Visualizing the Toughening Mechanism of Nanofiller with 3D X-ray

Sep 8, 2017 - Adding silica nanofiller in silicone rubber can toughen the matrix 3 orders in terms of fracture energy, which is far larger than most o...
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Visualizing the Toughening Mechanism of Nanofiller with 3D X‑ray Nano-CT: Stress-Induced Phase Separation of Silica Nanofiller and Silicone Polymer Double Networks Lixian Song,†,‡ Zhen Wang,† Xiaoliang Tang,† Liang Chen,*,† Pinzhang Chen,† Qingxi Yuan,§ and Liangbin Li*,† †

National Synchrotron Radiation Lab and CAS Key Laboratory of Soft Matter Chemistry, University of Science and Technology of China, Hefei 230026, People’s Republic of China ‡ State Key Laboratory Cultivation Base for Nonmetal Composites and Functional Materials, Southwest University of Science and Technology, Mianyang 621010, Sichuan, People’s Republic of China § Beijing Synchrotron Radiation Facility, Institute of High Energy Physics, Chinese Academy of Sciences, Beijing 100049, People’s Republic of China S Supporting Information *

ABSTRACT: Adding silica nanofiller in silicone rubber can toughen the matrix 3 orders in terms of fracture energy, which is far larger than most other nanofiller−rubber systems. To unveil the astonishing toughening mechanism, we employ in situ synchrotron radiation X-ray nanocomputed tomography (Nano-CT) technique with high spatial resolution (64 nm) to study the structural evolution of silica nanofiller in silicone rubber matrix at different strains. The imaging results show that silica nanofiller forms threedimensional connected network, which couples with silicone chain network to construct a double-network structure. Stress-induced phase separation between silica nanofiller and silicone polymer chain networks is observed during tensile deformation. Unexpectedly, though the spatial position and morphology of nanofiller network changes greatly at large strains, the connectivity of nanofiller network shows negligible reduction. This indicates that nanofiller network undergoes destruction and reconstruction simultaneously, during which silica nanofiller serves as reversible high functionality cross-linker. The reversible bonding between silica nanofiller and silicone rubber or between nanofiller particles can dissipate mechanical energy effectively, which may account for the 3 orders enhancement of toughness.

1. INTRODUCTION Silicone rubbers are high performance polymeric materials due to their unique molecular structure and strong interaction with reinforcing particles, which can be tailored for a variety of applications in virtually every industry.1−5 Nevertheless, the polymer network of silicone rubber is mechanically weak relative to many other elastomers even after being cross-linked. To remedy this, silica nanofiller is incorporated into the polymer for mechanical reinforcement and toughening purposes, which enhances both strength and stiffness. The mechanical enhancement of silica nanofiller on silicone rubber is astonishing, which can increase fracture energy more than 3 orders. However, the microscopic reinforcement and toughening mechanisms of silica nanofiller on silicone rubber still remain elusive.2,6−17 After intensive study for decades, it is generally agreed that nanoparticles aggregate in rubber and form three-dimensional network structure through the bounded rubber around the filler surface, as referred by Lame and Mark et al.3,5,18−30 The mechanical reinforcement and toughening essentially stem © XXXX American Chemical Society

from the synergetic effect of the nanofiller network and the matrix polymer chain network, which resemble an interpenetrated double-network structure.2,4,6,31−41 Thanks to recent active study on double-network hydrogels,42−50 the reinforcement and toughening mechanisms in those systems have achieved significant progress, which may be borrowed for understanding that in filler−rubber double-network composite. A recent review summarizes four different mechanisms for dissipating mechanical energy in hydrogels, summarized by Gong et al.,51−55 namely (i) fracture of polymer chains, (ii) reversible cross-linking of polymer chains, (iii) transformation of domains in polymer chains or cross-linkers, and (iv) fracture and pullout of fibers or fillers. As the second network in filler− rubber composite is constructed by aggregates of nanoparticles and bounded rubber, which serve as both chains and crosslinkers, the above four dissipating mechanisms may all involve Received: March 14, 2017 Revised: August 1, 2017

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DOI: 10.1021/acs.macromol.7b00539 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules in the mechanical reinforcement and toughening of filler− rubber composite. Recently, with synchrotron radiation X-ray nanocomputed tomography (Nano-CT), we have quantitatively assigned the reinforcement and toughening effect of carbon black on natural rubber to three mechanisms, i.e., elastic deformation, destruction, and friction of filler network.18 The fracture of filler network indeed plays an important role in dissipating mechanical energy while the other three mechanisms in hydrogels are not explicitly demonstrated. Compared with the toughening effect of carbon black on natural rubber where only several times enhancement is obtained,56,57 filling silica in silicone rubber can lead to 1 or 2 orders more enhancement on fracture energy.58,59 Thus, it is expected that the toughening mechanism in the silica−silicone rubber system is essentially different from that in the carbon black−natural rubber system. Aiming to unveil the toughening mechanisms of silica nanofiller on silicone rubber, we employ synchrotron radiation X-ray Nano-CT technique to study the structural evolution of silica-filled silicon rubber under different strains. Based on the three-dimensional (3D) structural information on silica nanofiller through 3D image processing, stress-induced phase separation between silica nanofiller and silicone polymer chain networks was observed during tensile deformation, which undergoes a structural evolution involving fracture, migration, and reversible cross-linking of nanofiller aggregates and network. The stress-induced phase separation supports that the four mechanisms for dissipating mechanical energy all participate in the silica−silicone rubber system, which may account for its astonishing enhancement on fracture energy.

(BSRF). An elliptically shaped capillary condenser was used to focus the incident X-ray onto the sample, and a phaser was placed in the back focal plane of the objective to provide the Zernike phase contrast. The images were recorded by a charge coupled device camera (1024 × 1024 pixels). In our experiment, a layout with a large field of view was used: the field of view was 60 μm × 60 μm with a spatial resolution of 64 nm, and the photon energy was 8 keV.60 In experiment, the rectangular-shaped rubber sample was mounted between two chucks of a homemade miniature tensile device which could draw thin films with different ratios. The initial sample length or distance between two chucks was 2 mm. Five different preset strains of 0, 0.5, 2.0, 3.0, and 4.0 were imposed on samples for X-ray 3D imaging measurements with a low stretching speed of 20 μm/s. Note that the deformed sample was kept for 1 h for relaxation after it reached the preset strain, which ensured a motionless situation throughout the sample. Before placing the loading sample inside the TXM, some gold particles were carefully pasted on sample surface under an optical microscope, which were used for calibrating the serial tilts and the tomographic reconstruction. During experiments, the sample was fixed on a rotation stage, and each exposure time of X-ray imaging was 15 s. A series of two-dimensional (2D) projections were acquired at title angles ranging from −70° to +70° with an interval of 0.5° by the rotation of sample. Approximate 281 serial 2D projections were obtained for each strain, which were further reconstructed into 3D tomograms with 300 × 300 × 350 cubic voxels (voxel size of 64 nm) through a commercial software package provided by Xradia XMReconstructor software.61 2.4. 3D Visualization and Quantitative Analysis. 3D visualization and quantitative analysis of the tomograms were carried out using Avizo Fire VSG software (Visualization Sciences Group, Bordeaux, France).62 First of all, both the 3D median filtering and the smooth filtering were applied to reduce the noise of image. Then the filtered image was processed to find different phases in sample by the segmentation operation. Since the silica aggregates have higher mass density (∼2.2 g/cm3) to cause more X-ray absorption than the silicon rubber matrix (∼1.0 g/cm3) during imaging, there exists gray value difference (imaging contrast) between these two components. By choosing an appropriate threshold of gray value (or pixel intensity) to separate the two components, the spatial distribution of silica aggregates in sample can be extracted.63,64 Note that the threshold chosen is very important for the extraction of silica aggregates from 3D images, which was modified and determined until the calculated volume fraction of silica filler from 3D image was very close to the known actual value (∼15.4%) in sample.65,66 Figure 1a shows a representative 3D imaging of silica aggregates in silicon rubber matrix at strain of 0, where the rubber matrix is made transparent. In large scale, silica aggregates form a three-dimensional network, which undergoes deformation, destruction, and reconstruction during deformation. On the basis of the X-ray imaging, we constructed the

2. EXPERIMENTS 2.1. Materials. Methyl vinyl silicone rubber (vinyl, 0.23%) with viscosity average molecular weight of about 650 000 g/mol was purchased from Zhejiang Wynca Chemical Industrial Group Co., Ltd. Precipitated silica nanoparticles (Evonik-Degussa AG, Germany) were chosen as the nanofillers in silicone rubber with a specific surface area of approximately 140 m2/g (measured by N2 adsorption) and a primary diameter of about 20 nm.20 Hydroxyl silicone oil (GY-209-3) was used as structure controlling and coupling agent, which was kindly provided by Chenguang Research Institute of Chemical Industry, China. Dicumyl peroxide (DCP, 99%, AR) used as a vulcanization agent was purchased from Sinopharm Chemical Reagent Co., Ltd., China. The final rubber compound was composed of methyl vinyl silicon rubber, silica nanoparticles, hydroxyl silicone oil, and dicumyl peroxide with a weight ratio of 100:40:4:3. 2.2. Sample Preparation. The silica-filled rubber compound was prepared by two-step mixing method in an internal mixer (HAAKE, PolyLab OS-RheoDrive 7) with a 75 cm3 mixing chamber. First, the methyl vinyl silicone rubber master batch (50 g), the silica nanoparticles (20 g), and the hydroxyl silicone oil (2 g) were mixed together under a mixing speed of 90 rpm for 30 min at 105 °C. Subsequently, the compound was cooled down to room temperature (25 °C) and kept for 2 weeks to ensure an equilibrium adsorption/ desorption. Second, the vulcanization agent DCP (1.5 g) was added into the compound under a mixing speed of 60 rpm for 15 min. The final rubber compound was cross-linked and molded into silicone rubber film in a compression molding machine (P300E, DR COLLIN Co., Ltd., Germany) under 20 MPa for 10 min at 160 °C. The thickness of the obtained rubber film was 50 μm, which matches with the focal depth of X-ray imaging system. The film was then cut into rectangular-shaped specimens with length and width of 10 and 2 mm, respectively, for uniaxial tensile testing and X-ray imaging. 2.3. X-ray 3D Imaging. The Nano-CT experiments were performed using the transmission hard X-ray microscope (TXM) on the 4W1A beamline station of Beijing Synchrotron Radiation Facility

Figure 1. 3D reconstruction of silica aggregates in silicon rubber matrix. (a) 3D visualization of sample by volume rendering (top) and corresponding schematic drawing (bottom). (b) Skeletonization model of silica aggregates (top) and corresponding schematic drawing (bottom). B

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Macromolecules pseudo-silica networks along the backbone of silica aggregates with the use of Avizo Fire Extension Skeleton command, as shown in Figure 1b. In order to quantitatively describe the enhancement mechanism of silica-filled silicon rubber, some crucial parameters such as the coordinate, size and number of silica aggregates, the coordinate and number of filler network cross-linked point, and the end coordinate of filler network were derived from 3D images at different strains. Then the average filler network mesh size, the connectivity of filler network, and the cross-linked point density of filler network were calculated using self-developed programs.

3. RESULTS 3.1. Tensile and Mechanical Fracture Measurements. Before presenting the X-ray Nano-CT results on silica filler network, we first compare the mechanical properties of filled and unfilled silicone rubbers. Figure 2 shows the stress−strain Figure 3. Constructed filler networks along the backbone of silica aggregates from X-ray three-dimensional reconstruction imaging at different strains of (a) 0, (b) 0.5, (c) 2.0, (d) 3.0, and (e) 4.0. Also shown is the corresponding engineering stress−strain curve. The tensile direction is vertical.

3b−e). At strain of 2.0, the density of filler network seems modulating along tensile axis (Figure 3c), while filler network assembles into bundle-like morphology aligning in tensile direction at strain of 3.0 (Figure 3d). Further increasing strain to 4.0 (Figure 3e), the bundle morphology is partially fractured and the separated filler network is remixed again. In order to quantitatively analyze the deformation, destructing and reconstructing of filler network under different strains, the structure parameters of filler network were derived from 3D image processing and computer program calculation. The average network distance or mesh size l was calculated by l = (V/N)1/3 with an assumption of the cubic occupation for a single network cross-linked point, where V is the volume of sample and N is the total number of cross-linked points. The average network distance is plotted as a function of strain in Figure 4a (open circle-black line). The variation of average network distance with tensile strain shows a nonmonotonicity. Before tensile deformation (at 0 strain), the average network distance is 622 nm, which increases up to 671 and 721 nm at strains of 0.5 and 2.0, respectively. With further increasing strain, the average network distance shows a decrease, which reduces to 703 and 629 nm at strains of 3.0 and 4.0, respectively. The magnitude of average network distance at strain of 4.0 is nearly the same as that at 0 strain. In addition to the average network distance, filler network connectivity (Tnet) is an important parameter to explore destructing and reconstructing of filler network structures under deformation.18 The silica aggregates dispersed in sample are fractal entities composed of a combination of primary silica nanoparticles, which are assumed to be spherical for the convenience of quantitative calculation. For a given silica aggregate, its equivalent diameter di is expressed as di = (6V3D/ π)1/3 with V3D being its real volume. Then the average diameter dAve of silica aggregates can be calculated by dAve = ∑Nidi/Nξ, where Ni is the number of aggregates with diameter of di and Nξ is the total number of aggregates in sample. According to the random arrangement,67 the average interparticle surface-tosurface distance between adjacent silica aggregates is dc = (0.86ϕ−1/3 − 1)dAve, where ϕ is the silica volume fraction. A cutoff factor β = di/dc is defined to describe the filler network connectivity. If β ≥ 1 (i.e., aggregates with di larger than dc), the

Figure 2. Engineering stress−strain curves of unfilled and filled silicon rubber samples. The insertion shows an enlargement of curves at the beginning small strain.

curves of both filled and unfilled silicone rubber samples during tensile deformation. The fracture strain and stress of unfilled silicone rubber are about 0.3 and 0.2 MPa, respectively, while adding silica nanoparticles renders a fracture strain and stress of about 4.5 and 10.1 MPa, respectively, indicating tremendous reinforcement and toughening effects of filler. For a quantitative comparison, we calculated the fracture energy Γ of filled and unfilled silicon rubbers, which is the work required to generate a crack on unit area of sample and expressed as51−54 Γ = L0

∫0

ε

σ (ε) d ε

(1)

where σ(ε) is the tensile stress at strain of ε and L0 is the initial length of sample (2 mm). Then the toughening factor can be defined as

ΔΓ = Γfilled/Γunfilled

(2)

where Γfilled and Γunfilled are fracture energies of filled and unfilled silicon rubbers, respectively. Based on eqs 1 and 2, Γunfilled, Γfilled, and ΔΓ are calculated to be 66 J/m2, 49 400 J/ m2, and 749, respectively. In other words, the fracture energy of filled silicone rubber is 749 times larger than that of the unfilled one. 3.2. Visualizing Structural Evolutions of Filler Networks during Deformation. Figure 3 shows the 3D images of filler network along the backbone of silica aggregates at different strains. For a better view of the 3D structure, corresponding movies are provided in the Supporting Information (Movies 1a−e). Before tensile deformation at zero strain, the highly connected filler network distributes homogeneously in silicone rubber matrix, as shown in Figure 3a. Imposing strain results in deformation of filler network, which gradually becomes loose and inhomogeneous (Figures C

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Figure 4. Structure parameters of filler networks at different strains and the corresponding mechanical property. (a) Average distance and connectivity of filler network at different strains. The corresponding error bars are plotted with considering the possible distribution of silica aggregates from the loosest to the closest arrangement in rubber matrix.67 (b) Tensile modulus as a function of filler network density. The solid line shows a linear fitting of experimental data. The values in yellow box correspond to different strains.

increase of filler network density implies that more chain segment units cooperate to transfer stress under deformation, which elevates the tensile modulus of sample.68 At strain of 4.0, which is close to the fracture strain of 4.5, the overstretching of sample is expected to happen. In this case, the excessive stress results in fracture of bundles or phase morphology of filler network (see Figures 3e and 5e), leading

filler network is considered to be connected. Then the percentage of network connectivity is expressed as Tnet = 100Ni(β ≥ 1)/Nξ, where Ni(β ≥ 1) is the number of silica aggregates with di larger than dc. Tnet at different strains are calculated and plotted in Figure 4a (open square-blue line). With the increase of strain, Tnet decreases slightly from 84% before deformation to 81% at strain of 3.0, which subsequently reverses the trend and increases back to nearly the same as the original value at 0 strain. Comparing to the carbon black filler in natural rubber,18 the reduction of network connectivity Tnet of silica filler in silicone rubber is negligibly small even at strain of 3.0. In the carbon black−natural rubber system, the reduction of Tnet is about 50% under the same strain. Note the negligible reduction of connectivity Tnet of silica filler network in silicone rubber matrix does not mean that fracture of filler network does not occur during tensile deformation. Considering large strain up to 4.0, fracture of filler network should inevitably take place, while nearly unchanged network connectivity stems from reversible bonding of silica aggregates. In other words, debonding and rebonding occur simultaneously in silica filler network, which is the most important contributor for its astonishing toughening effect as will be discussed later. To investigate the contribution of filler network to the macroscopic mechanical properties of rubber composite, Figure 4b presents the tensile modulus E as a function of filler network density n. The tensile modulus E at different strains is calculated by E = δσ/δε based on the stress−strain curve in Figure 2. The filler network density n is obtained by n = 1/l3, where l is the average network distance as shown in Figure 4a. Interestingly, except at the largest strain of 4.0 near macroscopic fracture of the composite, the tensile modulus E is linearly proportional to the filler network density n although the tensile strains are different. For the unfilled silicone rubber, the fracture strain is only about 0.3, which is far below that of general network polymers, indicating that the chain segment in silicone rubber is unable to bear large tensile strain before fracture. We speculate that the same situation also exists in the silica-filled silicone rubber. Although the apparent tensile strain increases from 0 to 4.0, the real microstrain of single chain segment is suggested to always keep at a low level (≤0.3) during the whole deformation process. Thus, the observed linear relationship between E and n at different strains (ε ≤ 3.0) indicates that the filler network serves as the main contributor on the mechanical behavior at strain larger than the fracture strain (0.3) of the unfilled silicone rubber. In this case, the filler network connectivity Tnet reflects the real mechanical connection, where whole filler network bears stress. The

Figure 5. Constructed filler networks along the backbone of silica aggregates at different strains, which correspond to the thickness of 2 μm (30 pixels) of samples. (a−e) Views of the x−z plane and (a′−e′) tilting 45° to the x−z plane. The tensile axis is in the z direction.

to a lower sample modulus than the prediction, i.e., deviation from the linear relationship in Figure 4b. However, the fracture of bundles occurs mainly on a large scale, which does not reduce the filler network density. Instead, the filler network density shows an increase at strain of 4.0 relative to that at strain of 3.0. The possible reason is that the silica aggregates originally restricted at strain of 3.0 are partially released to be free after the break of bundles at strain of 4.0, where the rebonding or remixing among them occurs. On the basis of the correlation between network structure and mechanical properties, we postulate that fracture and fully mechanical reconnection of filler network can occur simultaneously at strain below 3.0, while under larger strain fully mechanical reconnection of filler network is unattainable anymore, which may eventually result in the final fracture of the rubber composite. To demonstrate the details of structural evolution of the filler network during deformation, the enlarged 3D images of 300 × 30 × 350 pixel3 (x × y × z) at different strains are extracted, respectively. Figures 5a−e and 5a′−e′ are the perspectives at different strains, which are parallel to the x−z plane and at an angle of 45° to the x−z plane, respectively. The corresponding 3D movies are provided in the Supporting Information (Movies 2a−e). Similar to Figure 3, before deformation the filler network is finely and homogeneously dispersed in the rubber matrix (Figures 5a,a′). At strain of 0.5, the mesh size of filler D

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which stems from stretched filler network (Figure 6a, ε = 0.5− 4.0). Weak scattering maxima appear in the meridianal direction at strains of 2.0 and 3.0, indicating that concentration fluctuation or phase separation with the anisotropy direction parallel to the tensile direction occurs between filler network and polymer matrix.73 At strain of 3.0, a streak-like scattering signal appears close to the center in equator, which comes from the aligned bundles and is similar to the string phase in flowinduced phase separation of polymer blends.70 Further increasing strain to 4.0, the streak-like signal becomes weak and the isotropic scattering patterns partially restore, which resembles the flow-induced remixing process.74 Figure 6b depicts 1D intensity profiles integrated in the meridianal direction, in which the scattering peak is observed at strains larger than 0.5. The scattering intensity increases with strain in the observed Q range, while the Q value at peak position decreases from 0.75 × 10−3 A−1 at strain of 2.0 to 0.5 × 10−3 A−1 at strain of 4.0, indicating that increasing strain results in an increase of the average wavelength of concentration fluctuation of silica aggregates from about 830 to 1200 nm along the tensile direction. On the other hand, no obvious scattering peak is observed in 1D intensity profiles in the equatorial direction (Figure 6c), although about 3000 nm distance between filler bundles can be found perpendicular to the tensile direction in Figures 3d and 5d′. This may be attributed to a few repetitions of bundles in the probed zone of sample, the poor periodicity of bundle morphology, and the existence of background noise at low Q range. The evolution trend of phase separation between the silica aggregates and the silicone rubber matrix observed in this work is similar to the stress-induced phase separation process in polymer blends as reported by Hashimoto et al.69,70

network is obviously increased (Figures 5b,b′), while increasing strain to 2.0 filler network chains are aligned preferentially perpendicular to tensile axis (Figures 5c,c′). The most apparent change in the morphology occurs at strain of 3.0 (Figures 5d,d′), where bundle-like regions with high concentration of filler chains are generated. These bundles parallel to stretch direction and have diameter ranging from several hundred to several thousand nanometers, which are separated by regions with low concentration of filler chains. The average distance among the bundles is about 3000 nm. Further increasing strain to 4.0 results in destruction of such filler bundles (Figures 5e,e′). Meanwhile, the bundle morphology remains partially but becomes thicker aligning along the tensile axis. Figure 5 shows the structural evolution of filler network from a homogeneous phase to one with an oriented bundle morphology, which closely resembles stress-induced phase separation of polymer blends.69,70 Judging from the direct morphological observation by 3D X-ray Nano-CT imaging, we tentatively take this process as the phase separation between filler and matrix rubber networks, since applying stress field has induced the appearance of filler network rich zones. The observed separation between filler and matrix rubber networks happens periodically throughout the whole sample, similar to that of microphase separation. Note the “phase separation” used here does not mean a strict thermodynamic phase transition like that in polymer blends, which depicts the morphologic reorganization of filler network. Nevertheless, as the mechanism of driving filler particles to reorganize is rather similar to the squeezing effect of stress in polymer blends, “phase separation” in the current system does have its nonequilibrium thermodynamic origin. To further verify the speculation of stress-induced phase separation between the filler and rubber chains, fast Fourier transform (FFT) analyses and corresponding 1D intensity profiles were performed on the silica filler images at different strains,71 which are shown in Figure 6. The FFT patterns are similar to small-angle light or neutron scattering patterns obtained on flow-induced phase separation in polymer blends, as referred by Moses et al.72 Before deformation, the scattering pattern is isotropic, indicating a homogeneous distribution of filler network (Figure 6a, ε = 0). After strain exceeds 0.5, the isotropic scattering pattern transforms into elliptical and lozenge shapes gradually,

4. DISCUSSION The above X-ray Nano-CT and tensile measurements clearly demonstrate that the 3D nanofiller network plays a dominant role in the silica reinforcement of the rubber matrix. The fracture stress and strain of unfilled rubber are only 0.2 MPa and 0.3, which increases to 10.1 MPa and 4.5 after adding silica filler, respectively. Consequently, the fracture energy increases from 66 to 49 400 J/m2. In other words, the fracture energy of filled silicone rubber is 749 times larger than of the unfilled one. Meanwhile, a good linear relationship meets between tensile modulus E and nanofiller network density n as shown in Figure 4b, indicating nanofiller network indeed plays a major role on the enhanced mechanical performance of rubber composite. How does silica nanofiller network toughen silicone rubber with the increase of fracture energy near 3 orders? What are the mechanisms to dissipate mechanical energy in silica−silicone rubber composite during tensile deformation? As revealed by the above synchrotron radiation X-ray NanoCT results, stress-induced phase separation between silica nanofiller network and silicone polymer chain network is responsible for enhancing the toughness of rubber composite, during which all of the four mechanisms for dissipating mechanical energies in tough hydrogels, summarized by Zhao et al.,51 involve in. As shown in Figure 5d, to reach the phase separated bundle morphology at strain of 3.0, silica nanofillers are required redistribution in the rubber matrix inevitably, during which fracture and migration of nanofiller network chains and aggregates have to take place. Note as nanofiller network chains are connected by physical interactions among nanoparticles and bounded rubbers, here fracture may occur between nanofiller particles as well as between nanofiller and

Figure 6. (a) Fast Fourier transform (FFT) images of silica aggregates dispersed in rubber matrix at different strains. (b) 1D intensity profiles along the meridianal direction as a function of strain. (c) 1D intensity profiles along the equatorial direction as a function of strain. Q is the scattering vector as defined in the scattering experiments. E

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rebonding does take place in long time after releasing strain. Correspondingly, without simultaneously reversible bonding the toughening effect of carbon black fillers on natural rubber normally results in less than 10 times enhancement on fracture energy,19,75−79 while the presence of reversible bonding in silica−silicone rubber composite leads to 1−2 orders stronger toughening effect than that in carbon black−natural rubber system. The difference of toughening effects in silica−silicone and carbon black−natural rubber composites is similar to that in double-network hydrogels with and without reversible bonding during deformation, confirming reversible bonding is responsible to the astonishing toughening effect in silica− silicone rubber composite. The reversible bonding not only determines the toughening factors but also controls the macroscopic fracture of silica− silicone rubber composite. As strain reaches 4.0, close to the macroscopic fracture strain of 4.5, the cross-linked density, connectivity, and nmorphology of nanofiller network return back nearly to those at the undeformed state. However, as shown in Figure 4b, the modulus at strain 4.0 deviates largely from the linear correlation with cross-linked density of nanofiller network at lower strains. This indicates that the 3D filler network at strain 4.0 is only morphologically but not fully mechanically connected. Here rebonding can only partially succeed. In other words, interactions in nanofiller network and between nanofiller−rubber are gradually destroyed at large strain, which eventually leads to macroscopic fracture of the composite.

rubber. Migration of silica nanofiller can be considered as squeezing of the matrix rubber, while stress-induced phase separation results in transformation of network domains from homogeneous distribution to concentrated bundles of nanofiller chains. Thus, three toughening mechanisms summarized by Zhao et al.,51 namely fracture of filler network, transformation of domains, and pullout of fillers, indeed occur in silica−silicone rubber composite to dissipating mechanical energy. The fourth mechanism, reversible cross-linking or bonding, does also happen in silica−silicone rubber composite and will be discussed later, which may be the most peculiar one comparing with other filler−rubber systems like carbon black− nature rubber composites. Though the morphology of silica nanofiller network changes dramatically at strain of 3.0, the network connectivity only shows a negligible reduction. This suggests that debonding and rebonding of silica nanofiller network occur simultaneously. In other words, silica nanofiller−nanofiller and silica nanofiller− silicone rubber interactions are reversible under deformation. More specifically, silica nanoparticles are reversible highfunctionality cross-linkers. Note the cross-linking function not only works on building nanofiller chain network itself but also plays a critical role in connecting polymer chain network and sustaining the integrity of silicone rubber matrix. As shown in Figure 2, the fracture strain of unfilled silicone rubber is only 0.3, while the filled one reaches 4.5, which is 15 times that of the unfilled one. Taking the simple affine assumption, the microstrain of silicone rubber in the filled composite must be less than 0.3 before the sample is macroscopically fractured at strain of 4.5. Two possible mechanisms may realize this requirement, namely (i) squeezing solvent out of the polymer network or phase separation and (ii) reversibly cross-linking two networks. Squeezing solvent out of the polymer network can be an important factor in tough hydrogel, which may follow the mechanism of elastic phase separation to dissipate mechanical energy. As discussed early, stress-induced phase separation is the essential mechanism for toughening in silica− silicone rubber composite, which can help to release microstrain of polymer chain network partially. Microscopically, for dissipating mechanical energy in silica−silicone rubber composite, reversibly cross-linking of networks plays a more important role than the simply solvent-squeezing effect. As silica nanoparticles are not free small molecular solvent but present as physically connected nanoparticle network, redistributing silica nanofiller network must undergo fracture and rebonding between nanofiller and nanofiller as well as nanofiller and rubber matrix. Upon deforming, these reversible physical bonds not only dissipate mechanical energy but also guarantee the integrity of both nanofiller chain network and polymer chain network as well as stress transmission between them, which prevents rubber composite from losing mechanical stability before macroscopic fracture. The reversible high-functionality silica nano-cross-linkers or reversible filler−filler and filler−rubber bonds are the essential factor for toughening silica−silicone rubber composite with 3 orders increase of fracture energy. As shown in our early work,18 deformation with strain of 3.0 induces significantly reduction of network connectivity (about 50%) in carbon black−natural rubber composite, while only 3% reduction of network connectivity is generated under the same strain in silica−silicone rubber composite. This indicates that the reversible bonding seems not occurring in carbon black− natural rubber composite under deformation, although

5. CONCLUSIONS The silica filler−network structures and their evolution in silicon rubber matrix under loading are studied in situ using synchrotron-radiation X-ray Nano-CT and tensile measurements. The results show that adding silica nanofiller increases the fracture energy to 749 times larger than the unfilled silicone rubber. Moreover, a good linear relationship between tensile modulus E and silica nanofiller network density n is obtained at strain up to 3.0, indicating nanofiller network plays a critical role in the enhanced mechanical performance of rubber composite. Stress-induced phase separation between silica nanofiller network and silicone polymer chain network was observed directly through X-ray Nano-CT during tensile deformation, during which nanofiller network undergoes debonding and rebonging simultaneously and negligible reduction of network connectivity occurs even at large strains up to 3.0. Compared with the evolution of network morphology and toughening effect in other filler−rubber composites like the carbon black−natural rubber system, toughening silicone rubber with 3 orders of increase of fracture energy by silica nanofiller can be attributed to reversible bonding between nanoparticles as well as between nanofiller and rubber, which is in line with general toughening mechanisms as observed in double-network hydrogel. Our findings identify the key structural features of the highly deformed rubber composite and provide new insights into the evolution of nanofiller network during deformation, which may help to design high performance elastomer composite.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00539. F

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Movies of structural evolution about filler network at different strains (ZIP)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (L.L.). *E-mail [email protected] (L.C.). ORCID

Liangbin Li: 0000-0002-1887-9856 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by National Natural Science Foundation of China (51325301, 51573177, 51473151, and U1530114). The synchrotron beam time from Beijing Synchrotron Radiation Facility (BSRF) and National Synchrotron Radiation Lab (NSRL) is acknowledged.

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