Molecular Interactions and Toughening Mechanisms in Silk Fibroin

Article Cite This: Biomacromolecules 2019, 20, 2295−2304

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Molecular Interactions and Toughening Mechanisms in Silk Fibroin− Epoxy Resin Blend Films Kang Yang,†,‡ Kenjiro Yazawa,‡ Kousuke Tsuchiya,‡ Keiji Numata,*,‡ and Juan Guan*,†,§ †

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International Research Center for Advanced Structural and Biomaterials, School of Materials Science and Engineering, Beihang University, Beijing 100191, China ‡ Biomacromolecules Research Team, RIKEN Center for Sustainable Resource Science, 2-1 Hirosawa, Wako-shi, Saitama 351-0198, Japan § Beijing Advanced Innovation Center for Biomedical Engineering, Beihang University, Beijing 100083, China S Supporting Information *

ABSTRACT: Natural silkworm silks have been applied to reinforce epoxy resin to achieve sub-ambient and impact toughness in the composite. However, the molecular interactions at the silk fiber−matrix interface of the composite are poorly understood. In this work, silk fibroin extracted from Bombyx mori silk is blended with an epoxy resin polymer system to study the molecular interactions between silk fibroin, epoxy compounds, and hardeners. The effects of chemical crosslinks between epoxy groups and hardeners or silk fibroin, as well as physical crosslinks in the β-sheet structure of silk fibroin, were discussed on the thermal stability, glass transition behavior, and mechanical properties of the blend films. A relationship between the crosslinking structure and mechanical properties for the films is proposed to enlighten on the toughening mechanisms. The findings would provide insights into forming strong and tough silk fibroin material as well as understanding the interface interactions of silk−epoxy composites. Silk fibroin or SF regenerated from silk fiber has been used as a popular biomaterial for a wide range of biomedical applications, owing to its “green” origin, biocompatibility, and biodegradability. Similar to silk fiber composite, the SF-based blend materials have a potential for biomedical applications such as bone fixture devices or bone tissue engineering materials. SF comprises of over 20 amino acids in varied sequence and composition depending on the producing organism. Bombyx mori or B. mori SF contains about 5.3% tyrosine and 0.2% lysine amino acid residues, whose active phenolic hydroxyl group and amino group could react with epoxy groups. B. mori SF can also form a characteristic β-sheet structure,16−18 a physical crosslinking mechanism that endows regenerated silk materials with versatile and tunable properties.19,20 Nevertheless, the reactivity of SF with epoxy/epoxy resin polymers and the molecular interactions between SF and epoxy resin remain to be studied. In addition, it was noticed that many earlier works on reconstituted silk fibroin (RSF) products used aqueous solutions.21,22 The aqueous solution cast RSF films turned out to be mechanically weak and brittle due to prevalent “entanglements” in the silk I structure and poorly organized silk II structure.8,23−25 In contrast, using hexafluoroisopropanol (HFIP) solvent to dissolve freeze-dried silk fibroin for film

1. INTRODUCTION Epoxy resin polymers are widely used as adhesives, coatings, and matrix materials in high-performance polymer composites for automotive, marine, and aerospace applications.1−3 High modulus and strength of most epoxy resin thermosets are caused by high-density chemical crosslinks between oxidative epoxy groups and active hydrogen-containing species (i.e., double amino or hydroxyl groups) of the hardeners. However, one disadvantage of such highly cross-linked epoxy resins might be the poor ductility and toughness.4 In addition, the epoxy resin thermosets with high modulus and strength are popularly applied as a matrix for carbon fibers as long-term implants in bone tissue engineering,5 owing to its stable structure and nondegradable behavior. Recently, natural silk fibers have been applied as effective reinforcements for epoxy resin matrix in structural composites due to excellent tensile and flexural ductility and toughness combined with respectable modulus and strength.6−10 These advantages can be implemented to improve the ductility and toughness of epoxy resin matrix composites.7,11,12 Interface properties between the fiber and the matrix are critical in defining the mechanical performance of fiber reinforced plastics. Thus far, a few studies13−15 have focused on surface modifications of silk fibers to enhance the interface adhesion between the fiber and the epoxy resin matrix. However, the detailed molecular interactions between silk fibers and epoxy polymers still lack investigation. © 2019 American Chemical Society

Received: February 19, 2019 Revised: May 5, 2019 Published: May 9, 2019 2295

DOI: 10.1021/acs.biomac.9b00260 Biomacromolecules 2019, 20, 2295−2304

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Scheme 1. (a) Chemical Structures of the Epoxy System 1564 Including a Bisphenol Epoxy and an Aliphatic Epoxy; (b) Chemical Structures of the Hardener System 3486; (c) Proposed Curing Reactions between Epoxy Groups and Amino Groups; (d) Proposed Curing Reactions between Epoxy Groups and Phenyl Hydroxyl Groups and Amino Groups from Silk Fibroin

Scheme 2. Schematic Illustration of Silk−Epoxy Blend Film Preparation

performance of SF-based or silk fiber reinforced epoxy composites.

casting may result in much higher tensile ductility and improved toughness of RSF films, as reported in previous works.26,27 Therefore, in this work, HFIP solvent has been chosen for ductile and tough silk−epoxy resin films. This work set out to blend the “thermoplastic” biopolymer silk fibroin with the thermoset polymer of epoxy resin and the hardener system and explored the compatibility and molecular interactions of the two polymers. These molecular interactions are complex but critical for understanding the properties and functions of silk-based composite materials. Most importantly, the chemical and physical crosslinking mechanisms in silk fibroin and epoxy polymers could lead to new strategies in improving the interface strength and overall mechanical

2. EXPERIMENTAL SECTION 2.1. Materials. Epoxy polymer 1564 and hardener 3486 from Huntsman Company (UT) with a mass ratio of 100:34 were mixed in the formation of epoxy resin thermoset. The chemical structures of the compounds are shown in Scheme 1a,b. The curing reactions between epoxy groups and diamine or polyols from the hardeners are shown in Scheme 1c. Degummed B. mori silk fabrics were purchased from Huzhou Yongrui Textile Co. Ltd. (Zhejiang Province, China). The possible reactions between epoxy groups and hydroxyl groups or amine groups from silk fibroin are shown in Scheme 1d. 2296


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Figure 1. FT-IR spectra of the silk−epoxy blend films with different ratios (10:0, 7:3, 5:5, 3:7, and 0:10) between silk and epoxy resin by weight; (a, c) unheated films and (b, d) heated films. The contents of different secondary structures are shown in (e) unheated films and (f) heated films. For each composition, three spectra were used for peak deconvolution and the means ± standard error were calculated for the conformation contents. 2.2. Blend Film Preparation. Regenerated silk fibroin or RSF aqueous solution was prepared from degummed B. mori silk fiber according to a previously reported protocol.28 Dry silk fibroin sponge was acquired through lyophilization of the aqueous silk solution. The silk fibroin and the premixed epoxy resin system were weighed and mixed at varied ratios, and further mixed in 1,1,1,3,3,3-hexafluoro-2propanol (HFIP). To cast films, the mixtures were placed in Teflon Petri dishes and then dried for 24 h in a fume hood to allow evaporation of HFIP. Mass ratios between silk fibroin and the epoxy resin system for the blends were chosen as follows: 10:0, 7:3, 5:5, 3:7, and 0:10. A heat treatment (80 °C for 8 h) on dry films was added to alter the degree of curing reactions of the epoxy groups. The preparation process of silk−epoxy blend films is shown in Scheme 2. The thickness of the films was ∼60 μm. 2.3. Thermal Analyses. Thermal gravimetric analysis (TGA) and differential scanning calorimetry (DSC) of all of the films were tested simultaneously by TGA/DSC2 (Mettler Toledo, Greifensee, Switzerland). About 5−8 mg of the samples was placed in an aluminum pan. Heating was applied from 30 to 500 °C at a rate of 20 and 5 °C/min under a nitrogen/argon atmosphere. Dynamic mechanical thermal analysis (DMTA) measurements were conducted on the films on a rheometer (Modular Compact Rheometer 102, Anton Paar, Graz, Austria) under a torsion mode at a frequency of 1 Hz and a heating rate of 4 °C/min from 25 to 300 °C. The dynamic strain was set at 0.5%. The size of all of the rectangular samples was 5 × 20 mm2. Additional DMTA measurements were conducted on blend films from SF and epoxy compounds on a dynamic mechanical analysis (DMA) Q800 instrument (TA Instruments, DE) under a tension mode at a frequency of 1 Hz, a heating rate of 3 °C/min, and a dynamic strain of 0.1%.

2.4. Synchrotron Wide-Angle X-ray Scattering (WAXS) Measurements. WAXS spectra were measured at BL45XU beamline of SPring-8, Harima, Japan using an X-ray with a wavelength of 0.1 nm and an energy of 12.4 keV. The sample−detector distance and exposure time were set as 187 mm and 10 s. The crystallinity was calculated by dividing the whole spectra area into crystalline and amorphous areas through an automatic and standard process.8 Lorentzian functions were used for curve fitting with software Igor Pro 6.3 (WaveMetrics, Inc., Portland, OR). 2.5. Fourier Transform Infrared (FT-IR) Spectroscopy Tests and Analyses. The infrared spectroscopic measurements of all of the films were carried out on an IRPrestige-21 Fourier transform infrared spectrophotometer (Shimadzu Corporation, Kyoto, Japan) over a wavenumber range of 700−4000 cm−1 under an ATR mode. 32 scans were accumulated for each spectrum at 25 °C. Three specimens were tested to obtain triplicates for each film sample. The peak deconvolution of conformations in the amide I region (1600−1700 cm−1) was conducted using PeakFit 4.0. The peak type was Gaussian; the peak width at half height (full width at half maximum) was set as 5 cm−1; the peak number was 9; and the peak positions were at 1605− 1615, 1622−1627, 1628−1637, 1638−1646, 1647−1655, 1656− 1662, 1663−1670, 1671−1685, 1686−1696 cm−1. The peaks were assigned to conformations according to ref 29. 2.6. Tensile Mechanical Tests. Tensile mechanical tests were conducted on a mechanical testing apparatus (EZ-LX/TRAPEZIUM X, Shimadzu, Kyoto, Japan). The displacement rate was set at 0.5 mm/min. The relative humidity was controlled at 60% and the size of all of the samples was 3.5 × 15 mm2. 5 samples of each formulation were tested. The fracture morphologies of the films were observed on 2297


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Figure 2. WAXS results of the pure silk film (10:0), silk−epoxy blend films with varied ratios (7:3, 5:5, and 3:7) and a pure epoxy resin film (0:10). (a) WAXS 1D profiles of unheated films; (b) WAXS 1D profiles of heated films; (c) degree of crystallinity calculated from WAXS 1D profiles; (d) schematic map of two different nucleation paths for β-sheet crystallization process: homogeneous nucleation and heterogeneous nucleation.

groups in silk fibroin and epoxy groups, a peptide poly-L-lysine containing multiple amino groups was synthesized and mixed with epoxy resin polymers. From MALDI-TOF-mass spectrometry (Figure S1), the peptide was shown to have zaveraged molecular weights of 696 (n = 5) and 824 (n = 6). Butyl glycidyl ether containing one epoxy group and poly-Llysine were then mixed to react at 80 °C for 8 h (the same conditions were applied for the heat treatment of blend films). The 1H NMR spectra for butyl glycidyl ether, poly-L-lysine, and the reaction product are compared in Figure S2. The chemical shifts of hydrogen H (g) associated with the epoxy group in the reactant and ether bond in the product could prove that the crosslinking reaction between amino groups and epoxy groups occurred. The blend ratio was found to dramatically affect the β-sheet content of silk. The conformation structure of silk fibroin in the amide I region was not affected by the absorbance of epoxy groups or the ether products from epoxy resin polymers. From Figure 1e,f, the β-sheet content (∼34%) in the silk−epoxy blend film (7:3) was revealed to be the highest among all of the blend films for unheated films. This suggests that in the 7:3 composition epoxy resin polymers could induce more physical crosslinks of β-sheet in silk, probably through a heterogeneous nucleation process, as discussed later. Wide-angle X-ray scattering (WAXS) was used to investigate the crystalline structure of the silk−epoxy resin blend films. The one-dimensional (1D) WAXS results in Figure 2a,b of the epoxy resin films (both unheated and heated) showed a dominant amorphous phase. In contrast, silk films and blend films showed sharp crystalline peaks with d-spacing of 0.45 nm of the (110) plane, which is typical of the β-sheet structure in silk.33,34 A procedure8 was applied to calculate the crystallinity of these samples. The crystallinities of the heated silk films prepared from HFIP and heated epoxy resin films were 13.1 and 2.0%, respectively (Figure 2c). The crystallinity of blend films decreased with increased content of the epoxy resin polymer. It was thus assumed that the crystallinity in the blends was contributed mainly by the β-sheet crystallinity in silk polymers. Notably, heating of the silk−epoxy blend films

a scanning electron microscope (SEM), JCM-6000 (JEOL, Tokyo, Japan) at 5 kV acceleration voltage under a secondary electron mode. 2.7. Synthesis of Poly-L-lysine. Poly-L-lysine was synthesized using the papain catalyst in H-Lys(Boc)−OMe·HCl at 40 °C for 2 h according to a previously reported protocol,30 and then deprotected with trifluoroacetic acid. The product was characterized by matrixassisted laser desorption ionization time-of-flight (MALDI-TOF)mass spectrometry and 1H NMR spectra. The MALDI-TOF mass system (Autoflex Speed; Bruker, Bremen, Germany) was employed to check the molecular weight and degree of polymerization. The ScoutMTP ion source was adopted in a reflector mode with a nitrogen laser. 1H NMR spectra were acquired on a Varian system 500 NMR spectrometer (500 MHz) using VnmrJ software (Agilent Technologies, Santa Clara, CA).

3. RESULTS AND DISCUSSION 3.1. Molecular Structure and Crosslinks in Blend Films. FT-IR results are presented in Figure 1a−d. Amide I (1600−1700 cm−1) and amide III (1200−1300 cm−1) corresponding to the C−N, N−H, and CO bond stretch modes are focal regions for conformation analysis.31,32 Peaks centered at 1651 cm−1 in amide I and 1225 cm−1 in amide III, assigned as helices and random coils, are dominant in both unheated and heated silk films. Nevertheless, the 1624 cm−1 peak corresponding to β-sheet conformation for the unheated silk−epoxy blend films appeared stronger compared to the heated. Quantitative analyses of the secondary structures (βsheet, random coils, helices, and turns) of amide I bands were made through peak deconvolution following an established procedure.29 The results of conformation contents are shown in Figure 1e,f. Evidently, heating resulted in more β-sheet content for the pure SF film from 19% for the unheated to 25% for the heated, but less β-sheet content for the blend films. It suggests that heating in the presence of epoxy resin polymers might suppress the formation of β-sheet structure in silk. It is supposed that the heating process promoted the chemical crosslinks that affected the physical crosslinks in silk fibroin. The characteristic peaks of tyrosine and lysine residues that were expected to react with epoxy groups, showed little difference after heating. To prove the reaction between amine 2298


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Figure 3. (a, c) Thermogravimetric TG and (b, d) differential thermogravimetric curves of the pure silk film (10:0), silk−epoxy blend films with varied ratios (7:3, 5:5, and 3:7), and the pure epoxy resin film (0:10) from 30 to 500 °C at a heating rate of 20 °C/min in nitrogen gas.

the thermal properties of silk−epoxy blend films. The results are summarized in Figures 3 and S3. The pure silk film showed two major peaks centered at 100 and 273 °C (unheated)/280 °C (heated), whereas the two loss events of the epoxy resin film occurred at 137/160 °C and 394/392 °C for the unheated and heated films. The first weight loss event in pure silk can be attributed to water evaporation as proposed in ref 8, and the 273−280 °C loss event can be attributed to the thermal degradation of silk fibroin according to the literature.35−37 For the silk−epoxy blend films of ratios 7:3, 5:5, and 3:7, there appeared three major loss events, two of which corresponded well with the loss events of epoxy resin. The peak at 300 °C may be linked to silk degradation. It is found that the silk degradation peak temperature increased with the increasing epoxy resin content. The DSC results in Figure S3 also showed that this thermal degradation temperature for silk increased in the blends. The exothermal event of pure epoxy resin at 390 °C in DSC thermographs could be degradation due to the breakage of ether bridges and other weak bonds.38,39 The glass transition behavior of the films can be captured in DSC measurements as indicated in Figure S3. The glass transition of epoxy resin was invisible in the DSC thermographs. In contrast, the glass transition of silk around 170 °C was captured and agreed with previous works.40,41 The glass transition events of single components in blend systems can be used to infer the miscibility of the components.42,43 An increase in the Tg of silk for the heated 7:3 blend was observed, which agreed with the DMA results below. However, the glass transition of silk in other blend films was very weak in the DSC thermographs. Dynamic mechanical (thermal) analysis (DMA) is an effective tool in analyzing the glass transition behavior of polymers.44−46 Figure 4 shows the storage modulus (G′) and loss tangent tan δ (as a ratio of loss modulus G″ over storage modulus G′) as a function of temperature for the films. The sharp drop of storage modulus G′ denotes the glass transition, so does the loss tangent tan δ peak. The tan δ peak temperature is often taken as the glass transition temperature

appeared to induce slightly lower crystallinity. The coupled effects of blend ratios and heating resulted in the highest crystallinity (about 15.9%) for the unheated 7:3 blend film. To explain the crosslinking structure of the blend films, a two-step formation mechanism of the β-sheet crystal is proposed in Figure 2d according to the previous literature.18 A mesophase of the monolayer β-sheet in the nucleation process was initially formed, and then turned into stacked βsheet···β-sheet crystals in the second step. It is noted that the β-sheet content from FT-IR increased with the introduction of epoxy resin for the unheated films, but the overall crystallinity from WAXS decreased with the introduction of epoxy resin. This may be because some dispersed monolayer β-sheet cannot join the stacked β-sheet crystalline regions to become X-ray detectable. In addition, the crystallinity was calculated relative to the whole blend which included noncrystalline epoxy resin. Both FT-IR and WAXS results showed that the 7:3 blend showed the highest value in both the β-sheet conformation and β-sheet crystallinity, owing to the heterogeneous nucleation from the chemical crosslinks. Compared to homogeneous nucleation, the heterogeneous nucleation process reduced the nucleation Gibbs free energy barrier between the initial phase and the final phase. During the heterogeneous nucleation process, the epoxy resin polymer network could restrict the mobility of silk fibroin chains to facilitate the initial bonding of silk chains or nucleation of βsheets to enhance β-sheet formation. Thus, the epoxy resin could act as heterogeneous agents to promote the nucleation of silk fibroin. Heating could induce the glass transition of epoxy resin at 80 °C (as discussed in the next section), and the segmental mobility of epoxy resin could enable the silk structure to change locally. However, heating did not seem to affect the crystallinity much. We cannot explain why heating could reduce the β-sheet content, which will need further investigation. 3.2. Thermal Stability and Glass Transition Behavior of Blend Films. Thermal gravimetric (TG) and differential scanning calorimetry (DSC) measurements were analyzed for 2299


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phenolic hydroxyl groups as well as amino groups from silk fibroin. It is also noted that the Tg peak for the hardener-absent system was singular, suggesting a miscible system of silk and epoxies, which was different from the two-Tg systems of the silk−epoxy resin blends. 3.3. Tensile Mechanical Properties. Tensile mechanical properties of the films (five specimens for each composition) were characterized at room temperature under relative humidity (55%). Typical stress−strain curves of pure silk, silk−epoxy resin blend, and pure epoxy resin films are presented in Figure 5 and all of the curves can be found in

Figure 5. Tensile stress−strain curves of the pure silk film (10:0), silk−epoxy blend films with varied ratios (7:3, 5:5, and 3:7), and the pure epoxy resin film (0:10): (a) before heat treatment and (b) after heat treatment.

Figure S4. Generally, within the unheated group, the pure epoxy resin film was the most ductile and weak, whereas the pure silk film was stronger, but the three blends showed higher tensile strength than the pure films. For the heated group, the pure silk film showed the lowest strength, whereas pure epoxy resin and the blends showed a similar strength of ∼60 MPa. This indicated that the tensile strength for the heated film might be determined by the chemical crosslinks in the epoxy resin polymer network. The tensile fracture morphology of films in Figures 7 and S5 also revealed some tensile failure characteristics of these films. The heated pure epoxy resin film showed brittle fracture with a clean path without crazes. In contrast, close to the fracture region, regular crazes could be found in the surface of silk and silk−epoxy resin (7:3) films. The fracture surface showed rough morphologies and rich paths of crack propagation in silk and the silk−epoxy resin film (7:3). From the fracture morphology, we confirm that the two components were able to blend well in the microscale, and the interactions should be in the nano or molecular scale. The ductile and gradual fracture contributed to the improved toughness, and the deflective fracture path indicated that the nanoheterogeneous phases with either chemical or physical crosslinks responded differently to mechanical stretch. Overall, the 7:3 blends showed significantly enhanced tensile strength and toughness. The key mechanical properties including Young’s modulus, tensile strength, elongations, and breaking energy were compared in Figure 6. Breaking energy was calculated as the area under the stress−strain curve and the unit was converted from MPa to MJ/m 3 for easy comparison with the literature.47−49 It showed that pure epoxy resin films possessed the lowest Young’s modulus (0.5 GPa for the unheated and 1.2 GPa for the heated). All of the silk-containing films were stiffer, and the introduction of silk fibroin proved a stiffening effect on the blends. Unheated epoxy resin behaved ductile and weak, with the highest tensile elongation (84%) but with the lowest

Figure 4. DMA plots of (a, b) G′ and (c, d) tan δ as a function of temperature from 25 to 300 °C at 4 °C/min for the pure silk film (10:0), silk−epoxy blend films with varied ratios (7:3, 5:5, and 3:7), and the pure epoxy resin film (0:10) after heat treatment. Peak decomposition of tan δ (e, f). tan δ as a function of temperature from 25 to 250 °C at 4 °C/min for silk−epoxy blend films (7:3) without the hardener (g).

Tg. Glass transitions of both silk and epoxy resin are present in DMA graphs. The Tgs were 170 and 90 °C for the pure silk film and the epoxy resin film, respectively. The silk Tg agreed with previous reports.15,41 In the blend films, both glass transitions were suppressed, as shown by the reduced magnitude of the tan δ peak (maximum value ∼0.2). However, the two peaks remained distinguishable through peak decomposition analysis as shown in Figure 4c,f, indicating independent segmental motions from epoxy resin and silk. Interestingly, the Tgs in the 7:3 blend increased compared to the pure component. The separate glass transition events should infer that the two polymers were not fully miscible in the segmental level, although the glass transition behaviors were affected by the intermolecular interactions between these two polymers. To prove the chemical crosslinking reactions between silk and epoxies, silk fibroin was only mixed with epoxies in the mass ratio 7:3 without hardeners. The Tg of this blend was 191 °C, much greater than 170 °C for the pure silk fibroin film (Figure 4g). This demonstrated epoxy groups could react with 2300


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Figure 6. Tensile mechanical properties of the pure silk film (10:0), silk−epoxy blend films with varied ratios (7:3, 5:5, and 3:7), and the pure epoxy resin film (0:10) before and after heat treatment. (a) Young’s modulus, (b) tensile strength, (c) elongation, and (d) toughness calculated as the area enclosed by the stress−strain curves, and the unit was transformed from MPa to MJ/m3 for easy comparison with the literature.

Figure 7. SEM images of fracture morphologies: top view of the film surface close to the fracture region (a-1, b-1, c-1) and side view of fractured cross-sections (a-2, b-2, c-2) of the pure silk film (a), 7:3 silk−epoxy blend film (b), and the pure epoxy film (c). The inset shows the fracture path of films.

strength of 15.3 MPa. After heating/curing, the elongation decreased drastically to only 9.3%, but the strength increased to 56.8 MPa. Heating also helped the pure silk film to increase the strength to 42.5 MPa. Overall, heating proved to be beneficial in enhancing the strength of pure or blend films through chemical/physical crosslinks. On the significance level of p < 0.05, the heated 7:3 blend had significantly greater tensile modulus than pure epoxy resin film, significantly greater tensile strength (61 MPa) than pure silk film, and a relatively high tensile elongation of 36%. Therefore, this blend reached the highest toughness of 13.6 MJ/m3, 370% of the heated epoxy resin film (3.6 MJ/m3) (Figure 6, Figure 7). 3.4. Discussion: Crosslinking Structure of Silk−Epoxy Resin Blend. The smooth and even surface morphology in Figure S6 of the 7:3 silk−epoxy blend film suggested a homogeneous structure in the microscale. Phase separation was invisible in the microscale. The structure characterization showed that on the molecular level, silk and epoxy resin interacted through both physical crosslinks and chemical crosslinks, and separate glass transition behaviors indicated phase separation in the nanoscale. The chemical crosslinks of epoxy with silk could affect the kinetics of β-sheet formation of silk polymers, especially when heating was applied. The chemical crosslinking sites in silk could restrict the molecular motions of silk chains, which may prevent a single β-sheet structure developing into a stable crystalline structure. On the other hand, the fast-forming chemical crosslinks may act as heterogeneous nucleation sites for β-sheet formation. To illustrate the competition and cooperation of chemical and physical crosslinks in the silk−epoxy resin blend, a structure sketch is proposed in Figure 8. It demonstrates that as the silk/ epoxy resin ratio changes, the crosslinking structure evolves, including the β-sheet crosslinks within silk polymers and chemical crosslinks between epoxy resin and silk. This structure sketch also demonstrates that due to the much higher reactivity between epoxy and hardeners, silk polymers and epoxy resin polymers are partially miscible and form

relatively independent nanophases of crosslinking structures. This could be compared with a mechanism called the chemical reaction-induced phase separation in the literature.50 For the 7:3 blend, while the β-sheet crosslinks dominate and stabilize the silk polymer structure, the chemical crosslinks with the epoxy polymer network enhance the overall structure and improve the mechanical performance. The tensile mechanical properties of silk−epoxy resin blends may be explained by the structure data. Here, only heated films with fully developed chemical crosslinks are considered. It is interesting to find that the tensile modulus of the blend appeared to maintain a value similar to that of pure silk fibroin, whereas the tensile strength of the blend appeared to maintain a value similar to that of the pure epoxy resin film. Although the crystallinity from WAXS decreased with the increased content of the epoxy resin polymer, the tensile modulus or strength of the blends did not decrease. One possible reason is that the β-sheet structures were distributed as local crosslinks, but the β-sheet crystalline regions did not form the percolating network. In contrast, the chemical crosslinks from epoxy resin polymers formed a percolating network, which could determine the strength of the overall crosslinking network. Therefore, it is proposed that in the silk−epoxy resin blend the modulus is determined by both local physical and chemical crosslinks, but the strength is determined by the percolating chemical crosslinks. For the 7:3 blend, owing to the mechanism of chemical crosslinking-induced phase separation, the structure contains both chemical crosslinks and physical crosslinks. Thus, the mechanical properties including the tensile modulus and the strength reached an optimal performance, which could offer new routes to enhance silk polymers by introducing percolative chemical crosslinks. For the interface design in silk fiber reinforced epoxy resin composites, this work suggests that the interface interactions could be enhanced by chemical crosslinks if more amorphous silk structure could be exposed to epoxies in the matrix. 2301


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Figure 8. Schematic illustrations of the crosslinking structure in silk−epoxy resin blends with three mix ratios of 7:3, 5:5, and 3:7.

Through controlled solvent treatment (e.g., LiBr15), the highly ordered structure on the fiber surface could be partially unraveled without harming the internal fiber structure or the βsheet crystalline structure. Under elevated temperatures, the chemical crosslinks between silk and epoxy at the interface would help form a stronger interfacial phase, which would improve the overall mechanical performance of the silk−epoxy resin composite.

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blend films (stress−strain curve and fracture surface after the tensile test) (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (K.N.). *E-mail: [email protected] (J.G.). ORCID

4. CONCLUSIONS This work studies the molecular interactions between silk fibroin and an epoxy resin system, especially the evolution of physical and chemical crosslinks with blend ratios and heating treatment. We conclude that (i) the heterogeneous nucleation mechanism may explain the increased β-sheet content in silk after the introduction of epoxy resin; (ii) the blends showed independent glass transition events from epoxy resin segments and silk segments, suggesting phase separation of the two components on the segmental level; (iii) silk fibroin was proved to react with the epoxy group to induce higher-Tg glass transition; (iv) blending the thermoplastic silk polymer with thermosetting epoxy resin could result in a strong network structure with both physical crosslinks and chemical crosslinks to achieve enhanced mechanical properties. The 7:3 silk− epoxy resin blend possessed a crosslinking structure with the greatest β-sheet crystallinity (15.9%) and a percolative chemical crosslinking structure, which explained the highest strength (60.1 MPa) and toughness (13.6 MJ/m3) among all of the blends. The findings would help to understand the interface interactions and improve the interface strength between silk fibers and the epoxy resin matrix in highperformance structural composites.

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Keiji Numata: 0000-0003-2199-7420 Juan Guan: 0000-0002-9784-8125 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Prof. Zhengzhong Shao from Fudan University for constructive comments to improve the manuscript. K.Y. thanks the International Program Associate (IPA) office and the Biomacromolecules Group in RIKEN. K.N. acknowledges the Impulsing Paradigm Change through Disruptive Technologies Program (ImPACT) and RIKEN Engineering Network Program. J.G. acknowledges the Fundamental Research Funds for Central Universities. The authors thank Drs Takaaki Hikima and Hiroyasu Masunaga for SPring-8 measurements.

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REFERENCES

(1) Jin, F. L.; Li, X.; Park, S. J. Synthesis and application of epoxy resins: A review. J. Ind. Eng. Chem. 2015, 29, 1−11. (2) Carter, J. T.; Emmerson, G. T.; Faro, C. L.; Mcgrail, P. T.; Moore, D. R. The development of a low temperature cure modified epoxy resin system for aerospace composites. Composites, Part A 2003, 34, 83−91. (3) Feng, J.; Guo, Z. Temperature-frequency-dependent mechanical properties model of epoxy resin and its composites. Composites, Part B 2016, 85, 161−169. (4) Wong, D. W. Y.; Lin, L.; Mcgrail, P. T.; Peijs, T.; Hogg, P. J. Improved fracture toughness of carbon fibre/epoxy composite laminates using dissolvable thermoplastic fibres. Composites, Part A 2010, 41, 759−767. (5) Ding, G. Latest progress of carbon-fiber-reinforced polymer energy-storing transtibial prostheses. Chin. J. Tissue Eng. Res. 2013, 17, 8228−8234. (6) Wang, Y.; Wang, X.; Shi, J.; Zhu, R.; Zhang, J.; Zhang, Z.; Ma, D.; Hou, Y.; Lin, F.; Yang, J.; Mizuno, M. A biomimetic silk fibroin/

ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.biomac.9b00260. Additional characterization of poly-L-lysine and reaction product (MALDI-TOF-mass spectrum and 1H NMR spectrum), differential scanning calorimetry (DSC) thermographs and additional mechanical properties of 2302


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Molecular Interactions and Toughening Mechanisms in Silk Fibroin

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