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Anisotropic Nonlinear Mechanical Behavior in Carbon Nanotubes/ Poly(1,4-cis-isoprene) Nanocomposites Silvia Agnelli,*,† Stefano Pandini,† Andrea Serafini,‡ Sara Musto,‡ and Maurizio Galimberti‡ †

Department of Mechanical and Industrial Engineering of Brescia, University of Brescia, Via Branze 38, 25123 Brescia, Italy Department of Chemistry, Materials and Chemical Engineering G. Natta, Politecnico di Milano, via Mancinelli 7, 20131 Milano, Italy

‡

ABSTRACT: This work shows the remarkable anisotropic nonlinear mechanical behavior of nanocomposites based on carbon nanotubes and poly(1,4-cis-isoprene), prepared by mechanical mixing, calendering, and compression molding. The dynamicmechanical moduli were measured as a function of the applied strain and along specific directions. The specimens revealed, at all filler contents, an orthotropic and transversally isotropic response, which means properties very similar inside the sheet plane and very different from those in the orthogonal direction. This response was correlated to the material structure by means of bright field transmission electron microscopy analysis, coupled with electron diffraction measurements, so to observe the nanofiller structuring through the specimen thickness. Preferential orientation of nanotubes and alternate areas containing large or low CNT amount were revealed. This work shows that energy dissipation is not isotropic in CNT-filled polymer nanocomposites and aims at giving a contribution for controlling such an important phenomenon.

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filler, the pronounced Payne effect, much larger than for composites with the same volume fraction of carbon black, is usually explained with the presence of CNT agglomerates and with the weak interfacial interaction between the filler and the matrix, studied through spectroscopic analysis, mainly through Raman spectroscopy.18,31,32 Whatever is the mechanism at its origin, the Payne effect has a dramatic influence on the composites properties. In fact, it is widely acknowledged that the larger Payne effect means larger dissipation of energy in dynamic-mechanical applications of elastomeric materials. Huge research activity is performed in order to reduce the Payne effect of the composites, particularly with silica and carbon black as the fillers.16 As a consequence of the high aspect ratio of CNT, their orientation along specific directions in a polymeric matrix is a fundamental aspect for understanding their reinforcement mechanisms, since it may lead to anisotropic properties of the polymer nanocomposites. In the case of an aligned CNT structure their reinforcing potential can be thus exploited along

INTRODUCTION In the polymer field, a hot research topic is polymer nanocomposites based on nanosized1,2 carbon allotropes, such as carbon nanotubes (CNT), single walled3,4 and multiwalled,5,6 graphene,7−10 and graphites made by few layers of graphene, also named graphite nanoplatelets.11−14 A carbon nanotube is characterized by very high aspect ratio (i.e., the tube length prevails over the diameter) and large surface area.15 This allows to have large CNT−polymer interfacial area and a higher reinforcing efficiency with respect to carbon black or silica, whose primary particles are fused together into nearly spherical aggregates.16 Because of their structure, CNT bring about substantial improvement of the material initial modulus of an elastomerbased composite.2,17−22 In the case of polymer melts and elastomers,17,18,21,22 it has also been shown the pronounced nonlinearity of the modulus: after a plateau at small strains, the storage modulus strongly decreases by increasing the strain amplitude. Such nonlinearity of the modulus, ascribed to the filler networking phenomenon, is known as the Payne effect16,23,24 and is traditionally interpreted with agglomeration and deagglomeration of filler particles, above the filler percolation threshold,25,26 and/or with filler−matrix bonding and debonding mechanism.27−30 In the case of CNT as the © 2016 American Chemical Society

Received: August 3, 2016 Revised: October 20, 2016 Published: November 3, 2016 8686

DOI: 10.1021/acs.macromol.6b01682 Macromolecules 2016, 49, 8686−8696

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mechanical response at various strain levels, along all the directions needed to fully describe an anisotropic behavior. Preferential CNT orientation was also investigated by bright field transmission electron microscopy (BF-TEM). Recently, it has been demonstrated that electron diffraction pattern can be used as a texture analysis tool in the case of nanostructured materials.38,39 In this work, for the first time in the literature on polymer nanocomposites, electron diffraction pattern has been combined with conventional TEM imaging and dynamic mechanical measurements to evaluate CNT orientation in a polymer matrix and the orientation effect on polymer properties.

a preferential direction, and the properties of CNT-based polymer nanocomposites can be specifically enhanced. The promotion of highly aligned CNT structures in an elastomeric matrix and the derived effects are topics investigated in the literature.18,31,33−37 A high level of CNT alignment has been obtained with several techniques,31,36 such as electrospinning, plasma-enhanced chemical vapor deposition (PECVD), filtration, magnetic field-induced alignment, liquid crystalline phase-induced alignment, and mechanical stretching, but also with common industrial production processes,31,33−35 such as extrusion or injection molding, thanks to the shear deformation occurring in polymer melt flow. CNT orientation may be also altered by the application of strain: in a styrene− butadiene rubber, CNT orientation was observed to decrease after a second stress−strain, and this finding was commented as an indication of weak CNT−polymer interaction.18 CNT orientation can influence electrical,31,35−37 thermal,34 and mechanical34,35 properties of elastomer-based composites. Literature works investigating the anisotropic properties induced by CNT are mostly focused on the increase of properties along the alignment direction of CNT, whereas a proper description of the properties along other directions has mostly been neglected. As an exception, in the work of Kueseng et al.35 CNT alignment in elastomeric films (natural rubber/ nitrile rubber = 50/50) was controlled by milling parameters, and anisotropic mechanical and electrical properties were studied in different directions. Further studies on the effect of CNT anisotropy on the properties of polymer nanocomposites appear thus of great interest for the large scale development of CNT-based materials. The aim of the present work was to investigate in a CNT-reinforced elastomeric matrix the dependence on CNT orientation of the nonlinear mechanical response. More specifically, the dynamic mechanical properties and their dependence on strain amplitude (i.e., Payne effect) were studied along different directions and not only along CNT alignment direction, in systems in which a significant degree of order of the filler was achieved. An elastomeric matrix was used, based on natural rubber (NR), poly(1,4-cis-isoprene) from Hevea Brasiliensis. NR is known for its outstanding elastic properties. To promote the mechanical reinforcement of NR with little amount of a reinforcing filler, obtaining a reinforced and light material would be highly desirable. CNT offer this opportunity. At the same time, it would be highly desirable to maintain the outstanding elasticity of NR, thus without modifying, at least to some extent, the balance between storage and loss part of the dynamic-mechanical modulus. In fact, as mentioned above, CNT are also known to promote remarkable enhancement of the loss modulus. The effect of CNT on dynamic-mechanical properties of NR is not yet fully understood and deserves to be investigated. Nanocomposites were prepared, through mechanical mixing, with various amounts of CNT, both in the absence and in the presence of continuous filler network. Before curing with a peroxide, nanocomposites experienced calendering and compression moldingprocessing techniques commonly used in both research and industrial laboratories. Usually, measurements are done in the plane of the plate, whereas in the present work they were also performed in the direction perpendicular to the plate, with the specific aim to investigate the material anisotropy. The nonlinear mechanical behavior of the cross-linked CNTreinforced NR was investigated through its shear dynamic

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EXPERIMENTAL PART

Materials. Poly(1,4-cis-isoprene) from Hevea Brasiliensis (natural rubber, NR) was SMR GP, with 65 Mooney units as Mooney viscosity (ML(1 + 4)100 °C), from Lee Rubber. Multiwall carbon nanotubes (CNT) were Baytubes C150 P from Bayer Material Science. Characteristics of CNT reported on the technical bulletin were chemical purity ≥95 wt %, length of 1−10 μm, number of walls of 3−15, outer and inner diameters of 10−16 and 4 nm, respectively. CNT were handled in order to avoid their direct contact with the operator. They were kept and weighed in a sealed container, a glovebox. The operator had gloves and a mask. The internal mixer used for composites preparation had a steel funnel on top of the mixing chamber, and ingredients were fed through such a funnel. The mixing chamber was not open to the air during mixing. 2,5-Dimethyl-2,5-di(tert-butylperoxy)hexane (DCUP) supported on silica/CaCO3 (45 mass %) was from Arkema Inc. Composition and Preparation of Nanocomposites. NR-based nanocomposites with different CNT contents were prepared by incorporating the filler in a matrix with 100 php (parts per hundred parts of polymer) of NR and 3.5 php of DCUP. Peroxide was used as the cross-linking agent in order to obtain a composite substantially based only on NR and CNT, thus avoiding the addition of several ingredients that would be required by a sulfur-based cross-linking system. CNT contents in php for the various composites (denomination in parentheses) were 0 (NR), 4 (CNT-4), 15 (CNT-15), and 35 (CNT-35). The compositions in volume fractions are reported in Table 1.

Table 1. Composition in Volume Fraction of the NR-Based Composites NR CNT DCUP

NR

CNT-4

CNT-15

CNT-35

0.98 0.00 0.02

0.96 0.02 0.02

0.91 0.07 0.02

0.83 0.15 0.02

Composites were prepared using a Brabender type internal mixer, with 50 mL mixing chamber. 50 g of NR was introduced into the internal mixer and masticated at 80 °C for 1 min with rotors rotating at 60 rpm. The filler was then added, and mixing was performed for further 4 min. Peroxide was added, and the final composite was discharged after 2 min. In consideration of the half-life of DCUP (10 h at 115 °C) and of the short mixing time, premature cross-linking was considered unlikely. Indeed, experimental evidence of premature scorch was not observed. Composites were finally further homogeneized by passing them five times through a two roll mill operating at 50 °C, with the front roll rotating at 30 rpm and the back roll rotating at 38 rpm and 1 cm as the nip between the rolls. The sheet of the compounds was rotated at every step. Dynamic-Mechanical Tests. Samples Preparation. Square sheets of cured rubber, with 10 × 10 × 3 mm3 as the dimensions, were obtained by a compression molding machine, operating for 10 min at 170 °C with 3.5 MPa pressure. Rectangular specimens were cut from the rubber sheets and had the following nominal dimensions: 8687

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Macromolecules height (h) = 6 mm, width (w) = thickness (t) = 3 mm (see Figure 1A). The actual dimensions of each specimen in its unstrained state were measured before testing by a traveling optical microscope.

to the maximum strain level that may be applied, according to the machine bearing capabilities (maximum load: 18 N). Therefore, the actual maximum strain value experimentally achieved was dependent on the specimen stiffness, geometry, and size. The measurement of dynamic properties was carried out under a shear sandwich configuration, shown in Figure 1B. In such configuration, the two outer clamps were fixed and the inner was moving. Two nominally identical parallelepiped rubber specimens were compressed between the moving and the fixed clamps, and a compressive strain of about 15% was employed in order to maintain the specimen surfaces fixed to the clamp faces by friction forces. The storage shear modulus was recorded as a function of strain amplitude. Before each test, the specimen was mechanically conditioned by means of a first sweep of shear amplitude so that the following phases preceded each test: (1) mounting of the sample; (2) equilibration of the sample by keeping it for 15 min at the minimum strain (0.02% shear strain amplitude, 1 Hz); (3) strain sweep (0.02% maximum shear strain amplitude, 1 Hz); (4) further equilibration (0.02% shear strain amplitude, 1 Hz). For each test, at least three repetitions were performed. Test Configurations. The orientation of the specimens and of the strain directions with respect to the rubber sheet is schematically depicted in Figure 1C,D. The rubber sheet is sketched with a broken line. The two specimens are the parallelepipeds drawn with the solid line: they are perpendicularly oriented with respect to each other. A Cartesian coordinate system is displayed to describe the loading direction of the specimens. Axes 1 and 2 are arbitrarily assigned to the direction of the two main dimensions of the rubber sheet and are not associated with any specific processing direction of the sheet. The plane of the rubber sheet is thus defined plane 1−2. Axis 3 refers to the direction perpendicular to the sheet main plane, crossing the 3 mm thickness, and it is the direction of application of the pressure in the molding process. To study the shear response in different directions, the tests were carried out according to four test configurations (see Figure 1C,D). Such configurations can be divided in two groups. When the shear stress was applied on a and a′ faces of the sample, deformation occurred through the sample thickness and tests are named in the text below as “through thickness” (TT) tests (see Figure 1C). When the shear stress was applied on b and b′ faces of the sample, deformation occurred in planes parallel to the sample surface and tests are named in the text below as “in-plane” (IP) tests (see Figure 1D). Such configurations are labeled “through thickness” and “in-plane” with reference to basic definitions of composites laminate,40 widely used also for fiber-reinforced polymer composites. For both TT and IP configurations, two specimens, perpendicularly oriented to each other, were tested. Bright Field Transmission Electron Microscopy (BF-TEM) and Selected Area Electron Diffraction (SAED) Analyses. BFTEM analysis was coupled with selected area electron diffraction (SAED) analysis. BF-TEM micrographs and SAED patterns were acquired using a Philips CM200 electron microscope operating at 200 kV equipped with a field emission gun filament. A Gatan US 1000 CCD camera was used, and 2048 × 2048 pixels images with 256 gray levels were recorded. Thin film cross sections (approximately 70−100 nm thick) were prepared using a Leica Ultracut 7 ultramicrotome (sample temperature: −130 °C) with a diamond knife and mounted onto 300 mesh copper grids. No heavy metal staining methods were used. The procedure for the preparation of a ultrathin slice for TEM

Figure 1. Scheme (A) of a specimen and (B) of the shear sandwich configuration used for the dynamic-mechanical tests; schematic picture of a rubber sheet with coordinates of the reference system and of two specimens (white parallelepipeds), with indication of the applied shear strains (solid arrows) in (C) “through thickness” and (D) “in-plane” configurations. Measurements. Dynamic mechanical tests were performed, at room temperature and 1 Hz, on the cured rubber nanocomposites by a dynamic-mechanical analyzer Q800 (TA Instruments), in strain sweep mode, for strain amplitude levels ranging from a minimum of 0.02% up

Figure 2. Procedure for the preparation of an ultrathin slice for TEM analysis, starting from the cut of a parallelepiped sample from a sheet. 8688

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Macromolecules analysis is shown in Figure 2. A sheet was taken, perpendicular to axis 3 (shown also in Figure 1C,D). A parallelepiped sample was taken from this sheet, and a slice (a section) was cut by ultramicrotomy from such bar. Analyzed sections were thus oriented parallel to axis 3 (shown also in Figure 1C,D) and allowed to observe the nanocomposite through the sheet thickness. Intensity integration along the Debye rings of the composite SAED pattern was performed using ImageJ plugin implemented directly in MAUD (Materials Analysis Using Diffraction), which allows to import directly two-dimensional diffraction data coming from the TEM camera or imaging-plate detectors.38,41,42 Intensity profile obtained summing up all the one-dimensional plots (full integration along Debye−Scherrer rings) was plotted as a function of the length of the reciprocal lattice vector Q (Å−1). Q is defined in accordance with Bragg’s law: Q = (4π/λ) sin(θ), where θ is half of the scattering angle and λ is the radiation wavelength. Circular intensity profile of reflection ring was obtained using ImageJ and plotted as a function of azimuthal angle.

2 (see Figure 1D), and an orthogonal plane is indexed as plane 3−2 (see Figure 1C). In dynamic-mechanical tests, simple shear amplitudes were applied: Figure 3 shows γ12 for an “in-plane” test and γ32 for a “through thickness” test (see Figure 1). In Figure 3, it is also shown the CNT organization in the nanocomposites. This aspect will be discussed below in the text. IP tests, as that shown in Figure 3, allowed to evaluate the dynamic shear property G12 (G indicates both G′ and G″). Additional IP tests were carried out on parallelepiped samples orthogonally taken from the material sheet, by applying a simple shear amplitude γ21 (see Figure 1D), determining the properties G21. The objective of these additional tests was to compare properties G12 and G21, which are expected to assume the same values, according to Cauchy hypothesis, to verify the reliability of the test configuration. TT tests were carried out by applying a shear load on the sheet faces lying normally to direction 3 (see Figure 1C) and acting along direction 1 or 2. The TT tests allowed to measure the shear properties G31 and G32. To sum up, γ12, γ21, γ31, and γ32 shear strains were applied to evaluate the G12, G21, G31, and G32 properties, respectively. Storage modulus G′, loss modulus G″, and tan δ (G″/G′) were measured as a function of shear strain amplitude. Storage moduli values, obtained for neat cross-linked NR matrix at low and high strain, are shown in Table 2: dependence on strain is seen to be identical independently from the testing direction, and this result is here shown by displaying the values measured at minimum and maximum strain for the four sample types. The results thus reveal that the material behaves isotropically. This allows to exclude the presence of significant polymer chains orientations and, as a consequence, to fully ascribe anisotropic properties of the nanocomposites to CNT organization in the polymer matrix. Figure 4 shows the dependence of storage modulus G′ (Figure 4A), loss modulus G″ (Figure 4B), and loss factor G″/ G′ (tan δ) on the strain amplitude for CNT-35 nanocomposite. Curves in Figure 4A reveal remarkable differences between G′ values obtained from either IP or TT tests. In Figure 4A, it can be seen that at small strains curves cluster in two bundles (corresponding to “in-plane” and “through thickness” loading configurations), and they tend to merge into a single curve at high strain levels. The decrease of G′ is thus appreciable whatever is the shear load direction. Table 3 reports the average values, with standard deviations, of storage modulus G′ obtained with the four different test configurations, at minimum strain (0.025%) and at large strain (28%). G′28% was obtained as average of the values of the only samples able to achieve such a strain value. Curves in Figure 4A and data in Table 3 reveal that “through thickness” G31 and G32 moduli differ from “in-plane” moduli, and this result is indicative of an orthotropic behavior. As theoretically expected, “in-plane” moduli are equal (G′12 = G′21) according to the conditions of symmetry in the Cauchy strain (γ) and stress (τ) tensors (i.e., γ12 = γ21; τ12 = τ21, respectively). “Through thickness” moduli have very similar

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RESULTS As mentioned in the Introduction, the effect of CNT orientation was studied in the absence (at 4 php CNT content) and in the presence (at 15 and 35 php CNT content) of continuous filler network. The CNT contents were chosen on the basis of a previous study by some of the authors,21 which showed a CNT percolation threshold in poly(1,4-cis-isoprene) matrix at about 7 php. Dynamic-Mechanical Tests. Dynamic-mechanical tests were carried out on peroxide cross-linked neat polymer matrix and on CNT-based nanocomposites to measure dynamicmechanical moduli in different directions. With these experiments, the effect of CNT on the isotropy of dynamicmechanical properties was investigated. As described in detail in the Experimental Part, parallelepiped samples were taken from material sheets that were first calendered on a two roll mill and then cross-linked in a press with parallel plates, under pressure. Two main test configurations (TT and IP, respectively) were studied, as is shown in Figure 3. The plane of the material sheet is indexed as plane 1−

Figure 3. Schematic representation of dynamic-mechanical tests. Parallelepiped samples were taken from a cross-linked nanocomposite sheet. Solid arrows indicate sample faces. Shear stresses, indicated by dashed arrows, were applied on a and a′ faces in “through thickness” tests and on b and b′ faces in “in-plane” tests. CNT organization is shown with a lateral view in “through thickness” tests and with a top view in “in-plane” tests.

Table 2. Storage Modulus G′ at Minimum (0.025%) and High (25%) Shear Strain Amplitude of the Neat Cross-Linked NR Matrix from Tests Performed on the Four Different Sample Configurations modulus [MPa]

G21

G12

G31

G32

G′0.025% G′25%

0.460 ± 0.009 0.451 ± 0.009

0.452 ± 0.002 0.444 ± 0.002

0.442 ± 0.008 0.437 ± 0.006

0.440 ± 0.005 0.435 ± 0.004

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Table 4. Average Values of Storage Modulus G′ at Minimum (0.025%) and High (28%) Shear Strain Amplitude, Payne Effect ΔG′ (= G′0.025% − G′28%), Normalized Payne Effect (ΔG′/G′0.025%), Maximum Value of Loss Modulus (G″max), and Anisotropy Index from Tests Performed with the Two Different Configurations for CNT-35, CNT-15, and CNT-4 Nanocomposites “in-plane” configuration G′0.025% G′28% ΔG′ = G′0.025% − G′28% ΔG′/G′0.025% × 100 G″max G′0.025% G′28% ΔG′ = G′0.025% G′28% ΔG′/G′0.025% × 100 G″max

G21

G12

G31

G32

20.1 ± 1.3 2.4 ± 0.1

21.4 ± 2.6 2.2a

10.7 ± 0.5 1.8 ± 0.1

10.4 ± 1.2 1.76 ± 0.01

a

1.95 ± 0.06 1.32 ± 0.05 1.96 ± 0.05

87.8 ± 0.3%

1.06 ± 0.01

82.9 ± 1.3%

3.4 ± 0.3 MPa 1.66 ± 0.07 MPa CNT-15 nanocomposite 3.6 ± 0.2 MPa 1.61 ± 0.09 MPa 1.1 ± 0.2 MPa 0.70 ± 0.01 MPa − 2.5 ± 0.2 MPa 0.90 ± 0.08 MPa 69 ± 4%

56 ± 2%

2.07 ± 0.04 2.25 ± 0.04 1.6 ± 0.1 2.76 ± 0.04 1.23 ± 0.05 2.35 ± 0.04 1.88 ± 0.04 1.31 ± 0.04 4.46 ± 0.05 2.38 ± 0.07 2.13 ± 0.04

over the low strain modulus value (ΔG′/G′0.025%), and the maximum values of loss modulus (G″max). In line with the literature on polymer melt and elastomers,16,23,24 ΔG′ was taken as the index of Payne effect. For every property, the degree of anisotropy (anisotropy index) was evaluated as the ratio of the property measured in IP configuration over that measured in TT configuration (e.g., G′IP/G′TT or ΔG′IP/ΔG′TT ratio). In the literature,35 for milled thin rubber films, anisotropy is analogously estimated as the ratio of the mechanical properties along the milling direction (called “machine” direction35) and transverse directions. Average G′0.025% values for CNT-35 in Table 4 are as follows: about 20.5 MPa and about 10.5 MPa for IP and TT configurations, respectively. Lower difference between G′IP and G′TT was observed at 28% strain: about 2.4 and 1.8 MPa, respectively. The degree of anisotropy in CNT-35, evaluated as the ratio of the property measured in IP configuration over that measured in TT configuration, is quite high at small deformations: at 0.025% as shear strain amplitude, the ratio G′IP/G′TT is 1.95. At high strain (28%) the degree of anisotropy is remarkably reduced, 1.32. The anisotropy index of the Payne effect is equal to about 1.96, and it appears to mirror the anisotropy of G′0.025% modulus (1.95). This might be the reason why the anisotropy of Payne effect, normalized with respect to low strain modulus (ΔG′/G′0.025%), vanishes: in fact, ΔG′ normalized with respect to G′0.025% is only slightly larger (88%

Table 3. Storage Modulus G′ at Minimum (0.025%) and High (28%) Shear Strain Amplitude from Tests Performed on CNT-35 with the Four Different Configurations G′0.025% G′28%

anisotropy index

CNT-35 nanocomposite 20.5 ± 1.8 MPa 10.5 ± 0.9 MPa 2.4 ± 0.1 MPa 1.8 ± 0.1 MPa 16.9 ± 0.7 MPa 8.6 ± 0.7 MPa

0.75 ± 0.06 0.32 ± 0.02 MPa MPa CNT-4 nanocomposite G′0.025% 1.12 ± 0.02 0.59 ± 0.04 MPa MPa G′28% 0.634 ± 0.008 0.49 ± 0.02 MPa MPa 0.11 ± 0.02 MPa ΔG′ = G′0.025% − 0.48 ± 0.02 G′28% MPa ΔG′/G′0.025% × 43.3 ± 0.8% 18 ± 3% 100 G″max 0.092 ± 0.005 0.043 ± 0.002 MPa MPa

Figure 4. (A) Shear storage modulus, (B) loss modulus, and (C) loss factor vs shear strain amplitude curves obtained with different test configurations (see Figures 1B and 1C) for CNT-35. A sketch of the expected CNT structure in IP and TT configuration is also displayed.

modulus [MPa]

“through thickness” configuration

Modulus evaluated on a single test.

values (G′31 = G′32), and this result seems to suggest that the properties in the sheet are independent of the shear load direction. Therefore, the sheet behaves as a transversally isotropic solid. In light of this result, a direct comparison of “through thickness” and “in-plane” properties is shown in Table 4, not only for CNT-35 but also for the other nanocomposites. In the following, for the sake of brevity, the properties referred to “inplane” and “through thickness” configuration will be indicated by IP and TT subscripts, respectively. Table 4 shows the values of “in-plane” G′ (G′IP) and “through thickness” G′ (G′TT), obtained as the average of G′21 and G′12 values and G′31 and G′32 values, respectively, for CNT-35, CNT-15, and CNT-4 nanocomposites. Table 4 also shows the average values of the overall storage modulus reduction, ΔG′, from 0.025% to 28% strain (ΔG′ = G′0.025% − G′28%), the same reduction normalized 8690

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Macromolecules versus 83%) for “in-plane” configuration than for “through thickness” configuration. Curves in Figure 4B show that also the specific values of G″ and its dependence on strain amplitude are a function of the load direction. G″ curves, similarly to G′ curves, are divided in two groups, for “in-plane” and for “through thickness” shear, and in the whole strain range explored, “in-plane” loss moduli are higher than “through thickness” moduli. This should be expected, in consideration of G′ values, as it is known the linear correlation between ΔG′ and the maximum value of G″. The loss modulus peak occurs at a lower strain amplitude (about 0.5%) for the “in-plane” testing than for the “through thickness” one (about 0.7%), indicating an earlier occurrence of maximum dissipation on the strain scale. This finding suggests a weaker behavior of the filler network when strained in IP configuration. Differently from G′ curves, loss moduli values do not tend to collapse at high strains but remain separated, with a value depending on the loading direction. It can be speculated that different mechanisms lead to the reduction of G′ values at moderately/high strain levels and therefore lead to different levels of energy dissipation depending on the direction of strain. Such mechanisms will be discussed below in the text. Loss factor curves, from “in-plane” and “through thickness” tests, are shown in Figure 4C. As tan δ is the G″/G′ ratio, the trend of tan δ vs strain amplitude curves can be rationalized in the light of G′ and G″ curves in Figures 4A and 4B, respectively. As both G″ and G′ moduli of CNT-35 composite are increased by almost the same factor from TT to IP configuration, the curves of their ratio overlap at low strain. Since G′ vs strain curves (from “in-plane” and “through thickness” tests) overlap at large strain, whereas G″ curves show larger values for “inplane” tests, tan δ curves diverge at large strain and the “inplane” configuration gives higher tan δ values. This finding indicates that at high strains the filler network of a nanocomposites based on NR and 35 php of CNT dissipates energy differently depending on the strain direction. Figures 5A and 5B show the dependence of storage modulus G′ and loss modulus G″ on the strain amplitude for CNT-15 and CNT-4 nanocomposites, respectively. To allow a direct comparison, representative curves for CNT-4 nanocomposites are also shown in Figures 5A and 5B. Average “in-plane” and “through thickness” values are in Table 4. Absolute values of G′ and G″ were found to be lower for CNT-15 with respect to CNT-35, as expected. Besides that, CNT-15 nanocomposite revealed many similarities with CNT35: storage and loss moduli are higher for IP configuration; maximum of G″ occurs at lower strain in IP tests; the loss factor curves are overlapped at small strain while they tend to diverge for higher strains; a remarkable Payne effect was observed independently of the shear load direction. Interesting difference between CNT-35 and CNT-15 appears by inspecting the anisotropy of the nanocomposites: anisotropy indexes are higher for CNT-15. G′IP/G′TT is only slightly higher: 2.25 for CNT-15 and 1.95 for CNT-35. The anisotropy index clearly increases for ΔG′, from 2 (CNT-35) to 2.8 (CNT-15). Indeed, ΔG′ normalized with respect to G′0.025% is clearly larger (69% versus 56%) for IP configuration than for TT configuration, and the anisotropy of the Payne effect, normalized with respect to low strain modulus (ΔG′/G′0.025%), is appreciably larger than 1. A lower CNT content with respect to 35 php seems thus to promote higher anisotropy level in the polymer nanocomposite. CNT-4 nanocomposite showed the lowest absolute values for the storage and loss moduli and most of the above-reported

Figure 5. (A, C) Shear storage modulus and (B, D) loss modulus vs shear strain amplitude curves obtained with different test configurations (see Figures 1C and 1D) for (A, B) CNT-15 and (C, D) CNT-4. Curves representative of CNT-4 are also displayed in (A) and (B) as dashed lines. A sketch of the expected CNT structure in IP and TT configuration is also displayed.

similarities, with both CNT-15 and CNT-35. A remarkable difference with CNT-15 and CNT-35 appears by observing curves in Figure 5A: decrease of G′ with strain can be clearly detected only in IP configuration. However, it is definitely relevant for this configuration that Payne effect occurs for CNT content below the percolation threshold. As is shown in Table 4, anisotropy indexes for G′ and ΔG′ are appreciably larger than 1, confirming that CNT are able to develop a detectable anisotropy of small strain elastic moduli already at small CNT content, below CNT percolation threshold. In particular, values of anisotropy index of G′ for CNT-4 are close to those of CNT35, and this suggests that there is not an evident correlation 8691

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Figure 6. Bright field TEM micrograph of ultrathin slice of CNT-35 nanocomposite at (A) low and (B) high magnification. The red arrow in (B) indicates the long axis CNT preferential orientation extracted by SAED data analysis. (C) Selected area electron diffraction pattern of the area observed in (B). (D) One-dimensional intensity profile representing the full integration along the Debye rings in (C). (E) Electron diffraction intensity profile as a function of azimuthal angle for the (002) Debye−Scherrer ring.

responses are found in the sheet plane. Differences between CNT-35, CNT-15, and CNT-4 and, in particular, the larger values of anisotropic indexes, in particular of ΔG′, could be ascribed to a higher orientation degree of CNT in samples with lower CNT content. In the literature, it was reported that CNT alignment decreases as the CNT content increases.31,35 Restriction of CNT motions during the material processing, caused by their agglomeration, was hypothesized. However, before embarking on undue speculations, it is definitely worth examining results arising from TEM analysis and electron diffraction measurements. TEM Analysis and Electron Diffraction Measurements. BF-TEM analysis was performed on nanocomposites contain-

between such index and the CNT content. On the contrary, the anisotropy index of ΔG′ is much larger for CNT-4, and a consistent increase of this index appears by decreasing CNT content from 35 to 15 to 4 php: values are 1.96, 2.76, and 4.46 respectively. In the case of CNT-4, ΔG′ normalized with respect to G′0.025% is clearly larger (43% versus 18%) for “inplane” than for “through thickness” configuration, and the anisotropy of the Payne effect, normalized with respect to low strain modulus (ΔG′/G′0.025%), exceeds 2 (2.38). Results shown so far clearly indicate an orthotropic (more specifically a transversally isotropic) behavior of all investigated nanocomposites, which means different responses are found between “in-plane” and “through thickness” shear, and the same 8692

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Figure 7. Bright field TEM micrograph of ultrathin slice of CNT-15 nanocomposite at (A) low and (B) high magnification. The red arrow in (B) indicates the long axis CNT preferential orientation extracted by SAED data analysis. (C) Selected area electron diffraction pattern of the area observed in (B). (D) Electron diffraction intensity profile as a function of azimuthal angle for the (002) Debye−Scherrer ring.

ing 35 and 15 php of CNT. Such analysis provides information on CNT distribution within material sheets oriented parallel to axis 3 and gives a picture of sheets cut through the thickness of the specimen (see Figures 1C,D and 2). Analyses were performed on undeformed sheets. TEM micrographs of CNT-35, the system with the lowest anisotropy levels, are shown in Figures 6A and 6B. CNT appear highly entangled, forming a continuous network. Isolated tubes can be hardly detected, and it is not reliable to assess CNT orientation solely on the basis of image analysis techniques. BF-TEM investigations were thus combined with electron diffraction measurements, which allow to identify the presence of crystallographic textures. Since CNT exhibit features of a polycrystalline material, in this work, this technique has been tentatively used for the first time to highlight preferential CNT orientation inside the polymer matrix, even in the presence of agglomerated and entangled

nanostructures. Relevant SAED ring pattern of CNT-35 is shown in Figure 6C. The graph shows characteristic reflections of multiwalled CNT,43−45 which are present in the pattern at the following Q values (see Figure 6D): (002) at 1.85 Å−1, (100) at 2.98 Å−1, and the (004) at 3.65 Å−1. The (002) reflection ring is due to the concentric arrangements of nanotube walls and corresponds to a d-spacing of 0.339 nm. The intensity distribution along this specific ring exhibits the presence of two lower intensity sectors (blue boxes area in Figure 6C) and two higher intensity sectors (red boxes area in Figure 6C). This inhomogeneity is highlighted by the circular intensity profile plot of the (002) reflection ring (reported in Figure 6E) that clearly shows high (a, c) and low (b, d) intensity zones. This experimental evidence reveals a pronounced crystallographic texture due to anisotropic distributions of (002) concentric crystallographic planes belonging to the carbon nanotubes. This means that CNT 8693

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These findings are in agreement with results from BF-TEM analysis combined with electron diffraction measurements. Preferential orientation of nanotubes was detected inside “through thickness” sheets and TEM images of CNT-15 showed alternate areas with high CNT content and almost free of CNT. Interestingly, the anisotropy index of the Payne effect (ΔG′IP/ΔG′TT) was clearly larger than 1 and was observed to increase as the CNT content decreased: values are 1.96, 2.76, and 4.46 for CNT-35, CNT-15, and CNT-4, respectively. It is worth reminding here that the anisotropy index of ΔG′, normalized with respect to G′ at minimum strain, increases as well with the decrease of CNT content. Such increase of anisotropy level with reduced CNT content can be ascribed to a higher CNT alignment for low filler content due to a progressive restriction of CNT motions during the material processing as the filler amount increases, as tubes are entangled and give rise to agglomerates, to an extent that increases with the CNT content, as already commented. Also, these results are in line with the picture of the material structure, as interconnection of layers by CNT is expected to decrease as the CNT content decreases. Therefore, the interpretation of the different Payne effects depending on the loading directions could be based only on the orientation of CNT structures, whereas the soft elastomeric matrix, which should have weak interactions with the filler particles, should not affect the anisotropy levels. However, the remarkable anisotropic effects of G′ moduli found at small strain values appear to vanish at high strain for CNT-35 and to decrease with increasing strain for CNT-15 and CNT-4. TEM analysis was useful to provide insights into the behavior at small strain, but in order to justify the results of the properties measured at high strains, further investigations would be needed. Nevertheless, a couple of hypothesis can be formulated. First, it seems reasonable to assume that the strain modifies CNT organization, leveling off differences in CNT orientation and thus reducing the anisotropy level. In the nanocomposites studied in this work, application of a shear strain is likely to modify CNT organization and orientation and different mechanisms can be hypothesized for “in-plane” and “through thickness” loadings. In the “in-plane” testing configuration, shear stresses (see Figure 3) are perpendicular to CNT bundles, randomly distributed in the plane, and promote a pronounced CNT reorganization, that can be at the origin of higher dissipative behavior and modulus drop. It is worth observing that nonlinear mechanical behavior is detected in IP configuration at 4 php, below the threshold of CNT percolation. In the “through thickness” testing configuration, the shear stresses (see Figure 3) are parallel to CNT bundles, leading to a mutual sliding of the CNT layers. Also these results are in line with the picture of the material structure, as interconnection of layers by CNT is expected to decrease as the CNT content decreases. Effect of strain on CNT organization is documented in the literature. CNT bundles orientation and breakdown at the highest tensile strains occur.18,46 Reorientation of CNT under strain was proposed as one of the mechanisms influencing the mechanical properties of CNT/NR nanocomposites at high strain levels47 and was measured by Raman spectroscopy.32 A second hypothesis that could be suggested to explain the anisotropy results at high strains is that the well-known strain induced crystallization behavior of natural rubber48,49 could

assume a preferential orientation within the nanocomposite sheet. The direction of CNT preferential orientation can be obtained by SAED ring pattern analysis and can be correlated with the corresponding BF-TEM image, acquired on the same area. The red arrow in Figure 6B represents the extrapolated preferential CNT orientation and is indeed parallel to most tubes, which belong to entangled nanostructures. It could be assumed that CNT are organized in a layered structure and that they connect with interpenetrated layers. However, for CNT35, it is evident that the high CNT content does not allow to verify CNT orientation extrapolated from electron diffraction data by the observation of TEM image. Therefore, in order to have a confirmation of the reliability of SAED technique for the measurement of CNT orientation, a composite with a lower CNT content (CNT-15) was also analyzed. Figure 7 shows the results of BF-TEM and electron diffraction measurements performed on CNT-15. The BF-TEM image of a representative sample, shown in Figure 7A, reveals areas with different CNT concentration. In Figure 7B, a channel almost free of CNT can be observed in between two regions with higher CNT content, and a zone with low CNT content is shown: single tubes (and their ends) are visible. Preferential orientation can be clearly identified for most CNT. SAED pattern and the circular intensity profile of the (002) reflection in Figures 7C and 7D, respectively, reveal inhomogeneous intensity distribution along (002) ring as a clear indication of CNT orientation. The red arrow in Figure 7B indicates the preferential long axis tube direction extrapolated from the electron diffraction pattern: the good agreement between TEM image and electron diffraction data appears unequivocal. For both CNT-15 and CNT-35 samples the CNT orientation was collected on several regions of the ultramicrotomated slice. CNT orientation direction closely resemble the direction of the reference axis 2 (see Figures 1−3). This result suggests a preferential orientation of CNT along a direction which is perpendicular to the pressure applied in the compression molding phase. These results confirm the reliability of such technique to detect the presence of preferentially aligned nanotubes inside CNT based nanocomposites and support the choice of diffraction data analysis as a method to explore the alignment of crystalline anisotropic nanofillers inside a polymer matrix.

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DISCUSSION Experimental data presented above support the following picture of nanocomposites based on natural rubber, with different CNT content (35, 15, and 4 php), prepared via mechanical mixing and through calendering and compression molding. CNT are organized in a layered structure. They are randomly distributed inside a layer and connect different layers to an extent that depend on CNT content. Such organization is sketched in Figures 3−5. In fact, property values obtained from either “in-plane” or “through thickness” tests, particularly at low strains, are very different from each other: for example, the anisotropy index (G′IP/G′TT)0.025% is between 1.9 and 2.3. Hence, nanocomposites have an orthotropic behavior: different responses are found for “in-plane” and “through thickness” shear deformations. Vice versa, very similar values have been found for “through thickness” G31 and G32 moduli, suggesting a transversally isotropic material. 8694

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mask the effects of CNT orientation. NR gives rise to straininduced crystallization (SIC) after a certain level of strain, and it has been reported that CNT (at 1.75 wt %) promote such crystallization at lower strain levels with respect to the neat NR.50 This hypothesis is in line with the observation that at the highest CNT content (35 php) the high strain modulus values tend to overlap. On the basis of the results obtained in this work, the effect of CNT organization seems to prevail at low strain. The lower anisotropy degree found at higher strains might be due to the destructuring of filler network by the applied strain and also to the effect of the strain crystallization of the NR matrix.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected], tel +390303715908, fax +390303702448. Present Address

S.M.: Pirelli Tyre, Viale Piero e Alberto Pirelli 25, 20126 Milano, Italy. Notes

The authors declare no competing financial interest.

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REFERENCES

(1) PAS 71:2005 - Vocabulary. Nanoparticles. (2) Araby, S.; Meng, Q.; Zhang, L.; Zaman, I.; Majewski, P.; Ma, J. Elastomeric composites based on carbon nanomaterials. Nanotechnology 2015, 26, 112001. (3) Iijima, S.; Ichihashi, T. Single-shell carbon nanotubes of 1-nm diameter. Nature 1993, 363, 603−605. (4) Bethune, D. S.; Kiang, C. H.; De Vries, M. S.; Gorman, G.; Savoy, R.; Vazquez, J.; Beyers, R. Cobalt-catalysed growth of carbon nanotubes with single-atomic-layer walls. Nature 1993, 363, 605−607. (5) Iijima, S. Helical microtubules of graphitic carbon. Nature 1991, 354, 56−58. (6) Monthioux, M.; Kuznetsov, V. L. Who should be given the credit for the discovery of carbon nanotubes? Carbon 2006, 44, 1621−1623. (7) Novoselov, K. S.; Geim, A. K.; Morozov, S. V.; Jiang, D.; Zhang, Y.; Dubonos, S. V.; Grigorieva, I. V.; Firsov, A. A. Electric field effect in atomically thin carbon films. Science 2004, 306, 666−669. (8) Geim, A. K.; Novoselov, K. S. The rise of graphene. Nat. Mater. 2007, 6, 183−191. (9) Allen, M. J.; Tung, V. C.; Kaner, R. B. Honeycomb carbon: A review of graphene. Chem. Rev. 2010, 110, 132−145. (10) Zhu, Y.; Murali, S.; Cai, W.; Li, X.; Suk, J. W.; Potts, J. R.; Ruoff, R. S. Graphene and graphene oxide: Synthesis, properties, and applications. Adv. Mater. 2010, 22, 3906−3924. (11) Geng, Y.; Wang, S. J.; Kim, J. K. Preparation of graphite nanoplatelets and graphene sheets. J. Colloid Interface Sci. 2009, 336, 592−598. (12) Kavan, L.; Yum, J. H.; Graetzel, M. Optically transparent cathode for dye-sensitized solar cells based on graphene nanoplatelets. ACS Nano 2011, 5, 165−172. (13) Nieto, A.; Lahiri, D.; Agarwal, A. Synthesis and properties of bulk graphene nanoplatelets consolidated by spark plasma sintering. Carbon 2012, 50, 4068−4077. (14) Paton, K. R.; Varrla, E.; Backes, C.; Smith, R. J.; Khan, U.; O’Neill, A.; Boland, C.; Lotya, M.; Istrate, O. M.; King, P.; Higgins, T.; Barwich, S.; May, P.; Puczkarski, P.; Ahmed, I.; Moebius, M.; Pettersson, H.; Long, E.; Coelho, J.; O’Brien, S. E.; McGuire, E. K.; Sanchez, B. M.; Duesberg, G. S.; McEvoy, N.; Pennycook, T. J. Scalable production of large quantities of defect-free few-layer graphene by shear exfoliation in liquids. Nat. Mater. 2014, 13, 624− 630. (15) Peigney, A.; Laurent, C.; Flahaut, E.; Bacsa, R. R.; Rousset, A. Specific surface area of carbon nanotubes and bundles of carbon nanotubes. Carbon 2001, 39, 507−514. (16) Donnet, J. B.; Custodero, E. Reinforcement of elastomers by particulate fillers. In The Science and Technology of Rubber, 3rd ed.; Mark, J. E., Erman, B., Eirich, F. R., Eds.; Elsevier Academic Press: 2005; Chapter 8, pp 367−400. (17) Bokobza, L. Multiwall carbon nanotube-filled natural rubber: Electrical and mechanical properties. eXPRESS Polym. Lett. 2012, 6, 213−223. (18) Bokobza, L. Multiwall carbon nanotube elastomeric composites: A review. Polymer 2007, 48, 4907−4920. (19) Coleman, J. N.; Khan, U.; Blau, W. J.; Gun’ko, Y. K. Small but strong: A review of the mechanical properties of carbon nanotube− polymer composites. Carbon 2006, 44, 1624−1652.

CONCLUSIONS

Nanocomposites were prepared, via mechanical mixing, with natural rubber and different CNT contents: 4, 15, and 35 php. They experienced calendering and compression molding. Dynamic-mechanical measurements were carried out in the shear mode on peroxide cross-linked samples, investigating dynamic mechanical response, at various strain levels, in different directions, here defined as “in-plane” and “through thickness”. Structure of nanocomposites was studied by means of BF-TEM coupled with electron diffraction measurements. Dynamic-mechanical behavior was typical of a transversally isotropic material: moduli evaluated “in-plane” were remarkably different from those evaluated “through thickness”, revealing the anisotropy of nanocomposites, but the “through thickness” moduli measured along different strain directions in the sheet had very similar values, indicating a transversally isotropic material. The anisotropy index of G′ at minimum strain (G′IP/ G′TT)0.025% was between 1.9 and 2.3, and the anisotropy index of ΔG′ (ΔG′IP/ΔG′TT) was from 1.96 for CNT-35 to 4.46 for CNT-4. Anisotropy was thus observed to increase as CNT content decreased. BF-TEM analysis coupled with electron diffraction measurements revealed preferential orientation of nanotubes inside sheets cut “through thickness”, with alternate areas containing large or very low amount of CNT. CNT seem to be prevailingly organized in bundles randomly distributed inside layers disposed almost parallel to each other. They connect such layers, to an extent that depends on their amount in the nanocomposite. This work shows a remarkable anisotropic nonlinear mechanical behavior of CNT-based nanocomposites. Thanks to the combination of dynamic-mechanical tests with TEM and electron diffraction measurements, it is shown that such anisotropic behavior is the response of the organization in the polymer matrix of nanofillers with high aspect ratio, such as CNT. Hence, this work shows that energy dissipation is not an isotropic phenomenon in polymer nanocomposites based on CNT and aims at giving a contribution in order to achieve a better control of energy dissipation in polymer nanocomposites based on carbon nanofibers. As working hypothesis for future research, it could be hypothesized that reduction of energy dissipation and of the anisotropy of such phenomenon could be pursued by a promoting a proper structuring of CNT, reducing their entanglement, and improving the CNT−polymer interaction at the expenses of CNT−CNT interaction. Tailormade CNT functionalization could thus play a key role. 8695

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Macromolecules (20) Maiti, M.; Bhattacharya, M.; Bhowmick, A. K. Elastomer nanocomposites. Rubber Chem. Technol. 2008, 81, 384−469. (21) Galimberti, M.; Coombs, M.; Riccio, P.; Riccò, T.; Passera, S.; Pandini, S.; Conzatti, L.; Ravasio, A.; Tritto, I. The Role of CNTs in Promoting Hybrid Filler Networking and Synergism with Carbon Black in the Mechanical Behavior of Filled Polyisoprene. Macromol. Mater. Eng. 2013, 298, 241−251. (22) Agnelli, S.; Cipolletti, V.; Musto, S.; Coombs, M.; Conzatti, L.; Pandini, S.; Riccò, T.; Galimberti, M. Interactive effects between carbon allotrope fillers on the mechanical reinforcement of polyisoprene based nanocomposites. eXPRESS Polym. Lett. 2014, 8 (6), 436−449. (23) Payne, A. R. In Reinforcement of Elastomers; Kraus, G., Ed.; Wiley-Interscience: New York, 1965. (24) Payne, A. R.; Whittaker, R. E. Low Strain Dynamic Properties of Filled Rubbers. Rubber Chem. Technol. 1971, 44, 440−478. (25) Bohm, G. A.; Tomaszewski, W.; Cole, W.; Hogan, T. Furthering the understanding of the non linear response of filler reinforced elastomers. Polymer 2010, 51, 2057−2068. (26) Heinrich, G.; Klüppel, M. Recent advances in the theory of filler networking in elastomers. Adv. Polym. Sci. 2002, 160, 1−44. (27) Sternstein, S. S.; Zhu, A.-J. Reinforcement Mechanism of Nanofilled Polymer Melts As Elucidated by Nonlinear Viscoelastic Behavior. Macromolecules 2002, 35, 7262−7273. (28) Montes, H.; Lequeux, F.; Berriot, J. Influence of the glass transition temperature gradient on the nonlinear viscoelastic behavior in reinforced elastomers. Macromolecules 2003, 36, 8107−8118. (29) Jancar, J.; Douglas, J. F.; Starr, F. W.; Kumar, S. K.; Cassagnau, P.; Lesser, A. J.; Sternstein, S. S.; Buehler, M. J. Current issues in research on structure-property relationships in polymer nanocomposites. Polymer 2010, 51, 3321−3343. (30) Gauthier, C.; Reynaud, E.; Vassoille, R.; Ladouce-Stelandre, L. Analysis of the non-linear viscoelastic behaviour of silica filled styrene butadiene rubber. Polymer 2004, 45, 2761−2771. (31) Ponnamma, D.; Sadasivuni, K. K.; Grohens, Y.; Guo, Q.; Thomas, S. Carbon nanotube based elastomer composites − an approach towards multifunctional materials. J. Mater. Chem. C 2014, 2, 8446−8485. (32) Frogley, M. D.; Ravich, D.; Wagner, H. D. Mechanical properties of carbon nanoparticle-reinforced elastomers. Compos. Sci. Technol. 2003, 63, 1647−1654. (33) Kim, Y. A.; Hayashi, T.; Endo, M.; Gotoh, Y.; Wada, N.; Seiyama, J. Fabrication of aligned carbon nanotube-filled rubber composite. Scr. Mater. 2006, 54, 31−35. (34) Kim, J. H.; Kim, G. H. Effect of Orientation and Content of Carbon Based Fillers on Thermal Conductivity of Ethylene-PropyleneDiene/Filler Composites. J. Appl. Polym. Sci. 2014, 131, 41000. (35) Kueseng, P.; Sae-oui, P.; Sirisinha, C.; Jacob, K. I.; Rattanasom, N. Anisotropic studies of multi-wall carbon nanotube (MWCNT)filled natural rubber (NR) and nitrile rubber (NBR) blends. Polym. Test. 2013, 32, 1229−1236. (36) Sun, X. M.; Chen, T.; Yang, Z. B.; Peng, H. S. The alignment of carbon nanotubes: An effective route to extend their excellent properties to macroscopic scale. Acc. Chem. Res. 2013, 46, 539−549. (37) Courty, S.; Mine, J.; Tajbakhsh, A. R.; Terentjev, E. M. Nematic elastomers with aligned carbon nanotubes: new electromechanical actuators. Europhys. Lett. 2003, 64, 654−660. (38) Boullay, P.; Lutterotti, L.; Chateigner, D.; Sicard, L. Fast microstructure and phase analyses of nanopowders using combined analysis of transmission electron microscopy scattering patterns. Acta Crystallogr., Sect. A: Found. Adv. 2014, 70, 448−456. (39) Gemmi, M.; Voltolini, M.; Ferretti, A. M.; Ponti, A. Quantitative texture analysis from powder like electron diffraction data. J. Appl. Crystallogr. 2011, 44, 454−461. (40) Analysis and Performance of Fiber Composites, 3rd ed.; Agarwal, B. D., Broutman, L. J., Chandrashekhara, K., Eds.; John Wiley & Sons: 2006.

(41) Ischia, G.; Wenk, H. R.; Lutterotti, L.; Berberich, F. Quantitative Rietveld texture analysis of zirconium from single synchrotron diffraction images. J. Appl. Crystallogr. 2005, 38, 377−380. (42) Lutterotti, L.; Vasin, R.; Wenk, H.-R. Rietveld texture analysis from synchrotron diffraction images. I. Calibration and basic analysis. Powder Diffr. 2014, 29, 76−84. (43) Reznik, D.; Olk, C. H.; Neumann, D. A.; Copley, J. R. D. X-ray powder diffraction from carbon nanotubes and nanoparticles. Phys. Rev. B: Condens. Matter Mater. Phys. 1995, 52, 116−124. (44) Cao, A.; Xu, C.; Liang, J.; Wu, D.; Wei, B. X-ray diffraction characterization on the alignment degree of carbon nanotubes. Chem. Phys. Lett. 2001, 344, 13−17. (45) Saito, Y.; Yoshikawa, T.; Bandow, S.; Tomita, M.; Hayashi, T. Interlayer spacings in carbon nanotubes. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 48, 1907−1909. (46) Bokobza, L.; Belin, C. Effect of Strain on the Properties of a Styrene−Butadiene Rubber Filled with Multiwall Carbon Nanotubes. J. Appl. Polym. Sci. 2007, 105, 2054−2061. (47) Deng, F.; Ito, M.; Noguchi, T.; Wang, L.; Ueki, H.; Niihara, K.; Kim, Y. A.; Endo, M.; Zheng, Q. S. Elucidation of the reinforcing mechanism in carbon nanotube/rubber nanocomposites. ACS Nano 2011, 5, 3858−3866. (48) Huneau, B. Strain-induced crystallization of natural rubber: a review of X-ray diffraction investigations. Rubber Chem. Technol. 2011, 84, 425−452. (49) Tosaka, M. Strain-Induced Crystallization of Crosslinked Natural Rubber As Revealed by X-ray Diffraction Using Synchrotron Radiation. Polym. J. 2007, 39, 1207−1220. (50) Fu, D. H.; Zhan, Y. H.; Yan, N.; Xia, H. S. A comparative investigation on strain induced crystallization for graphene and carbon nanotubes filled natural rubber composites. eXPRESS Polym. Lett. 2015, 9, 597−607.

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