Molecular interactions and toughening mechanisms in silk fibroin

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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 Biomacromolecules, Just Accepted Manuscript • DOI: 10.1021/acs.biomac.9b00260 • Publication Date (Web): 09 May 2019 Downloaded from http://pubs.acs.org on May 9, 2019

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Biomacromolecules

Molecular interactions and toughening mechanisms in silk fibroin-epoxy resin blend films Kang Yang,1,2 Kenjiro Yazawa,2 Kousuke Tsuchiya,2 Keiji Numata,2* Juan Guan,1,3* 1.

Intl. Research Center for Advanced Structural and Biomaterials, School of Materials Science

and Engineering, Beihang University, Beijing, 100191, China 2.

Biomacromolecules Research Team, RIKEN Center for Sustainable Resource Science, 2-1

Hirosawa, Wako-shi, Saitama 351-0198, Japan 3.

Beijing Advanced Innovation Center for Biomedical Engineering, Beihang University, Beijing,

100083, China

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 fibre-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 behaviour 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 silkepoxy composites.

Keywords: 1

ACS Paragon Plus Environment

Biomacromolecules 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

glass transition, biopolymer, polymer matrix composites, thermoset, toughness, cross linking.

1. Introduction Epoxy resin polymers are widely used as adhesives, coatings and matrix materials in high-performance polymer composites for automotive, marine and aerospace applications1-3. High modulus and strength of most epoxy resin thermosets are caused of 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 toughness4. In addition, the epoxy resin thermosets with high modulus and strength are popularly applied as a matrix for carbon fibres as long-term implants in bone tissue engineering5, owing to its stable structure and non-degradable behaviour. Recently, natural silk fibres 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 strength6-10. These advantages can be implemented to improve the ductility and toughness of epoxy resin matrix composites

7, 11, 12.

Interface properties between the fibre and the matrix are

critical in defining the mechanical performance of fibre reinforced plastics (FRPs). Thus far, a few studies13-15 have focused on surface modifications of silk fibres to enhance the interface adhesion between the fibre and the epoxy resin matrix. However, the detailed molecular interactions between silk fibres and epoxy polymers still lack investigation. Silk fibroin or SF regenerated from silk fibre has been used as a popular biomaterial for a wide range of biomedical applications, owing to its “green” origin, 2


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biocompatibility and biodegradability. Alike silk fibre 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 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 HFIP

solvent to dissolve freeze-dried silk fibroin for film 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 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

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improving the interface strength and overall mechanical performance of SF-based or silk fibre reinforced epoxy composites.

2. Experimental 2.1 Materials

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. Epoxy polymer 1564 and hardener 3486 from Huntsman Company (Utah, USA) 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 1(a) and 1(b). The curing

4


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reactions between epoxy groups and diamine or polyols from the hardeners are shown in Scheme 1(c). 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 1(d).

2.2 Blend film preparation

Scheme 2 Schematic illustration of silk-epoxy blend film preparation. Regenerated silk fibroin or RSF aqueous solution was prepared from degummed B.mori silk fibre according to a previously reported protocol28. Dry silk fibroin sponge was acquired through lyophilization of the aqueous silk solution. The silk fibroin and the pre-mixed epoxy resin system were weighed and mixed at varied ratios, and further mixed in 1,1,1,3,3,3-hexafluoro-2-propanol (HFIP). To cast films, the mixtures were placed in Teflon petri-dishes, and then dried for 24h in a fume hood to allow evaporation of HFIP. Mass ratios between silk fibroin and 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 oC for 8h) on dry films was added to alter the degree of curing reactions of the

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epoxy groups. The preparation process of silk-epoxy blend films is shown in Scheme 2. The thickness of the films was ~60 microns.

2.3 Thermal analyses Thermal gravimetric analysis (TGA) and differential scanning calorimetry (DSC) of all 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 oC at a rate of 20 oC/min and 5 oC/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 torsion mode at a frequency of 1 Hz and a heating rate of 4 oC/min from 25 to 300 oC. The dynamic strain was set as 0.5%. The size of all the rectangular samples were 5 mm × 20 mm. Additional DMTA measurements were conducted on blend films from SF and epoxy compounds on a DMA Q800 instrument (TA Instruments, USA) under tension mode at a frequency of 1 Hz, a heating rate of 3 oC/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 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 FT-IR tests and analyses 6


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The infrared spectroscopic measurements of all the films were carried out on IRPrestige-21 Fourier transform infrared spectrophotometer (Shimadzu Corporation, Kyoto, Japan) over a wavenumber range of 700-4000 cm-1 under ATR mode. 32 scans were accumulated for each spectrum at 25 oC. Three specimens were tested to obtain triplicates for each film sample. The peak de-convolution 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 (FWHM) 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 reference29.

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 as 0.5 mm min-1. The relative humidity was controlled as 60% and the size of all the samples was 3.5 mm× 15 mm. 5 samples of each formulation were tested. The fracture morphologies of the films were observed on a scanning electron microscope (SEM), JCM-6000 (JEOL, Tokyo, Japan) at 5 kV acceleration voltage under secondary electron mode.

2.7 Synthesis of Poly-L-lysine Poly-L-lysine was synthesized using papain catalyst in H-Lys(Boc)-OMe·HCl at 40 oC for 2h according to a previously reported protocol30, and then deprotected with TFA. The product was characterized by MALDI-TOF-mass spectrometry and 1H NMR spectra. The MALDI-TOF-mass system (Autoflex Speed;Bruker, Bremen, Germany)

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was employed to check the molecular weight and degree of polymerization. ScoutMTP ion source was adopted in reflector mode with nitrogen laser. 1H NMR spectra was acquired on Varian system 500 NMR spectrometer (500 MHz) by 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 Figures 1(a-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 (β-sheet, random coils, helices and turns) of amide I band were made through peak de-convolution following an established procedure

29.

The results of

conformation contents are shown in Figure 1(e) and (f). Evidently, heating resulted in more β-sheet content for 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 which 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. In order to prove the

8


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reaction between amine groups in silk fibroin and epoxy groups, a peptide poly-Llysine 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 z averaged molecular weight of 696 (n=5) and 824 (n=6). Butyl glycidyl ether containing one epoxy group and poly-L-lysine were then mixed to react at 80 oC for 8h (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.

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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, 0:10) between silk and epoxy resin by weight. (a,c) Unheated films; (b,d) Heated films. The contents of different secondary structure were shown in (e) Unheated films; (f) Heated films. For each composition, three spectra were used for peak deconvolution and the Means ± SE were calculated for the conformation contents. The blend ratio was found to dramatically affect the β-sheet content of silk. The conformation structure of silk fibroin in Amide I region was not affected by the 10


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absorbance of epoxy groups or the ether products from epoxy resin polymers. From Figure 1(e,f), the β-sheet content (~34 %) in silk-epoxy blend film (7:3) was revealed to be the highest among all the blend films for unheated films. This suggests that in the 7:3 composition epoxy resin polymers could induce more physical crosslinks of βsheet in silk, probably through a heterogeneous nucleation process, as discussed later.

Figure 2 WAXS results of pure silk film (10:0), silk-epoxy blend films with varied ratios (7:3, 5:5, 3:7) and 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. Wide angle X-ray scattering (WAXS) was used to investigate the crystalline structure of the silk-epoxy resin blend films. The 1D WAXS results in Figure 2(a,b) of the epoxy resin films (both unheated and heated) showed a dominant amorphous 11



phase. In contrast, silk films and blend films showed sharp crystalline peaks with dspacing 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 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 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 β-sheet crystal is proposed in Figure 2d according to previous literature18. A meso-phase of monolayer β-sheet in the nucleation process was initially formed, and then turn into stacked β-sheet…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 βsheet 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 non-crystalline epoxy resin. Both FT-IR and WAXS results showed the 7:3 blend showed the highest value in both β-sheet conformation and β-sheet crystallinity, owing to the heterogeneous nucleation from the chemical crosslinks. Compared to homogeneous nucleation, heterogeneous nucleation process reduced the nucleation Gibbs free energy barrier between the initial phase and final phase. During the heterogeneous nucleation process, epoxy resin polymer network could restrict the

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mobility of silk fibroin chains to facilitate the initial bonding of silk chains or nucleation of β-sheets to enhance β-sheet formation. Thus, the epoxy resin could act as heterogeneous agents to promote the nucleation of silk fibroin. On the effect of heating, it could induce the glass transition of epoxy resin at 80 oC (as discussed in the next section), and the segmental mobility of epoxy resin enabled the silk structure to change locally. However, heating didn’t seem to affect much the crystallinity. We cannot explain why heating could reduce β-sheet content, which will need further investigation.

3.2 Thermal stability and glass transition behaviour of blend films Thermal gravimetric (TG) and differential scanning calorimetry (DSC) measurements were analysed for the thermal properties of silk-epoxy blend films. The results are summarized in Figure 3 and Figure S3. Pure silk film showed two major peaks centered at 100 oC and 273 oC (unheated) /280 oC (heated), whereas the two loss events of epoxy resin film occurred at 137 oC/160 oC and 394 oC/392 oC for the unheated and heated. The first weight loss event in pure silk can be attributed to water evaporation as proposed in 8, and the 273-280 oC loss event can be attributed to the thermal degradation of silk fibroin according to the literature35-37. 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 oC may be linked to silk degradation. It is found that the silk degradation peak temperature increased with increasing epoxy resin content. The DSC results in Figure S3 also showed this thermal degradation temperature for silk increased in the blends. The exothermal event of pure epoxy resin at 390 oC in DSC thermographs could be degradation due to the breakage of ether bridges and other weak bonds38, 39.

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

was captured, 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. 14


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Dynamic mechanical (thermal) analysis (DMA) is an effective tool in analyzing the glass transition behaviour of polymers

44-46.

Figure 4 shows the storage modulus

(G') and loss tangent tanδ (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 tanδ peak. The tanδ peak temperature is often taken as the glass transition temperature Tg. Glass transitions of both silk and epoxy resin are present in DMA graphs. The Tgs were 170 oC and 90 oC for pure silk film and 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 tanδ peak (maximum value ~0.2). However, the two peaks remained distinguishable through a peak decomposition analysis in Figure 4(c,f), indicating independent segmental motions from epoxy resin and silk. Interestingly, the Tgs in the 7:3 blend increased compared to 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 behaviours were affected by the inter-molecular interactions between these two polymers. In order to prove the chemical crosslinking reactions between silk and epoxies, silk fibroin was only mixed with epoxies in mass ratio 7:3 without hardeners. The Tg of this blend was 191 oC, much greater than 170 oC for pure silk fibroin film (Figure 4(g)). This demonstrated epoxy groups could react with 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.

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

3.3 Tensile mechanical properties

Figure 5 Tensile stress−strain curves of pure silk film (10:0), silk-epoxy blend films with varied ratios (7:3, 5:5, 3:7) and pure epoxy resin film (0:10): (a) before heat treatment and (b) after heat treatment. Tensile mechanical properties of the films (5 specimens for each composition) were characterized at room temperature under relative humidity 55%. Typical stressstrain curves of pure silk, silk-epoxy resin blend and pure epoxy resin films were presented in Figure 5 and all the curves can be found in Figure S4. Generally, within the unheated group, pure epoxy resin film was the most ductile and weak whereas pure silk film was stronger, but the three blends showed higher tensile strength than the pure films. For the heated group, pure silk film showed the lowest strength, whereas pure epoxy resin and the blends showed similar strength of ~60 MPa. This

17



indicated 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 Figure 7 and Figure S5 also revealed some tensile failure characteristics of these films. 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 facture surface showed rough morphologies and rich paths of crack propagation in silk and 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 the nano heterogeneous 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 literature47-49. 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 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 the lowest 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 pure silk film to increase 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

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significance level of p

Molecular interactions and toughening mechanisms in silk fibroin

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