Nanocomposites of poly(l -lactide) and surface-modified chitin whiskers with improved mechanical properties and cytocompatibility

Accepted Manuscript Nanocomposites of poly(L-lactide) and surface-modified chitin whiskers with improved mechanical properties and cytocompatibility Cairong Li, Hua Liu, Binghong Luo, Wei Wen, Liumin He, Mingxian Liu, Changren Zhou PII: DOI: Reference:

S0014-3057(16)30575-4 http://dx.doi.org/10.1016/j.eurpolymj.2016.06.015 EPJ 7394

To appear in:

European Polymer Journal

Received Date: Revised Date: Accepted Date:

29 January 2016 7 June 2016 13 June 2016

Please cite this article as: Li, C., Liu, H., Luo, B., Wen, W., He, L., Liu, M., Zhou, C., Nanocomposites of poly(Llactide) and surface-modified chitin whiskers with improved mechanical properties and cytocompatibility, European Polymer Journal (2016), doi: http://dx.doi.org/10.1016/j.eurpolymj.2016.06.015

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Nanocomposites of poly(L-lactide) and surface-modified chitin whiskers with improved mechanical properties and cytocompatibility Cairong Li 1 & Hua Liu 2, Binghong Luo 2,3 *, Wei Wen 2,3, Liumin He 4, Mingxian Liu 2,3, Changren Zhou 2,3 1

Translational Medicine R&D Center, Institute of Biomedical and Health Engineering,

Shenzhen Institutes of Advanced Technology, Chinese Academy of Sciences, Shenzhen, P.R. China 2

Biomaterial Research Laboratory, Department of Material Science and Engineering,

College of Science and Engineering, Jinan University, Guangzhou 510632, P.R. China 3

Engineering Research Center of Artificial Organs and Materials, Ministry of

Education, Guangzhou 510632, P.R. China 4

Department of Biomedical Engineering, College of Life Science and Technology,

Jinan University, Guangzhou 510632, P.R. China *Corresponding authors: Tel: +86-20-85226663, Fax: +86-20-85223271, e-mail: [email protected] &

These authors contributed equally to this work and should be considered co-first

authors.

1

ABSTRACT Rod-like chitin whiskers (CHWs) were prepared by acid hydrolysis and then surface-modified with L-lactide to obtain grafted CHWs (g-CHWs). The structures, morphologies and properties of the CHWs and g-CHWs were studied by FTIR, solid state

13

C NMR, TG, XRD, FESEM and TEM. Subsequently, CHWs and g-CHWs

were used as reinforcing agents for a poly(L-lactide) (PLLA) matrix to fabricate CHW/PLLA and g-CHW/PLLA nanocomposites via solution casting. The resulting nanocomposites were fully characterized in terms of crystallization behavior, thermal transitions, morphology, mechanical properties and cytocompatibility. The nucleating effect of the CHWs and g-CHWs was confirmed by POM and DSC analysis, and as a result, the degree of crystallinity of PLLA in the CHW/PLLA and g-CHW/PLLA nanocomposites was higher than that of neat PLLA. FESEM and TEM observations indicated that g-CHWs were better dispersed throughout the matrix than CHWs due to the surface-modifications of the whiskers. Correspondingly, using an optimized g-CHW content of 5 wt%, the tensile strength, tensile modulus and fracture energy of g-CHW/PLLA nanocomposites were found to be 30.5 MPa, 1.4 GPa and 333.7 J/m2, respectively, values that were significantly superior to those of the CHW/PLLA nanocomposites and neat PLLA. Cell culture results revealed that the g-CHW/PLLA nanocomposites were more suitable for cell adhesion and differentiation than those of the neat PLLA and CHW/PLLA nanocomposites. Keywords: Poly(L-lactide); Chitin whiskers; Nanocomposite; Surface-modification; Mechanical properties; Cytocompatibility

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1. Introduction Bone fracture and damage has become a serious and worldwide health problem, and more than one million bone grafting procedures are performed in the United States yearly [1]. Generally, small-scale bone injuries can be healed through a self-repair system, whereas the reconstruction of large-volume bone defects still remains a challenging clinical problem. Autologous bone grafting and allograft bone transplantation may offer two major conventional approaches for bone repair [2-4], but both of these cures have serious limitations in clinical practice. Recently, significant attention has been devoted to developing artificial bone material substitutes to treat defects. Poly(L-lactide) (PLLA), approved for biomedical materials applications by the Food and Drug Administration (FDA), has been considered for use as a bone scaffold due to its good biocompatibility and biodegradability [5, 6]. Scaffolds for the repair of bone defects must have mechanical strength that is sufficiently similar to that of natural bone, but the mechanical properties of PLLA, particularly its toughness, are not ideal for use in hard bone tissue repair [7]. Additionally, the hydrophilicity and cytocompatibility of such polyester materials still need to be further improved. To address these issues, different types of organic and inorganic whiskers, such as cellulose and magnesium oxide [8-11] whiskers, have been incorporated into the PLLA matrix, and markedly enhanced the properties of the nanocomposites. As Wen et al. demonstrated, the mechanical strength and toughness of PLLA can be significantly improved by introducing untreated and treated MgO whiskers [11].

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As nanofillers, chitin whiskers (CHWs) have attracted particular interest for use in polymer-based nanocomposites in recent years. Chitin, the second most abundant natural polysaccharide derived from insects, crabs, shrimps and lobsters, has been recognized as a biocompatible polymer and is widely used in biomedical fields due to its biodegradability, lack of environmental hazards, low antigenicity, nontoxicity, and good absorption properties [12]. CHWs that are nano-sized in average length and diameter can be prepared by acid hydrolysis of chitin to remove amorphous domains and retain nanoscale crystallites. Because of their nanoscale size, outstanding mechanical properties (high longitudinal modulus (150 GPa) and transverse modulus (15GPa)), natural abundance and renewability, CHWs have been widely introduced into various types of polymers, such as carboxylated styrene-butadiene rubber [13], natural rubber [14, 15, 16], poly(vinyl alcohol) [17], chitosan [18] and poly(ε -caprolactone) [19], to obtain nanocomposites with improved performance. Dufresne et al. reported that nanocomposite materials were prepared using poly-( ε -caprolactone) as the matrix and CHWs as the reinforcing phase, and the mechanical properties and thermal stabilization of poly-(ε-caprolactone) were improved by the addition of chitin whiskers. Moreover, CHWs significantly reinforced either unvulcanized or prevulcanized natural rubber via the formation of a rigid three-dimensional CHW network, which has also been systematically reported by Dufresne [15]. However, as a nanofiller, CHWs tend to aggregate in a matrix, leading to a low efficiency of reinforcement in nanocomposites. Moreover, there are many hydrophilic

4

hydroxyl groups on the surfaces of CHWs, which result in incompatibility of CHWs with

hydrophobic

polymer

matrices.

To

overcome

those

drawbacks,

surface-modification by introducing certain functional groups into CHWs is a promising method. For example, modification of CHWs with hydrophobic functional groups will improve their compatibility in nonpolar solvents as well as their adhesion properties with hydrophobic matrices in composite materials [20]. According to reports in the literature, the hydrophobicity of CHWs was improved by reaction with poly(3-hydroxybutyrate-co-3-hydroxyvalerate) (PHBV) via chlorination [21] and bromohexadecane [22], respectively. Despite the plethora of papers reporting the preparation of CHW/polymer nanocomposites and surface modifications of CHWs, the poor dispersion of CHWs and weak interfacial interaction between fillers and the matrix phase are still the main challenges. Additionally, until recently, little has been reported on the preparation of CHW/PLLA nanocomposites, nor have there been many reports on the surface-modification of CHWs with L-lactide with consideration of the effects of the surface-grafted PLLA layer on the mechanical properties and cytocompatibility of CHW/PLLA nanocomposites. The major objective of this study is to use CHWs as reinforcing agents for the PLLA matrix. Moreover, this study aims to improve the dispersion of the whiskers and interfacial adhesion between the whiskers and PLLA matrix by modifying the surfaces of CHWs, as a result, to take full advantage of the whiskers and achieve optimal strengthening effects. For this purpose, CHWs were prepared and then

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surface-modified with L-lactide to obtain g-CHWs. The CHWs and g-CHWs were further introduced into the PLLA matrix to prepare CHW/PLLA and g-CHW/PLLA nanocomposite films via solution casting. The microstructures and properties of the CHWs and g-CHWs were studied. Meanwhile, the crystallization behavior, thermal properties,

whisker

dispersion,

interfacial

interaction,

mechanical

performance and cytocompatibility of the CHW/PLLA and g-CHW/PLLA nanocomposite films were evaluated in detail. 2. Materials and methods 2.1. Materials Chitin from crab was supplied by Sanland-chem International, Inc. Poly(L-lactide) (PLLA, Mw=300,000, optical purity: -156°, polydispersity index: 1.5), and L-lactide (purity: ≥99.5%, optical rotation：-261°) was purchased from Jinan Dai Gang Biological Engineering Co., Ltd. Stannous octoate (Sn(Oct) 2) (purity: approx 95%) was provided by Sigma-Aldrich and used as received. All other reagents were of analytical grade and used without further purification. 2.2. Preparation of CHWs and g-CHWs CHWs were prepared from chitin powder based on the acid hydrolysis method described by Nair et al. [14]. First, chitin powder was dispersed in a 5% KOH solution and boiled for 6 h while being stirred. The suspension was kept at room temperature overnight and then filtered and washed several times with distilled water to obtain chitin powder that was nearly free of protein. Second, the chitin powders were bleached with a 0.23 mol/L NaClO water solution containing 0.3 M sodium acetate

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buffer for 6 h at 75 C, and the bleaching solution was changed every 1 h. After being rinsed with distilled water, the sample suspension was subsequently kept in a 5 % KOH solution for 48 h to remove residual proteins. Lastly, the sample was hydrolyzed in 3 mol/L HCl at its boiling point (i.e., 102 C) for 5 h with vigorous stirring. After being diluted with distilled water and centrifuged several times, the suspension was transferred to dialysis bags and dialyzed against distilled water until the pH was close to 7. The suspension was subsequently treated by ultrasonic dispersion and freeze-dried to obtain the CHWs. g-CHWs were synthesized by a solution polymerization method through the ring-opening polymerization of L-lactide, using CHWs as macroinitiators [23]. Briefly, a mixture of CHWs and L-lactide (the molar ratios of the glucoside unit of CHWs to L-lactide were 1:2.5, 1:5 and 1:10), as well as a certain amount of Sn(Oct) 2 (the molar ratio of Sn(Oct)2 to L-lactide was 0.05%) and toluene, were added into a three-necked flask. The mixture was heated to 90 C for 24 h with stirring under a nitrogen atmosphere. After the reaction, the product was subsequently purified by Soxhlet extraction, using chloroform as a solvent to remove PLLA homopolymer and unreacted L-lactide. The residual solvent was removed under vacuum at 45 C for 24 h to obtain the g-CHWs. The synthetic route used to generate the g-CHWs (1) and the corresponding mechanism of the Sn(Oct) 2-catalyzed ring-opening polymerization of L-lactide in the presence of active hydroxyls (2) are shown in Scheme 1.

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Scheme 1. The synthetic route used to generate g-CHWs (1) and the mechanism of the Sn(Oct)2-catalyzed ring-opening polymerization of L-lactide in the presence of active hydroxyls (2). 2.3. Preparation of CHW/PLLA and g-CHW/PLLA nanocomposite films A series of CHW/PLLA and g-CHW/PLLA nanocomposite films were prepared by solution-casting [24, 25]. Briefly, a 50 g/L PLLA solution in chloroform was first prepared. Then, a certain amount of CHWs (1:5) or g-CHWs (1:5) was added to the solution with magnetic stirring and ultrasonic treatment. The resulting mixture was poured into a Petri dish and dried in a vacuum for 24 h at 45 C, and then, CHW/PLLA and g-CHW/PLLA nanocomposite films were obtained. The average thickness of the resulting nanocomposite films was 0.14 ± 0.02 mm. The compositions of the nanocomposite films are listed in Table 1 as mass fractions.

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Table 1 Mass fraction of CHWs and g-CHWs in the nanocomposite films Samples

Mass fraction of CHWs (wt.%)

Mass fraction of g-CHWs (wt.%)

Samples

CHWs1.25/PLLA

1.25

g-CHWs1.25/PLLA

1.25

CHWs2.5/PLLA

2.5

g-CHWs2.5/PLLA

2.5

CHWs3.75/PLLA

3.75

g-CHWs3.75/PLLA

3.75

CHWs5/PLLA

5

g-CHWs5/PLLA

5

CHWs7.5/PLLA

7.5

g-CHWs7.5/PLLA

7.5

CHWs10/PLLA

10

g-CHWs10/PLLA

10

2.4. Characterization 2.4.1. Characterization of CHWs and g-CHWs The infrared spectra of the CHWs and g-CHWs were recorded on a Fourier-transform infrared (FTIR) spectrometer (Bruker Equinox 55, Germany). The samples were compressed into KBr pellets for FTIR measurements. Solid state

13

C NMR was performed on a

13

C NMR spectrometer (AVANCE III

HD400, Germany Bruker). Fifty to one hundred milligrams of each sample was used, and the spectrometer was operated at a recording frequency of 100.63 MHz and spinning rate of 6 kHz. The average degree of polymerization (DP) of the oligo(L-lactide) side chains was calculated based on Eq. (1) [26]: (1) where [I(Cγ’)] and [I(C8-Cγ)-I(C8)] are the NMR signal integral intensities of the terminal 9

and internal methyl carbon of oligo(L-lactide) side chains at 20.57 and 14.77 ppm, respectively. The mass content of the oligo(L-lactide) side chains on the g-CHWs were calculated from Eq. (2) [26]: (2) where 72 and 203 g/mol in Eq. (2) are the molecular weights of the L-lactide monomer and the glucoside unit of CHWs, respectively. X-ray diffraction analysis of the CHWs and g-CHWs was conducted using a Dmax-1200 X-ray diffractometer with Cu Kα radiation over the range of 5 to 60 ° (2θ) at a scanning speed of 8 °/min. The crystallinity indexes (CrI 2θ) of the CHWs and g-CHWs were calculated according to the Segal equation (3) [27]: (3) where I2θ is the overall intensity of the peak at 2θ and Iam represents the intensity of amorphous diffraction at 16.0 º. The crystalline dimensions of different planes of nanocrystals were calculated according to the Scherrer equation (4) [27]: (4) where Bhkl is the average crystalline width of a specific plane; K is a constant (indicative of crystallite perfection and assumed to be 0.9); λ represents the wavelength of incident X-rays (λ=0.15418 nm); θ is the center of the peak; and β1/2 (in radians) represents the full width at half-maximum (FWHM) of the reflection peak. Thermogravimetric (TG) and differential thermogravimetric (DTG) analyses of the CHWs and g-CHWs were performed on a TG instrument (TG 209, Netzsch Co.,

10

Germany) under a nitrogen atmosphere. The samples were tested at a heating rate of 10 C/min from 30 to 500 C. The morphology and microstructure of the CHWs and g-CHWs were observed by field emission scanning electron microscopy (FESEM, XL30 FESEM FEG, PHILIPS) and transmission electron microscopy (TEM, a Philips CM-120 instrument, Holland). The length and width of the CHWs and g-CHWs were calculated based on the FESEM images using software to determine the particle size distribution. 2.4.2. Characterization of the CHW/PLLA and g-CHW/PLLA nanocomposite films The crystal morphology of pure PLLA and nanocomposite films was analyzed by polarized optical microscopy (POM, Axioskop40, ZEISS). The dried samples were melted at 200 C for 3 min and then cooled to 135 C at a rate of 30 C/min. The samples were observed by POM as soon as they had cooled to 135 C [28]. The melting behavior of the pure PLLA and nanocomposite films was measured by differential scanning calorimetry (DSC-204, Netzsch Co., Germany) at a heating/cooling rate of 10 C/min in the temperature range of 20 to 250 C to investigate the thermal transitions. To eliminate the thermal history of the samples, a heating-cooling-heating cycle was applied and the thermal transitions were observed based on the cooling and the second heating cycle. Before the measurement, each sample was cut into small pieces, and approximately 5-8 mg of sample was used. The degree of the crystallinity (XC) of the PLLA matrix was calculated based on Eq. (5) [29]: (5)

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where ΔHm is the melting enthalpy for PLLA-based composite films, ΔHmo is the theoretical enthalpy of 100% crystalline PLLA, w is the mass fraction of PLLA in the nanocomposites, and the melting enthalpy of a totally crystalline PLA material (ΔHmo) was considered to be 93 J/g. The tensile properties of pure PLLA and nanocomposite films were measured with a universal testing machine (SHMADZU AG-1, Japan) at a cross-head speed of 2 mm/min and with a span of 15 mm. Rectangular specimens with dimensions of 50 mm × 10 mm × (0.1-0.2) mm were cut from the films for tensile testing. Every specimen was measured five times in parallel. The average and standard deviation were calculated based on the five raw data measurements. The tensile fracture morphologies of the pure PLLA and nanocomposite films were observed by FESEM (XL30 FESEM FEG, PHILIPS). A layer of gold was used to coat the surface before observation. The dispersion of the whiskers in the nanocomposite films was observed by TEM using a Philips CM-120 instrument (Holland). The nanocomposite films were trimmed to a thickness of approximately 100 nm and used to coat the copper grids. The coated grids were then used for TEM observation. 2.4.3. In vitro cytocompatibility of the CHW/PLLA and g-CHW/PLLA nanocomposite films Mouse embryo osteoblast precursor (MC3T3-E1) cells were used for cell culture experiments in vitro. The cells were cultured in Dulbecco’s modified Eagle’s medium (DMEM) supplemented with 10 % (v/v) fetal calf serum (Gibco). Cell cultures were

12

maintained in an incubator with 5% CO2 at 37 C. After three passages, confluent cells were rinsed with phosphate-buffered saline (PBS) and detached from the culture dishes using 0.25% (w/v) EDTA-trypsin, which was neutralized with complete medium. The collected cell suspension, with a density of 1 × 104 cells/mL, was seeded onto the sample films, which had been sterilized by ultraviolet irradiation. Three replicates were used for each sample. The culture media were changed every two days. After 1, 3, 5 or 7 days of culture, the cell/material constructs were transferred into another dish, rinsed three times with PBS for 10 min each, and fixed with 2.5 % glutaraldehyde (GA) at room temperature for 4 h. The fixed cell/material constructs were then rinsed three times with PBS for 10 min each. The samples were sequentially dehydrated with ethanol and freeze-dried, and then, they were coated with gold for SEM (a Philips XL-30 scanning electron microscope, Holland) observation. A Cell Counting Kit-8 (CCK-8, Dojindo, Japan) assay was used to determine the level of cellular metabolism and thus indirectly indicate the cell proliferation status. MC3T3-E1 cells were cultured in 24-well plates for 1, 3, 5 and 7 days, and the media were changed every two days. Fifty microliters of CCK-8 (10 µL/mL) was added to each well and incubated with the cells for an additional 2 h with 5 % CO 2 at 37 C. Then, 100 µL of the solution was transferred to a 96-well plate and the absorbance of each well at 450 nm was measured by with a Thermo MULTISKAN-MK3. 3. Results and discussion 3.1. Microstructure and properties of the CHWs and g-CHWs

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The IR spectra of the CHWs and g-CHWs with different compositions are shown in Fig. 1. For the CHWs, the peaks at 3450, 3258 and 3109 cm-1 were ascribed to the O-H and N-H stretching vibrations, and the peak in the 3000~2700 cm-1 range was attributed to C-H stretching vibrations. In addition, the peaks at 1664 and 1629 cm-1 for amide I, 1564 cm-1 for amide II (a combination of the C-N-H stretching vibration and the N-H bending vibration), 1318 cm-1 for amide III and 1154 cm-1 for the C-N stretching vibration can also be observed, suggesting the existence of α-chitin [30]. The peak at 1540 cm-1, corresponding to proteins, was absent from the spectrum of the CHWs, which indicated that proteins had been removed thoroughly by the treatments [14]. In contrast to the spectrum of the CHWs, the characteristic stretching vibration peak of the ester carbonyl attributed to oligo(L-lactide) side chains at 1758 cm-1 appeared in the spectra of g-CHWs. Moreover, as the amount of L-lactide monomer gradually increased (the molar ratio of glucoside unit of CHWs to L-lactide changed from 1:2.5 to 1:10), the intensity of the absorption peak at 1758 cm-1 increased, suggesting that the amount of oligo(L-lactide) side chains on the surface of the resulting g-CHWs increased.

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Fig. 1. FTIR spectra of the CHWs and g-CHWs. The solid state

13

C NMR spectra of the CHWs and g-CHWs are shown in Fig. 2.

For the CHWs, the resonance peaks at 101.64, 80.82, 73.45, 71.04, 58.56 and 52.92 ppm were assigned to the C1, C4, C5, C3, C6 and C2 in the CHW units, respectively. The signals derived from C=O and -CH3 (C7 and C8) can be observed at 171.11 and 20.61 ppm, respectively [26]. According to reports in the literature, PLLA exhibits typical signals around 169.4-169.8 ppm, 68.8-69.2 ppm and 16.4-16.8 ppm, which are ascribed to internal C=O, -CH- and -CH3, respectively [31]. However, the peaks attributed to the internal C=O, -CH- and -CH3 (Cα, Cβ and Cγ) of the oligo(L-lactide) side chains overlapped with those of the C7, C6 and C8 ascribed to CHWs, respectively, and could not be distinguished in the spectra of g-CHWs. It is worth mentioning that three new shoulder peaks belonging to the terminal C=O, -CH- and -CH3 (Cα’, Cβ’ and Cγ’) of the oligo(L-lactide) side chains were observed at 173.34, 59.93 and 14.77 ppm in the spectra of the g-CHWs, respectively, which were similar to those reported in the literature [32]. The result confirmed the successful grafting of 15

oligo(L-lactide) chains onto the surface of the CHWs via ring-opening polymerization of L-lactide. The calculated values of DP and mass content of the oligo(L-lactide) side chains on the g-CHWs, obtained by quantifying the

13

C NMR spectra, are listed in Table 2. It

can be seen that the DP of the oligo(L-lactide) side chains on the g-CHWs was 1.54, 1.85 and 2.08, with feed molar ratios of 1:2.5, 1:5 and 1:10，respectively. It is well known that the grafting of oligo(L-lactide) side chains onto CHWs can be controlled only by the surface hydroxyl or amine groups, which should result in a low DP of the oligo(L-lactide) side chains corresponding to all of the hydroxyl groups. This result is consist with that reported in the literature [33]. In addition, when the feed molar ratio changed from 1:2.5 to 1:10, the DP values and corresponding mass content of the oligo(L-lactide) side chains showed a tendency to increase from 1.54 to 2.08 and 35.33% to 42.45%, respectively. Combining the results of FTIR and solid state of the 13

C NMR spectra of the g-CHWs, it could be inferred that g-CHWs had been

successfully synthesized. Moreover, the composition of g-CHWs can be controlled to some extent by changing the feed molar ratio, and a higher content of L-lactide would lead to a higher amount of surface-grafting of oligo(L-lactide) side chains [34].

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Fig. 2. Solid state 13C NMR spectra of the CHWs and g-CHWs. Table 2 The DP values and grafting of oligo(L-lactide) side chains on the g-CHWs calculated from the solid state 13C NMR spectra Samples

DPPLLA

Mass content (PLLA) (wt%)

g-CHWs(1:2.5)

1.54

35.33

g-CHWs(1:5)

1.85

39.62

g-CHWs(1:10)

2.08

42.45

The TG and corresponding DTG curves of CHWs and g-CHWs are shown in Fig. 3 (a, b). It can be seen that CHWs showed a two-stage degradation pattern in the TG curves. The first ranged from room temperature (approximately 32 C) to 110 C can be attributed to the evaporation of physically absorbed water. The second stage, which started from approximately 253 C and showed major degradation at 357 C, was ascribed to the thermal decomposition of the CHWs and was somewhat different from that reported in the literature [35]. In contrast to the pristine CHWs, the g-CHWs 17

showed a three-stage degradation pattern in the TG and DTG curves. Among them, the second stage, occurring from approximately 180 to 300 C, can be attributed to the thermal degradation of the surface-grafting oligo(L-lactide) side chains, which provided further evidence of the existence of grafted oligo(L-lactide) side chains on the surface of CHWs.

Fig. 3. TG (a) and DTG (b) curves of the CHWs and g-CHWs Fig. 4 shows the change in the crystalline structure of the CHWs after being surface-grafted with oligo(L-lactide) side chains, as observed by the XRD patterns, and the crystallinity indexes (CrI, %) and crystalline dimensions (B hkl) in different planes are summarized in Table 3. The characteristic diffraction peaks located at 9.6 º, 12.6 º, 19.8 º, 23.7 º and 26.3 º in 2θ corresponded to the (020), (021), (110), (130) and (013) planes, respectively, which is consistent with previously reported values for α-chitin crystal [36, 37]. The diffraction patterns of the g-CHWs basically coincided with those of CHWs in spite of the layer of oligo(L-lactide) side chains grafted to the surface of the whiskers. This suggested that the grafted oligo(L-lactide) side chains cannot intercalate into the interlayer of the CHWs to change the interlayer distance because the grafting reaction occurred only on the surface -OH and -NH2 groups of 18

CHWs. Moreover, the content of grafted oligo(L-lactide) side chains in the g-CHWs was too low to lead to any crystallization of the PLLA. As a result, there were no characteristic diffraction peaks at 17.02 º (110, 200) and 19.48 º (203) corresponding to the crystalline phase of PLLA, as can be seen in Fig. 4 [38, 39]. The related crystalline indices of g-CHWs were determined from two peaks, CrI020 and CrI110, according to the Segal method. As shown in Table 3, unmodified CHWs maintained a relative high crystallinity index of 89.59% and 93.09% in CrI020 and CrI110, respectively. Compared to the CHWs, there was no distinct difference in the crystallinity indexes of g-CHWs. When the feed molar ratio of the glucoside unit of CHWs to L-lactide was changed from 1:2.5 to 1:10, the crystallinity indices of g-CHWs decreased slightly. This indicated that minimal damage to the crystalline region occurred after the surface grafting of short PLLA chains, which is consistent with reports in the literature [35]. In addition, the crystalline dimensions of different planes (B020 and B110) of the CHWs and g-CHWs can be calculated according to the Scherrer equation. That the surface-grafting of oligo(L-lactide) side chains resulted in only a slight decrease in the crystalline sizes in the (110) and (020) planes, which demonstrated that the crystalline structure of the CHWs was preserved.

19

Fig. 4. XRD patterns of the CHWs and g-CHWs. Table 3 Crystallinity indices (CrI020, CrI110) and crystalline dimensions (B020, B110) of the CHWs and g-CHWs calculated by XRD Samples

Crystalline indexes

Crystalline dimensions

CrI020(%)

CrI110(%)

B020(nm)

B110(nm)

CHWs

89.59

93.09

8.76±0.04

6.05±0.03

g-CHWs(1:2.5)

86.41

91.19

8.63±0.03

5.85±0.02

g-CHWs(1:5)

84.86

90.97

8.70±0.04

5.99±0.03

g-CHWs(1:10)

82.60

89.11

8.52±0.06

5.59±0.03

The morphologies of the CHWs and g-CHWs were investigated by FESEM and TEM, respectively. CHWs have rod-like shapes with sharp points, as shown in Fig. 5 (a, c), which is in accordance with the morphologies reported in literature [32]. Moreover, several agglomerated CHWs can be observed, which may be mainly due to the strong hydrogen bonding interaction of the whiskers [40]. After surface-grafting 20

with PLLA, there was no significant change in the slender rod-like shape of the nano-whiskers, but the g-CHWs dispersed more uniformly in distilled water than CHWs and more single g-CHWs nano-whiskers of smaller size were obtained (Fig. 5 (b, d)). This may be explained by the fact that the hydrogen bonding interaction between the g-CHW nano-whiskers was weakened due to the hydrophobic nature of PLLA.

Fig. 5. FESEM micrographs of CHWs (a), and g-CHWs (b), and TEM micrographs of CHWs (c) and g-CHWs (d). Using FESEM imaging and grading analysis, the length and diameter distributions of the CHWs and g-CHWs (1:5) were determined and are shown in Fig. 6. The length of the CHWs ranged from 200 to 800 nm and the diameter ranged from 20 to 70 nm. The average length and width of the CHWs were estimated to be 400 and 30 nm, respectively, and the corresponding average aspect ratio (L/d, L being the length and d 21

the diameter) was therefore approximately 13.3. These dimensions are in agreement with the values for CHWs reported in literature [41-43]. The g-CHWs were 150-400 nm in length and 5-55 nm in diameter, and approximately 60% of the g-CHWs had a diameter below 25 nm. The average length and diameter of the g-CHWs was estimated to be approximately 270 and 20 nm, and the corresponding average aspect ratio (L/d) was therefore approximately 13.5, which is nearly the same as that of the CHWs. Obviously, the average length and diameter of the g-CHWs were smaller than those of the CHWs, indicating that smaller whiskers were obtained, which is in agreement with morphological observations.

Fig. 6. The length and diameter distributions of the CHWs (a, b) and g-CHWs (c, d) (1:5). 3.2. Structure and properties of the CHW/PLLA and g-CHW/PLLA nanocomposite

22

films The effects of the CHWs and g-CHWs on the crystallization rate and crystal size of the PLLA matrix were investigated by POM, and the micrographs are shown in Fig. 7. Well-developed and large spherulites were observed in pristine PLLA as the crystallization time increased. The nucleation period and spherulite size of PLLA was clearly reduced by introducing some CHWs, and more and smaller spherulites were obtained in the CHW/PLLA nanocomposite due to the nucleation effect of CHWs. For the g-CHWs, both CHWs and surface-grafted oligo(L-lactide) side chains can act as effective heterogeneous nucleating agents for the PLLA matrix. As a result, the total crystallization rate of the g-CHW/PLLA nanocomposite was much higher than that of the CHW/PLLA nanocomposite or pristine PLLA, and the former formed the smallest spherulites, which may be expected to endow the g-CHW/PLLA nanocomposite with improved toughness.

23

Fig. 7. POM micrographs of the PLLA, CHW5/PLLA and g-CHW5/PLLA nanocomposite films (bar = 10 μm) . The effects of the CHWs and g-CHWs on the crystalline structure of the PLLA matrix were further illustrated by XRD investigation, and the XRD patterns of pure PLLA and the resulting nanocomposite films with CHW and g-CHW mass contents in the range of 2.5 to 10% are shown in Fig. 8. The XRD pattern for pure PLLA had three diffraction peaks, (010), (110/200) and (203), at approximately 2θ=14.03 o, 16.33 o and 18.67 o, which corresponded to the basal spacings of 0.63, 0.54 and 0.47 nm, respectively, based on the Bragg equation. For the resulting CHW/PLLA and g-CHW/PLLA nanocomposite films, the three diffraction peaks of the PLLA matrix, as mentioned above, shifted slightly to a higher 2θ. This suggested that the crystalline structure of the PLLA matrix may be transformed from α` to α in the CHW/PLLA and g-CHW/PLLA nanocomposite films due to the heterogeneous nucleation effects of CHWs and g-CHWs [38, 44]. Additionally, when the content of the CHWs and g-CHWs increased to 10%, the XRD patterns for CHW/PLLA and g-CHW/PLLA nanocomposite films showed an increase in the intensity of the diffraction peak (203), which may be the result of the overlap between the diffraction peaks of the PLLA matrix (203) and the CHWs (110).

24

Fig. 8. XRD patterns of the CHW/PLLA and g-CHW/PLLA nanocomposite films. The DSC curves of pure PLLA and CHW/PLLA and g-CHW/PLLA nanocomposite films obtained from the second heating cycle at 10 C/min are shown in Fig. 9. The three thermal events, glass transition, polymer cold crystallization and polymer melting, can be observed on the DSC curves of all samples. The glass transition temperature (Tg), cold crystallization temperature (Tcc), melting temperature (Tm), melting enthalpy (ΔHm) and degree of crystallinity (Χc) obtained from the DSC curves are summarized in Table 4. As a typical semi-crystalline polymer, the DSC curve of neat PLLA showed Tg and Tm values of 59.5 and 171.9 C, respectively, which are similar to those reported in the literature [45-47]. With the addition of CHWs and g-CHWs into the PLLA matrix, only a slight change in the Tg and Tm of the resulting CHW/PLLA and g-CHW/PLLA nanocomposite films was observed; moreover, the Tg and Tm of the g-CHW/PLLA are slightly higher than those of the CHW/PLLA with the same whisker content. It is generally known that Tg depends on several factors, including chain flexibility, intermolecular interactions, molecular mass, branching/crosslinking and steric effects

25

[48, 49]. This result suggested that the mobility of only a small fraction of the grafted oligo(L-lactide) side chains was modified by the presence of the whiskers, which resulted in the slight decrease in the Tg. The slightly higher Tg measured for the g-CHW/PLLA compared to CHW/PLLA may be due to the molecular interaction between the PLLA matrix and the oligo(L-lactide) side chains surface-grafted to the g-CHWs, which could hinder the motion of the PLLA matrix chains. Because the Tm mainly depends on the size and perfection of crystalline lamellae, the slight change in Tm for all samples may be reasonably explained by the fact that the crystalline structures of the pure PLLA and resulting nanocomposites are nearly the same. As shown in Fig. 9, the pure PLLA and the CHW/PLLA and g-CHW/PLLA nanocomposites displayed two exothermic peaks, one at approximately 100 C and another just before the melting point of PLLA. The first peak was attributed to cold crystallization, as the crystallization of the PLLA matrix is not complete during cooling. The second peak can be attributed to a melting/recrystallization mechanism due to an increase in the thickness of the crystalline lamellae formed during cold crystallization [29]. With the addition of CHWs and g-CHWs, the cold crystallization peak of the PLLA matrix shifted slightly to lower temperatures compared to that of pure PLLA, which may be due to the heterogeneous nucleation effect of the CHWs and g-CHWs in the PLLA matrix. Correspondingly, the Χc of the PLLA matrix in the CHW/PLLA and g-CHW/PLLA nanocomposites was higher than that of pure PLLA. This result is in agreement with reports in the literature regarding the thermal properties of PLLA and PLLA/natural rubber composites [49]. Further observation

26

revealed that a higher Χc was observed for g-CHW/PLLA than for CHW/PLLA. This may be because both the whisker and surface-grafted oligo(L-lactide) side chains can act as nucleating agents for the PLLA matrix and because more and smaller spherulites were observed in the g-CHW/PLLA nanocomposite. The result is in agreement with the POM observation in this study and reports from the literature [50].

Fig. 9. DSC curves of the neat PLLA and the CHW/PLLA and g-CHW/PLLA nanocomposites from the second heating cycle. Table 4 DSC data of the neat PLLA and the CHW/PLLA (a) and g-CHW/PLLA nanocomposite films corresponding to the second heating scan Sample

Tg (C)

Tcc (C)

Tm (C)

ΔHm(J/g)

Χc (%)

PLLA

59.5

101.9

171.9

54.2

58.2

CHWs2.5/PLLA

58.6

99.5

169.9

57.1

62.9

CHWs5/PLLA

58.1

100.4

168.5

48.4

54.9

CHWs10/PLLA

57.8

98.3

167.9

49.8

59.4

g-CHWs2.5/PLLA

58.8

100.5

171.7

56.1

61.9

27

g-CHWs5/PLLA

58.5

101.3

170.7

59.6

67.5

g-CHWs10/PLLA

58.3

99.8

169.3

52.5

62.7

FESEM was used to investigate the tensile fracture morphologies of the PLLA and the CHW/PLLA and g-CHW/PLLA nanocomposites. Pure PLLA exhibited a striated but relatively smooth, fractured surface (Fig. 10 (a)), suggesting the semi-brittle characteristic of the material. There were no significant differences in the fracture morphologies between the FESEM micrographs of pure PLLA and of the CHW/PLLA and g-CHW/PLLA nanocomposites with 2.5 wt% whiskers. When the CHWs and g-CHWs contents increased to 10 wt%, a long, gully like fractured surface was observed for the CHW/PLLA nanocomposite, and a relatively smooth surface was observed for g-CHW/PLLA nanocomposites. The CHWs and g-CHWs in the PLLA matrix with 2.5 wt% and 5 wt% contents were indistinguishable from the fractured surfaces, which may be attributed to the good dispersion and their nanoscale size. As the content of g-CHWs increased to 10 wt%, the g-CHW fillers dispersed uniformly and entangled tightly with the matrix, and they are faintly visible in the FESEM micrographs.

28

Fig. 10. FESEM micrographs of the tensile fracture surfaces of the neat PLLA (a), CHW/PLLA (the mass fraction of CHWs: 2.5% (b), 5% (c), and 10% (d)) and

29

g-CHW/PLLA (the mass fraction of g-CHWs: 2.5% (e), 5% (f), and 10% (g)) nanocomposite films. To further investigate the dispersion of the whiskers in the matrix, TEM was used to observe CHW5/PLLA and g-CHW5/PLLA sheets that were approximately 100 nm in thickness, and the images are shown in Fig. 11. Compared with the observations reported in the literature, the morphology and dispersion of the whiskers before and after surface modification with the PLLA matrix can be clearly observed in high-resolution and high-contrast TEM images [13]. As shown in Fig. 11 (a and c), the strong agglomeration and inferior dispersion of the CHWs in the PLLA matrix can be observed. By contrast, there are obvious agglomerations of the g-CHWs, indicating that the g-CHWs were uniformly dispersed in the PLLA matrix on the nanometer scale (Fig. 11 (b and d)). Moreover, the interface between the CHWs and PLLA matrix is clear, indicating that there is lack of strong interfacial interaction between the two phases. However, the g-CHWs were embedded in and entangled tightly with the matrix, which caused the fuzzy interface between g-CHWs and the PLLA matrix. This phenomenon could reasonably be explained by the grafted oligo(L-lactide) side chains forming molecular chain entanglements and H-bonds at the surfaces of g-CHWs, which facilitated improved dispersion of the g-CHWs and interfacial adhesion between the g-CHWs and PLLA matrix.

30

Fig. 11. TEM images of CHW5/PLLA (a, c) and g-CHW5/PLLA (b, d) nanocomposites. The mechanical properties of the CHW/PLLA and g-CHW/PLLA nanocomposite films with different contents of whisker fillers were evaluated by tensile testing. Typical stress-strain curves of the neat PLLA and the CHW/PLLA and g-CHW/PLLA nanocomposite films are shown in Fig. 12 (a, b). All of the specimens exhibited a similar stress-strain behavior that showed a linearly elastic region at low strain initially, followed by plastic deformation at almost the same stress before fracturing. However, further observation revealed that the plastic deformation of the CHW/PLLA and g-CHW/PLLA nanocomposite films with a 3.75 wt% of CHW and g-CHW was much more obvious than that of the neat PLLA, suggesting that the toughness of the PLLA matrix was improved by introducing whisker fillers. The effects of the content and surface modification of the whiskers on the tensile strength and modulus of the PLLA matrix were investigated as shown in Fig. 12 (c 31

and d). It can be observed that the tensile strength and modulus of the resulting nanocomposites strongly depended on the whisker content. As the contents of CHWs and g-CHWs increased, the tensile strength and modulus first increased and then decreased. When the filler content reached 5 wt%, the highest tensile strength of the CHW/PLLA and g-CHW/PLLA nanocomposites was observed, with strengths of 26.7 and 30.5 MPa, or 42.8% and 63.1% higher, respectively, than that of the neat PLLA film (18.7 MPa). The highest tensile modulus of the CHW/PLLA and g-CHW/PLLA nanocomposites was also observed when the CHW and g-CHW loading content was 5 wt%, with moduli improved by 116.7% (1.3 GPa) and 133.3% (1.4 GPa), respectively, compared to that of pure PLLA (0.6 GPa). It is worth mentioning that the tensile modulus of pure PLLA and resulting nanocomposites was significantly lower than that reported in our previous research [11]. The difference in this result may be mainly attributed to such differences as the PLLA matrix parameters (e.g., molecular weight and crystallinity), processing method and thickness of the film. Additionally, this result also showed that the tensile strength and modulus of g-CHW/PLLA were significantly higher than those of CHW/PLLA nanocomposites with similar filler contents. It is generally known that the mechanical properties of nanocomposites depend on many factors, including the filler content, degree of dispersion and adhesion between the filler and matrix. According to reports in the literature, the formation of a rigid CHW network above the critical volume fraction at the percolation threshold (VRC) within the evaporated nanocomposites has been confirmed [19, 51]. For a CHW

32

network with random orientation, VRC can be calculated based on the equation VRC = 0.7/(L/d) [14]. In this study, the L/d of the CHWs and g-CHWs was approximately 13.3, which leads to a percolation threshold of 5.2 vol%, i.e., approximately 6.0 wt% (taking 1.5 [51] and 1.3 [52] for the density of the whisker fillers and PLLA matrix, respectively) for both CHWs and g-CHWs. This result basically supports the changes in the tensile strength and modulus of the CHW/PLLA and g-CHW/PLLA nanocomposites with increasing whisker content. The initial improvement in the tensile strength and modulus of the PLLA matrix upon introduction of CHWs and g-CHWs may be due to the expected reinforcing effects of the whiskers, i.e., crack deflection, crack bridging and whisker pullout. However, when the CHW or g-CHW content is greater than 5 wt%, which is close to the percolation threshold of the CHWs and g-CHWs, the excess whiskers can restrict the whisker/matrix interfacial area. Moreover, whiskers tend to aggregate in the matrix, and the interfacial phase separation of whiskers from the matrix becomes severe. As a result, the tensile strength

and

modulus

of

the

resulting

CHW/PLLA and

g-CHW/PLLA

nanocomposites decreases, which is consistent with those reported in the literatures [53, 54]. Because the oligo(L-lactide) side chains grafted to the surface of g-CHWs can mix, crystallize and entangle with the molecular chains of the PLLA matrix, the g-CHWs were dispersed uniformly throughout the matrix and tethered strongly to the matrix. Additionally, the grafted oligo(L-lactide) side chains might inhibit the formation of a percolation network of chitin whiskers. Combining these factors, it is easy to understand that the g-CHW/PLLA nanocomposite has a better tensile strength

33

and modulus than the CHW/PLLA nanocomposite. The elongation at break of the CHW/PLLA and g-CHW/PLLA nanocomposite films were determined from the stress-strain curves and are shown in Fig. 12 (e). As the whisker content increased, the elongation at the break of the nanocomposite films first increased and then decreased. When the whisker content was 5 wt%, the highest elongation at break was observed, with values of 19.6% and 13.9% for CHW/PLLA and g-CHW/PLLA, or 225.2% and 130.1% higher, respectively, than the value for PLLA alone. Interestingly, the elongation at break of g-CHW/PLLA was lower than that of CHW/PLLA with the same whisker content, which may be due to the relatively high strength and modulus of g-CHW/PLLA. The fracture energy of the CHW/PLLA and g-CHW/PLLA nanocomposite films can be calculated from the integral area of the stress-strain curve, and the result is shown in Fig. 12 (f). The fracture energy of the PLLA matrix was significantly increased by introducing a certain amount of CHWs and g-CHWs. When the CHW content was 3.75 wt%, the highest fracture energy was observed, with 251.95 J/mm measured for CHW/PLLA, which was 288.99 % higher than the fracture energy of the PLLA matrix. Because the tensile fracture toughness of materials is a reflection of the fracture energy, the higher the fracture energy, the better the fracture toughness of the material will be. Thus, this result suggested that the toughness of the PLLA matrix improved significantly due to the whiskers, which could effectively dissipate the concentrated stress at the crack tip. However, when the CHW content further increased, the interfacial phase separation and the whisker aggregation was

34

exacerbated, and as a result, the toughness of CHW/PLLA decreased significantly. For the g-CHW/PLLA nanocomposite, the highest fracture energy of 333.7 J/mm was observed when the g-CHW content was 5 wt%, which was much higher than the fracture energy of the PLLA alone and CHW/PLLA nanocomposite. The high fracture energy could correlate with energy-dissipating processes owing to crack bridging by the g-CHWs. Additionally, the strong interfacial adhesion between the g-CHWs and PLLA matrix as a result of the surface-modification of g-CHWs and uniform dispersion are other key factors.

35

Fig. 12. Stress-strain curves (a and b) and tensile properties (c, d, e, and f) of the neat PLLA and the CHW/PLLA and g-CHW/PLLA nanocomposite films. 3.3. Cell culture In order for the CHW/PLLA and g-CHW/PLLA nanocomposites to be used in the biomedical field, it is necessary to investigate their biological behaviors. In this study, the proliferation and viability of MC3T3-E1 cells on the nanocomposites were evaluated by the CCK-8 assay, and a pure PLLA film was used as the control. The OD values of cells 1, 3, 5 and 7 days after being seeded on different films were obtained 36

and are shown in Fig. 13. The results revealed that the OD values of cells cultured on all of the samples increased slowly with time in the first three days, indicating growth of MC3T3-E1 cells in the lag phase. However, 5 and 7 days after seeding, the OD values of cells on all of the films increased significantly, suggesting that the cells had entered into a logarithmic growth phase. Moreover, the OD values of MC3T3-E1 cells on the CHW/PLLA and g-CHW/PLLA nanocomposite films were significantly higher than those of cells on PLLA, which suggested that the addition of CHWs and g-CHWs can obviously improve the cytocompatibility of the PLLA material. This result was in agreement with reports in the literature on cell proliferation on the surface of the cellulose and chitin whiskers or cellulose films [55]. Moreover, the superior proliferation of MC3T3-E1 cells on the g-CHW/PLLA nanocomposite film compared to that of cells on the CHWs/PLLA may be due to the uniform dispersion of g-CHWs and good interfacial adhesion between the g-CHWs and PLLA matrix.

Fig. 13. Proliferation of MC3T3-E1 cells seeded on the neat PLLA and CHW/PLLA and g-CHW/PLLA nanocomposite films. The FESEM micrographs of MC3T3-E1 cells cultured for 1, 3, 5 and 7 days on the neat PLLA and nanocomposite films are shown in Fig. 14. It can be seen that the

37

amount and spreading area of MC3T3-E1 cells on the g-CHW/PLLA nanocomposite film were significantly superior to those of the cells on the neat PLLA and CHW/PLLA films. After 3 days of culture, cells were anchored to the surfaces of the g-CHW/PLLA nanocomposite film via filopodia with a high spreading area, whereas cells on the neat PLLA and CHW/PLLA films still had a spindle shape and showed little spreading. After 7 days of culture, the adherent cells on the g-CHW/PLLA nanocomposite film were fully stretched and had irregular or polygonal shapes with pseudopodia, and the cell number and spreading areas were still significantly higher than those for cultures on the PLLA and CHW/PLLA films. These results suggested that the g-CHW/PLLA nanocomposite film is more suitable for cell adhesion, growth and differentiation than the neat PLLA and CHW/PLLA films.

38

Fig. 14. FESEM images of the morphologies of MC3T3-E1 cells after being seeded on the neat PLLA and the CHW5/PLLA and g-CHW5/PLLA nanocomposite films for 1, 3, 5 and 7 days. 4. Conclusion In this study, rod-like CHWs with lengths of 200 to 800 nm and diameters of 20 to 70 nm were first prepared by acid hydrolysis of chitin and then successfully modified with PLLA to obtain g-CHWs via ring-opening polymerization of L-lactide. The amount of surface-grafted PLLA can be controlled to some extent by changing the feed molar ratio of the glucoside unit of CHWs to L-lactide. As the feed molar ratio changed from 1:2.5 to 1:10, the DP and mass content of grafted oligo(L-lactide) side

39

chains increased from 1.54 to 2.08 and 35.33% to 42.45%, respectively. The crystallization properties, thermal transitions and hydrophilicity of the PLLA matrix were affected by the introduction of CHWs and g-CHWs. Due to the bridging ability of the whiskers and excellent interfacial adhesion between the g-CHWs and PLLA matrix, uniform g-CHW/PLLA nanocomposites were successfully obtained and exhibited excellent tensile properties. When the g-CHW content was 5 wt%, the highest tensile strength (30.5 MPa), tensile modulus (1.4 GPa) and fracture energy (333.7 J/mm) were observed for g-CHW/PLLA nanocomposites, all of which were obviously higher than those of CHW/PLLA and the neat PLLA. This implied that the main objective of this study to illuminate the potential of g-CHWs for both mechanical reinforcement and toughening effects on a PLLA matrix was successfully achieved. The g-CHW/PLLA nanocomposite also showed better cytocompatibility than CHW/PLLA and neat PLLA, ascribed to its good biocompatibility and the more uniform dispersion of g-CHWs. All of these results suggest that the g-CHW/PLLA nanocomposite may be a promising bone tissue repair material. Acknowledgements This work was supported by the National Natural Science Foundation of China (31570981, 31571030 and 51473069), Guangdong Provincial Natural Science Foundation of China (2016A030313086), Project on the Integration of Industry, Education and Research of Guangdong Province (2013B090500107), Guangdong Provincial Natural Science Foundation of China (2016A030313086), Science and Technology Program of Guangzhou, China (No. 201510010135).

40

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Graphical abstract

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Highlights 1.

The CHWs were prepared and further surface modified with PLLA to obtain g-CHWs.

2.

The g-CHWs can disperse more uniformly in PLLA matrix than CHWs.

3.

The interfacial adhesion between g-CHWs and PLLA matrix was obviously improved.

4.

Tensile property of g-CHW/PLLA nanocomposite is superior to PLLA and CHW/PLLA.

5.

The addition of g-CHWs to PLLA can obviously improve cell adhesion and growth.

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