chen, J.Y. Chueh, H. Tseng, H.M. Huang, S.Y. Lee, Preparation and characterization of biodegradable PLA polymeric blends, Biomaterials 24 (2003) 1167–1173. [39] A. Ravaglioli, A. Krajewski, G.C. Celoti, A. Piancastelli, B. Bacchini, L. Montanari, G. Zama, L. Piombi, Mineral evolution of bone, Biomaterials 17 (1996) 617–622. [40] M. Vallet-Regí, A.M. Romero, C.V. Ragel, R.Z. LeGeros, XRD, SEM–EDS, and FTIR studies of in vitro growth of an apatite-like layer on sol–gel glasses, J. Biomed. Mater. Res. 44 (1999) 416–421.