Dual Cross-linked Epoxidized Natural Rubber Reinforced by Tunicate

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Dual Cross-linked Epoxidized Natural Rubber Reinforced by Tunicate Cellulose Nanocrystals with Improved Strength and Extensibility Liming Cao, Jiarong Huang, and Yukun Chen ACS Sustainable Chem. Eng., Just Accepted Manuscript • DOI: 10.1021/ acssuschemeng.8b03331 • Publication Date (Web): 11 Oct 2018 Downloaded from http://pubs.acs.org on October 15, 2018

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ACS Sustainable Chemistry & Engineering

Dual Cross-linked Epoxidized Natural Rubber Reinforced by Tunicate Cellulose Nanocrystals with Improved Strength and Extensibility Liming Cao,# Jiarong Huang,# Yukun Chen* Lab of Advanced Elastomer, South China University of Technology, 381 Wushan Road, Tianhe District, Guangzhou 510640, China # These two authors contributed equally to this work. Corresponding Author: Yukun Chen ([email protected])

Abstract: Cellulose nanocrystals have been proved to be ideal reinforcing fillers for rubber. However, the increase in tensile strength always accompanies by a reduction in extensibility. In this paper, epoxidized natural rubber (ENR) nanocomposites with dual cross-linking networks by incorporating tunicate cellulose nanocrystals (t-CNs) were prepared. Carboxyl groups were introduced onto t-CNs surface via grafting reaction with maleic anhydride (MAH). Interfacial chemical reaction between carboxyl groups on modified t-CNs and epoxy groups on ENR chains led to formation of covalent crosslinking network. Hydrogen bonds between hydroxyl groups on t-CNs and epoxy groups on ENR chains served as physical cross-linking network. Therefore, uniform dispersion and improved interfacial adhesion were achieved. Meanwhile, the covalent crosslinking network could provide strength and elasticity, while the reversible physical cross-linking network could serve as sacrificing element during stretching and dissipate energy. Therefore, dual cross-linked nanocomposites showed improved strength, modulus, extensibility and decreased hysteresis compared with sole cross-linked nanocomposites. Keywords: rubber nanocomposites, dual cross-linking networks, tunicate cellulose nanocrystals, carboxylation modification Introduction Polymer nanocomposites composed of polymer matrixes and nanoscale fillers have attracted numerous attention during the past decades.1-6 Nanoparticles, e.g., silica, carbon nanotubes, clay and graphene, etc., have been widely exploited. Among the potential nanofillers, cellulose nanocrystals (CNs), which are rod-like nanomaterials with high crystalline and high aspect ratio, have been widely investigated because of their excellent mechanical properties and their sustainability, renewability and abundance. 7,8 Favier et al. explored the utilization of CNs as reinforcing component in polymer nanocomposites firstly, and observed formidable mechanical enhancement.9 Since then, numerous studies on using CNs extracted from different sources to reinforce a variety of polymers have been explored.10-14 Rubber represents a class of strategically important polymeric materials due to their high elasticity and soft characteristic.15-19 However, unfilled rubbers usually show poor mechanical properties and it is essential to incorporate nanofillers into rubbers, which could achieve a combination of high strength and modulus, etc. Because of the outstanding physical properties and sustainability of CNs, there is a growing interest in using CNs to reinforce rubber.20-22 Pasquini20 and Smyth21 et al. investigated the

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reinforcement of CNs for rubbers and found that the properties of the obtained nanocomposites largely depend on the source of CNs, the higher aspect ratio, the better reinforcement. Pei et al. designed interfacial covalent bonds between CNs and polyurethane (PU) by adding CNs during the preparation of prepolymer, and synthesized ultrahigh tensile strength PU nanocomposites with very low CNs content.12 Previously, our group also studied the reinforcement of CNs in polar rubbers and foamed rubber, and achieved good reinforcements via constructing hydrogen bonds (Hbonds) between rubber and CNs.23-25 However, the compatibility between rubber matrix and CNs is relatively poor, and the interfacial adhesion between them is not strong enough to transfer stress from matrix to filler efficiently. To enhance interfacial adhesion between them, the modification of matrix or CNs have been extensively investigated recently. For example, Kanoth et al. introduced mercapto-groups onto surface of CNs, and realized both reinforcing and cross-linking of CNs in natural rubber (NR) via thiol-ene reactions. Improved dispersion and much higher mechanical properties were achieved after modification.26 Mariano et al. prepared oxidized NR (ONR) and promoted the interfacial interaction via H-bonds between the -OH groups on ONR and those on CNs surface.27 It has been proved that both chemical and physical interactions are effective ways to improve interfacial adhesion between rubber and CNs.26-28 In recent years, hydrogels with dual cross-linking networks have attracted numerous interest because of the fascinating mechanical properties.29-31 The chemical crosslinking network could provide strength and elasticity, whereas the physical crosslinking network, such as ionic interactions, H-bonds and host-guest interactions, etc. could serve as sacrificing element and consume energy. In addition, the reversible nature of physical cross-linking points could allow the materials to self-heal after disruption. Inspired by the current research, in this work, we intend to design rubber based nanocomposites with dual cross-linking networks. To achieve this, inspired by the esterification reactions between carboxyl groups/anhydride and epoxy groups,32-35 epoxidized natural rubber (ENR) was chosen as matrix and tunicate cellulose nanocrystals (t-CNs), which shows large aspect ratio, was used as fillers. Carboxyl groups were introduced onto t-CNs surface via ringopening esterification reactions with maleic anhydride (MAH).36-38 Rubber nanocomposites were prepared via latex mixing and co-coagulation method. By utilizing esterification reactions between carboxyl groups and epoxy groups, interfacial covalent cross-linking points were formed. Meanwhile, the oxygenous groups and -OH groups between the two phases led to formation of H-bonds, which could be used as reversible physical cross-linking points. Various characterizations were conducted on the nanocomposites. And the results indicate that compared with sole physically crosslinked nanocomposites, dual cross-linked nanocomposites showed optimal combination of strength and extensibility with lower hysteresis. Experimental section Materials. ENR latex with 25% solid content and an epoxidation degree of 40% was obtained from the Agricultural Products Processing Research Institute, Chinese


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Academy of Tropical Agricultural Sciences, Zhanjiang, China. According to our published work, t-CNs was extracted from tunic of tunicates (Halocynthia roretzi Drasche) via sulfuric acid hydrolysis method.25,39 MAH, N, N’-dimethylformamide (DMF), sulfuric acid, toluene and chloroform were supplied by Guangzhou Chemical Reagent Co., Ltd. All the reagents were used as received. Preparation of carboxylated t-CNs (mt-CNs). Carboxyl groups were introduced onto t-CNs surface by grafting MAH on the hydroxyl groups of t-CNs via ring-opening esterification.36-38 Aqueous suspension of t-CNs was solvent changed with DMF via rotary evaporation and obtained a DMF suspension with a solid content of 2%. Predetermined MAH was added into the suspension (t-CNs : MAH=1:5 (w/w)), and the mixture was stirred until MAH dissolved completely. Followed by the addition of catalyst H2SO4 (98 wt %). Then the esterification was performed at 120 oC for 5 h with continuous stirring. Followed the esterification reaction, excess ethanol was added to the resultant mixture and then centrifuged at 10000 r/min for 10min. The sediment was re-diluted and centrifuged for another three times for the complete removal of unreacted MAH. The obtained mt-CNs were finally dispersed in deionized water via ultrasound under 300 W for 1 h before use. Preparation of ENR/mt-CNs nanocomposites. Desired amounts of mt-CNs aqueous suspension were added to ENR latex slowly, and stirred for 1 h to achieve homogeneous dispersion. Subsequently, the suspension was co-coagulated with excess ethanol and washed several times with deionized water. The obtained floccule was air dried under 50 oC overnight to obtain a dry compound. Finally, ENR/mt-CNs nanocomposites with thickness of ~ 0.5 mm were compression molded under 180 oC and 10 MPa for 1 h. For comparison, nanocomposites with unmodified t-CNs were also prepared in the same way. The obtained ENR/mt-CNs and ENR/t-CNs nanocomposites were coded as ENR/mX and ENR/tX, respectively, where X represents parts of t-CNs in 100 parts of rubber (phr). Characterizations. Fourier Transform Infrared spectroscopy (FT-IR) were conducted on a TENSOR 27 spectrometer (BRUKER, Germany) under attenuated total reflectance (ATR) mode. All the measurements were performed in wavenumber range of 4000 cm-1 ~ 600 cm-1 for 32 scans and with resolution of 4 cm-1. Atomic force microscopy (AFM) measurements were conducted by a Bruker Multimode 8 (Germany). Specimen was obtained by dropping t-CNs suspension onto a freshly cleaved mica substrate and blown dry before observation. Transmission electron microscopy (TEM) was also conducted to observe the structure of the t-CNs. A droplet of diluted t-CNs suspension was placed onto a copper grid coated with amorphous carbon, the copper grid was then stained with phosphotungstic acid solution for 2 min before observation. The prepared sample was completely dried and observed using a JEOL JEM 2100 F operated at 200 kV. X-ray diffraction (XRD) were carried out by using BRUKER D8 ADVANCE. CuKα radiation was operated at 45kV and 40mA. Scans were carried out from 5o to 40o with



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a scan speed of 2o/min. Curing process was monitored by using a UR-2010 Rheometer (U-CAN Dynates Inc. China) under temperature of 180 oC, and torques were recorded over time. Cross-link densities of the samples were measured according to the following procedure: three pre-weighed (m0) specimens were immersed in toluene under ambient conditions. About 5 days later, the swollen specimens were taken out from the solvent, wiped up with filter paper and weighed (m1) immediately. Cross-link densities (ν) were calculated according to the Flory-Rehner equation:25,26

ve = -

ln(1-Vr ) + Vr + xV 2r 1 V1 (V 3r

(1)

V — r) 2

where

Vr =

ms ρs

mr ρr

(2)

m + r ρr

V1 and x represent the molar volume and interaction parameter of toluene (V1 = 106.2 cm3/mol, x = 0.3440), respectively. ρr and mr represent the density and weight of rubber in the sample (for ENR, ρr = 0.96 g/cm3), ρs and ms represent the density and weight of toluene (ρs = 0.865 g/cm3). Dynamic mechanical analysis (DMA) was performed on a NETZSCH 242 C under tension mode. All the specimens were measured from -60 ~ 30 oC with heating rate of 3 oC/min and frequency of 1 Hz. Differential scanning calorimetric (DSC) analysis was conducted on a NETZSCH 204 F1 under nitrogen atmosphere from -70 ~ 30 oC under heating rate of 10 oC/min. Transmission electron microscopy (TEM) was carried out to analyze the dispersion of CNs in rubber matrix. Specimens were cryo-microtomed under -80 oC on Lecia EM UC 6 system. And images were recorded using JEOL JEM-2100F at operating voltage of 200 kV. Scanning electron microscopy (SEM) observation was conducted on ZEISS Merlin apparatus working under 10 kV. Before observation, a thin layer of platinum were coated on both the frozen quenched and tensile fractured surfaces. A UCAN UT-2080 universal tensile testing machine was used to test mechanical properties of the nanocomposites. The tests were conducted at ambient conditions using dumbbell samples under strain rate of 500 mm/min. And five specimens were tested for each sample. The modulus was determined at 10 % strain, and work-of-fracture was obtained by calculating the area under stress-strain curve. Results and discussion Characterization of carboxylated t-CNs. In this work, t-CNs extracted from a sea animal termed tunicate (Halocynthia roretzi Drasche) via sulfuric acid hydrolysis


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approach was used as reinforcing filler for ENR. Figure 1 (a) and (b) are AFM image and height profiles of t-CNs. As can be clearly seen, t-CNs show good dispersibility in aqueous suspension with several aggregates. They show length between 500 nm ~ 2 μm with characteristic diameter of 10 ~ 20 nm, which are comparable to the reported values.41,42 Figure 1 (c) is the TEM image of the obtained t-CNs, which also show similar sizes with the AFM image.

Figure 1. (a) AFM image of t-CNs; (b) transverse height profiles determined from AFM and (c) TEM image of t-CNs. Carboxyl groups could be introduced onto CNs surface by TEMPO-mediated oxidation43,44 or esterification with MAH.36 In this work, ring-opening esterification reaction with MAH was carried out to introduce carboxyl groups onto t-CNs surface. Figure 2 (a) displays FT-IR spectra of unmodified and modified t-CNs (mt-CNs). As can be seen, several typical absorption peaks of cellulose, e.g., 3340, 2920 and1060 cm1 are observed in both t-CNs and mt-CNs, which correspond to -OH stretching, C-H vibration of -CH2 groups and C-O stretching vibrations, respectively, indicating that cellulose structure has not been changed during modification.42 A new absorption peak at around 1723 cm-1 is observed in the spectrum of mt-CNs, corresponding to the stretching vibration of ester carboxyl groups.36,38 Furthermore, the peak at 1644 cm-1 assigns to the stretching vibration of absorbed water in t-CNs, shifts to a lower wavenumber slightly. This is attributed to the partially overlap with stretching vibration of C=C bonds introduced by MAH.45 Abovementioned observations confirm that MAH has been successfully grafted onto t-CNs surface. Figure 2 (b) is the XRD patterns for t-CNs and mt-CNs, respectively. Compared with t-CNs, no obvious change is observed in mt-CNs. And both of them show 2θ angles at 14.7o, 16.7o, 22.8o, and 34.3o, which are assigned to the lattice planes (110), (110), (002) and (040) of cellulose, respectively.25,26 In addition, the degree of crystallinity (CI) can be calculated according to the following equation: CI=

I200 -Iam I200

×100%

(3)

where I200 is the intensity of the lattice plane at 22.8o, Iam represents the intensity of amorphous phase at 2θ=18o.46 According to Eq. (3), the crystallinity of t-CNs is about 84.3%, while it decreases to 73.5% for mt-CNs. This is ascribed to that hydroxyl groups on t-CNs surface is partially substituted by hydrocarbon chains, which makes it difficult for the formation of H-bonds and is unfavorable for crystallization. Therefore, CI decreased for t-CNs after modification with MAH.



Figure 2. (a) FT-IR spectra; (b) XRD patterns of t-CNs and mt-CNs. Structure characterization of the nanocomposites. The interfacial interaction between ENR and t-CNs was studied by using FT-IR first, the results are revealed in Figure 3. The absorptions at 870 and 1260 cm-1 correspond to the symmetrical and asymmetrical stretching vibration of epoxy groups, respectively.47 For the nanocomposites, the absorptions between 3450-3220 cm-1 are assigned to the -OH groups of the incorporated CNs. Figure 3 (b) shows that compared with neat ENR, the absorptions around 870 and 1260 cm-1 shift to lower wavenumbers for nanocomposite with unmodified t-CNs. This indicates the formation of H-bonds between the epoxy groups on ENR and -OH groups on t-CNs, as shown in Scheme 1 (a). However, for nanocomposite with modified t-CNs, except both the peaks shift to lower wavenumbers, the intensity of them also decreases simultaneously. Accompanied by a newly emerging peak around 1732 cm-1 (as shown in Figure 3 (a)), which is attributed to the vibration of ester groups.48 These results reveal that carboxyl groups on mt-CNs surface reacted with epoxy groups on ENR and produced β-hydroxyl esters, which could serve as interfacial covalent cross-links in ENR/mt-CNs.49 Therefore, it can be concluded that dual cross-linking networks of covalent bonds network and H-bonds network have been formed in nanocomposites with modified t-CNs (as shown in Scheme 1 (b)), which are beneficial for improving their mechanical properties.


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Figure 3. (a) FT-IR spectra of neat ENR and nanocomposites; (b) evolutions of the symmetrical (900-800 cm-1) and asymmetrical stretching vibration (1300-1200 cm-1) regions of FT-IR spectra.

Scheme 1. Illustration of the cross-linking network in nanocomposites: (a) H-bonds network in ENR/t-CNs system; (b) dual cross-linking networks: H-bonds and covalent bonds in ENR/mt-CNs system. To monitor the curing process, curing curves were recorded and the results are displayed in Figure 4 (a). For neat ENR and ENR/t-CNs nanocomposite, only slight increase in torque is observed in the whole curing process. As no curing agent was incorporated with the samples, the increase in torque is attributed to the self-curing of ENR under high temperature. It has been reported that during epoxidization reaction, some ring-opened by-products (e.g., carboxyl groups, hydroxyl groups, etc.) may form due to the high temperature, lower pH, and long epoxidization time,47,50 the generated by-products are possible to react with adjacent epoxy groups. Therefore, ENR may be self-cured, especially under high temperatures. However, for nanocomposite with mtCNs, obvious increase in torque is observed during curing process. As discussed above, only weak self-curing occurred for neat ENR, therefore, the apparent increase in torque can be ascribed to the chemical cross-linking reaction between epoxy groups on ENR and carboxyl groups on mt-CNs, as evidenced by the FT-IR results (Figure 3).51 In addition, to further support the different cross-linking networks in the nanocomposites, the hysteresis were also conducted by stretching the samples to 100% strain, and then releasing force to zero under the same speed, the results are shown in Figure 4 (b). For neat ENR, a residual strain of ~ 47% is observed due to the weak cross-linking density. The incorporation of unmodified t-CNs leads to a larger residual strain of ~ 53%, which can be explained by the formation of H-bonds network in ENR/t5 nanocomposites. During stretching, the rapture and rapid reconstruction of Hbonds take place simultaneously, the neonatal and redistributed H-bonds restrict the instantaneous network. Therefore, higher hysteresis is generated. However, for nanocomposite with modified t-CNs, an obvious decreased residual strain of ~ 34% is observed. As discussed above, both covalent cross-links and H-bonds are formed for



samples with modified t-CNs. Although H-bonds may lead to a large residual strain, the formation of covalent cross-links could provide strong recovery force, which is beneficial for the rubber chains to recover after deformation. Therefore, a much lower hysteresis is observed for ENR/m5 nanocomposite. These results are well consistent with previous analysis.39

Figure 4. (a) Curing curves; (b) hysteresis loops of neat ENR and the nanocomposites. Moreover, the cross-link densities of neat ENR and nanocomposites were also calculated based on equilibrium swelling method, as displayed in Figure 5 (a). For ENR/t-CNs nanocomposites, only slight increase is observed compared with neat ENR due to the formation of H-bonds, which is believed to serve as physical cross-link points and restricts the penetration of solvent into matrix. Compared with ENR/t-CNs, ENR/mt-CNs nanocomposites show much higher cross-link densities. For example, with 10 phr CNs, the cross-link density increases from 7.0910-5 mol/cm3 of ENR/t10 to 10.2110-5 mol/cm3 of ENR/m10, which increases by 44%. The dual cross-linking networks: chemically cross-linked rubber chains at rubber-filler interface and H-bonds bind rubber chains together tightly, which make the solvent difficult to permeate into matrix and lead to increased cross-link densities.26 The increased cross-link density may also restrict the mobility of rubber chains, which could be characterized by glass transition temperatures indirectly (Tg). Therefore, Tgs of neat ENR and nanocomposites were obtained via DSC and DMA measurements, the results are displayed in Figure 5 (b) and (c). Figure 5 (b) shows that for ENR/t-CNs nanocomposites, slight increase in Tgs is observed with increasing t-CNs content because of the formation of H-bonds. However, for nanocomposites with modified tCNs, obvious increase in Tgs is observed because of the chains movement is largely restricted in the dual cross-linking networks. The Tgs obtained from tan δ-temperature curves also show a similar tendency (Figure 5 (c)). In addition, the tan δ peak values decrease with the incorporation of unmodified t-CNs, and a further decrease is observed for nanocomposite with modified t-CNs, which indicates restricted chain mobility due to the formation of H-bonds and interfacial covalent cross-links.


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Figure 5. (a) Cross-link densities; (b) DSC curves and (c) tan δ - temperature curves for neat ENR and the nanocomposites. Morphology of ENR/CNs nanocomposites. The properties of the nanocomposites largely depend on the dispersion state of nano-fillers in rubber matrix. A uniform dispersion is always desirable. Whereas, aggregated nano-fillers in matrix may form stress concentration, which is unfavorable for the final properties. Therefore, an examination of the cryo-fractured surface was conducted by using SEM and TEM. As shown in Figure 6 (a) and (b), CNs appear as white dots, similar to previous observations.25,52 And no large aggregates of CNs are observed in ENR matrix for both the nanocomposites, indicating good interfacial adhesion between rubber and CNs. However, compared with ENR/t5 nanocomposite, a rather homogeneous dispersion is presented in ENR/m5 nanocomposite. The long hydrocarbon chains introduced via grafting reaction with MAH reduce the hydrophilic nature of mt-CNs, which inhibits the tendency to form aggregates. Therefore, mt-CNs could be distributed more uniformly in matrix. In addition, the interaction between carboxyl groups on mt-CNs surface and epoxy groups on ENR chains also promotes the homogeneous dispersion. In addition, TEM images (Figure 6 (c) and (d)) also indicate that a more uniform dispersion is achieved for ENR/m5 nanocomposite compared with ENR/t5, which is in accordance with SEM observations.



Figure 6. SEM and TEM images for: (a) and (c) ENR/t5 nanocomposite; (b) and (d) ENR/m5 nanocomposite. The tensile fractured surfaces of the nanocomposites were also characterized via SEM. As shown in Figure 7 (a), the fractured surface of ENR/t10 shows that t-CNs or the aggregates are pulled out from the matrix, which indicates the weak interfacial adhesion between them. The sole H-bonds interface is not efficient enough to transfer stress from rubber to rigid fillers, leading to the separation of t-CNs from ENR matrix. However, it can be clearly seen that the surface of ENR/m10 is rough and many mtCNs are stuck to ENR matrix, with no evidence of pulling out after tensile test. Moreover, partially orientation of mt-CNs is also observed during tensile process for ENR/m10 nanocomposite, indicating that enhanced interfacial adhesion has been achieved. The dual cross-linking networks consist of H-bonds and covalent cross-links, especially the covalent cross-links could bear large force and transfer force from rubber to fillers effectively, which leads to the orientation of mt-CNs along tensile direction and contributes to the improved strength and modulus. This is beneficial for improving tensile properties of the nanocomposites, as will be analyzed in the following section.


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Figure 7. SEM images of the cross-section of tensile fractured surfaces: (a) ENR/t10 and (b) ENR/m10 nanocomposites. Mechanical properties. To study the reinforcing efficiency of t-CNs, both the tensile tests and dynamic mechanical analysis (DMA) were performed. The results are displayed in Figure 8, and the relevant tensile properties are summarized Table 1. It can be clearly seen, for neat ENR, the stress increases gradually at first, followed by a slow increase until break, and achieves tensile strength and strain at failure of 1.29 MPa and 493%, respectively. Compared with neat ENR, with 5 phr unmodified t-CNs, tensile strength increases by 80%, while the modulus increases 1.6 times. Meanwhile, the extensibility is remained. Moreover, for nanocomposites with modified t-CNs, significantly improved strength, modulus and strain at failure are observed compared with ENR/t-CNs nanocomposites. For example, with only 3 phr mt-CNs, tensile strength and strain at failure increase to 3.07 MPa and 568%, respectively, which are even superior to nanocomposite with 5 phr unmodified t-CNs. When the content of mtCNs increased to 10 phr, ENR/m10 shows modulus and tensile strength of 3.63 MPa and 4.66 MPa, which are 3.6 and 4.4 times of neat ENR, respectively. The fracture energy also increases by ~2-fold due to a more optimal combination of elongation at break and strength. Furthermore, obvious change in stress-strain profiles is observed with the addition of reinforcing filler, in which stress increases faster in the later stage of stretching, indicating the orientation of t-CNs or mt-CNs in matrix during stretching, as will be discussed below. Generally speaking, the incorporation of reinforcing fillers in rubber matrix will lead to an improved tensile strength, accompanied by a decreased strain at failure.12,20,21 However, in this work, both tensile strength and strain at failure increase simultaneously with the incorporation of nano-fillers. This is because of the formation of dual crosslinking networks for nanocomposites with mt-CNs, namely the rubber chains are held together by covalent cross-linking network via esterification reaction and physical cross-linking network via H-bonds. During stretching, the dissociation of H-bonds occurs first because of the weak bond strength, which orientates t-CNs partially and serves as sacrificing bonds that dissipate energy to improve the extensibility of the nanocomposites.49,53,54 With further stretching, interfacial debonding occurs for ENR/tCNs nanocomposites due to the weak interfacial adhesion with only H-bonds. However, for nanocomposites with mt-CNs, both rubber chains and mt-CNs will be further oriented due to the strong interfacial covalent bonds, which is beneficial for transferring stress from rubber to filler at interface efficiently. Finally, the mt-CNs are pulled out from matrix or raptured. The orientation of mt-CNs under dual cross-linking networks contributed to the faster increase of stress in the later stage of stretching (as observed in Figure 8 (a)). The proposed fracture mechanism during stretching process of the nanocomposites is illustrated in Scheme 2.



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Figure 8. (a) Typical tensile stress-strain curves; (b) storage modulus (E’) vs. temperature of neat ENR and nanocomposites. Table 1. Mechanical properties of neat ENR and nanocomposites. Tensile strength

Strain at

Modulus

Fracture

(MPa)

failure (%)

(MPa)

energy (KJ/m3)

ENR

1.29 ±0.07

493 ±35

0.83 ±0.05

402 ±18

ENR/t5

2.31 ±0.14

512 ±40

2.14 ±0.11

715 ±43

ENR/m3

3.07 ±0.18

568 ±32

1.48 ±0.07

970 ±52

ENR/m5

3.54 ±0.17

530 ±35

2.33 ±0.14

1036 ±42

ENR/m10

4.66 ±0.18

522±29

3.63 ±0.15

1248 ±49

Sample

Figure 8 (b) displays the E’ vs. temperature curves for neat ENR and nanocomposites. As can be seen, obvious improvement of E’ in both rubbery and glassy region is achieved with incorporation of CNs into ENR matrix. For example, with the addition of 10 phr CNs, E’ of ENR/m10 at -50 oC and 20 oC increase by 168% and 12-fold compared to neat ENR, respectively. At the same filler loading, the improvement in E’ of nanocomposites with mt-CNs is more obvious than nanocomposites with t-CNs. With respect to ENR/t10, an increment of 18% and 107% of E’ at -50 oC and 20 oC are achieved for ENR/m10. The increased interfacial interaction and cross-link density due to the covalent cross-linking at interface, as characterized by FT-IR and cross-link density measurements, contributes to the improved E’ for nanocomposites with modified t-CNs.


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Scheme 2. Illustration for uniaxial stretching process of nanocomposites: (A) ENR/mtCNs, the rubber chains are bonded together by dual cross-linking networks: covalent cross-linking network and physical cross-linking network via H-bonds; (B) ENR/t-CNs, the rubber chains are held together by only physical cross-linking network via H-bonds. Conclusions In this work, we have demonstrated the formation of dual cross-linking networks in bio-based ENR nanocomposites. By introducing carboxyl groups onto t-CNs surface via grafting reaction with MAH, interfacial covalent bonds were formed between epoxy groups and carboxyl groups on modified t-CNs, which could serve as chemical crosslinking network with high strength. Meanwhile, H-bonds between hydroxyl groups on t-CNs and epoxy groups on ENR could serve as physical cross-linking network. The formation of dual cross-linking networks achieved simultaneous improved strength and extensibility. With addition of 5 phr CNs, nanocomposite with dual cross-linking networks showed 172% and 158% increment in tensile strength and toughness compared with neat ENR, and 52% and 45% increment compared with sole physically cross-linked nanocomposite. It should be noted that if the rubber chains have also been cross-linked by commonly used cross-linking agents, such as sulfur and peroxide, the reinforcing efficiency would be even stronger. References 1. Liff, S. M.; Kumar, N.; McKinley, G. H. High-performance elastomeric nanocomposites via solvent-exchange processing. Nat. Mater. 2007, 6, 76-83. 2. Ma, L. F.; Bao, R. Y.; Liu, Z. Y.; Yang, W.; Xie, B. H.; Yang, M. B.; Fu, Q. A highperformance temperature sensitive TPV/CB elastomeric composite with balanced electrical and mechanical properties via PF-induced dynamic vulcanization. J. Mater. Chem. A 2014, 2, 16989-16996. 3. Papageorgiou, D. G.; Kinloch, I. A.; Young, R. J. Graphene/elastomer nanocomposites. Carbon 2015, 95, 460-484. 4. Chen, Y. K.; Huang, X. H.; Gong, Z.; Xu, C. H.; Mou, W. J. Fabrication of High Performance Magnetic Rubber from NBR and Fe3O4 via in Situ Compatibilization with Zinc Dimethacrylate. Ind. Eng. Chem. Res. 2017, 56, 183-190. 5. Fu, L. H.; Wu, F. D.; Xu, C. H.; Cao, T. H.; Wang, R. M.; Guo, S. H. Anisotropic Shape Memory Behaviors of Polylactic Acid/Citric Acid–Bentonite Composite with a Gradient Filler Concentration in Thickness Direction. Ind. Eng. Chem. Res. 2018, 57, 6265-6274. 6. Xu, C. H.; Zhan, W.; Tang, X. Z.; Mo, F.; Fu, L. H.; Lin, B. F. Self-healing chitosan/vanillin hydrogels based on Schiff-base bond/hydrogen bond hybrid linkages. Polym. Test. 2018, 66, 155-163. 7. Habibi, Y.; Lucia, L. A.; Rojas, O. J. Cellulose Nanocrystals: Chemistry, SelfAssemble, and Applications. Chem. Rev. 2010, 110, 3479-3500. 8. Kargarzadeh, H.; Mariano, M.; Huang, J.; Lin, N.; Ahmad, I.; Dufresne, A.; Thomas, S.; Recent developments on nanocellulose reinforced polymer nanocomposites: A review. Polymer 2017, 132, 368-393.



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Dual cross-linking networks achieves simultaneous improved strength and extensibility. Both rubber matrix and cellulose nanocrystals are obtained from natural resources, which are renewable and sustainable.


Dual Cross-linked Epoxidized Natural Rubber Reinforced by Tunicate

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