European Polymer Journal

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1 European Polymer Journal 45 (9) Contents lists available at ScienceDirect European Polymer Journal journal homepage: Macromolecular Nanotechnology Application of strain time correspondence as a tool for structural analysis of acrylonitrile butadiene copolymer nanocomposites with various organoclay loadings Jae Woo Chung, Seok Jong Han, Seung-Yeop Kwak * Department of Materials Science and Engineering, Seoul National University, 599 Gwanak-ro, Gwanak-gu, Seoul , Republic of Korea article info Article history: Received 4 January 8 Received in revised form 9 September 8 Accepted October 8 Available online 17 October 8 Keywords: NBR (acrylonitrile butadiene copolymer) Nanocomposite Percolation Silicate structure STC (strain time correspondence) abstract Acrylonitrile butadiene copolymer (NBR) nanocomposites were prepared with varied silicate loadings by the melt mixing between NBR and organoclays (OCs) containing intercalants with different polarity and chain length. WXRD exhibited that the NBR nanocomposites had an intercalated structure with distinct differences in gallery height depending on the intercalant characteristics. However, WXRD failed to show a structural change with increasing silicate contents. Hence, tensile strain stress measurements were carried out at various strain rates (.16,.975, and.187 s 1 ), and then the results of tensile measurement applied to the strain time correspondence (STC) principle, resulting in the tensile modulus master curves of the NBR nanocomposites as a function of time. For pure NBR, a master curve was constructed using only the horizontal shift factor, indicating that the material was structurally homogeneous. However, the NBR nanocomposites required both vertical shift (modulus shift, C(a)) and horizontal shift to form the master curves, indicating structural heterogeneity ascribed to the domain structure such as silicate tactoid. From master curves, we found that NBR nanocomposite with OC having polar organic intercalant, NBROC3B, had the lowest n value in the nanocomposites. This indicates that NBROC3B had the most dispersive silicate structure in the nanocomposites due to the polar interaction, being in good agreement with WXRD results. In particular, STC was not applicable at all nanocomposites with silicate loadings over 8 wt%, regardless type of organoclay, and tensile strength and toughness of the nanocomposites with silicate loading of 8 wt% were better than expected. These could be explained as the network-like percolation of the silicate tactoids in all nanocomposites with silicate loadings over 8 wt%, which were consistent with the results observed from HR-TEM. Ó 8 Elsevier Ltd. All rights reserved. 1. Introduction In the past decades, clay has attracted a great deal of attention as filler due to its ability to remarkably enhance the mechanical and thermal properties of rubber. Clay is composed of silicate layers with thicknesses on the order of 1 nm and very high aspect ratios (e.g., 1 1), and has interlayer spacing between the stacked layers of about 1nm[1]. The extent to which the clay acts as a reinforcing * Corresponding author. Tel.: ; fax: address: sykwak@snu.ac.kr (S.-Y. Kwak). agent to improve the properties of a rubber is determined by the structure of the layered silicate in the rubber matrix []. Thus, it is important to analyze the silicate structure in rubber-based nanocomposites in order to understand their characteristics. In general, the silicate structure of nanocomposites has been elucidated by wide angle X-ray diffraction (WXRD), transmission electron microscopy (TEM), and small angle neutron scattering (SANS) [3 6]. WXRD is a powerful and simple technique for evaluating materials containing a periodic arrangement of elements, such as the silicate structure in intercalated nanocomposites. WXRD, however, provides little information about /$ - see front matter Ó 8 Elsevier Ltd. All rights reserved. doi:1.116/j.eurpolymj.8.1.9

2 8 J.W. Chung et al. / European Polymer Journal 45 (9) the silicate structure in exfoliated nanocomposites where the silicate is not arranged periodically. Besides, WXRD cannot show the three-dimensional silicate structure on a micro scale. TEM and SANS are also powerful tools, but they are very expensive and require experienced operators. Thus, another simple and useful tool is required to investigate the silicate structure of rubber basednanocomposites. It has been known for a long time that the mechanical behavior of rubbers at large deformation can be represented in the formation of a single master curve [7 9]. This implies that the nonlinear behavior caused by large deformations can be reduced to the corresponding linear behavior by using a reduced time-scale. This reduction, known as strain time correspondence (STC), is analogous to the well-known time temperature reduction [1]. Nakajima et al. have studied the applicability of STC in gum rubbers with a variety of molecular structure (e.g., branching and network) and the carbon black filled rubbers. Then, on the basis of their findings, they could identify rubbers into types such as linear type, STC type, and strain-hardening type depending on the molecular structure in rubber [11,1]. This implies that STC can be used as an indicator of the molecular structure of rubber systems. To the best of our knowledge, however, no analysis has yet been made of the silicate structure in rubber-based nanocomposites using STC. Herein, the acrylonitrile butadiene copolymer (NBR) nanocomposites with organoclays (OCs) having different intercalants were prepared with various silicate loadings by the melt mixing process. Then, the STC was applied to the strain stress results of NBR nanocomposites obtained from tensile measurement at various strain rates in order to investigate the silicate structure with the type and content of organoclay. It can be seen that the linearization of nonlinear behavior of NBR nanocomposite via the application of STC is used as a useful tool to characterize the silicate structure in NBR nanocomposites.. Theoretical background In the tensile test, as the deformation becomes larger, the moduli become dependent on the magnitude of deformation, indicating nonlinear behavior. However, the strain stress curves do not provide an adequate description of the nonlinear behavior in a systematic manner. Therefore, various attempts have been made to systematically treat the nonlinear behavior [13 15]. Among them is to extend the description of linear viscoelastic behavior to nonlinear behavior, preserving the separability of strain (or stress) and time. Schapery [16] has given general expressions for the separability of time and strain (or stress), incorporating time temperature equivalence; the expressions for the nonlinear creep compliance, J n (t), and the nonlinear stress relaxation modulus, G n (t), are given by [13,16] J n ðtþ ¼eðtÞ=r ¼ g J u þ g 1 g DJðt=a r Þ and G n ðtþ ¼rðtÞ=e ¼ h e G r þ h 1 h DGðt=a e Þ ð1þ ðþ Here, r is the stress and e is the strain; J u is the unrelaxed compliance and G r is the relaxed modulus [13]; DJðtÞ ¼JðtÞ J u and DGðtÞ ¼GðtÞ G r ; the symbols g o, g 1, g and a a are material properties which are function of stress, and h e, h 1, h and a e are material properties analogous to g, g 1, g and a a but functions of strain. The a a and a e represent the effect of mechanical excitation on the time scale; t/a a and t/a e are the stress-dependent intrinsic time and the strain-dependent intrinsic time, respectively. Once recognizing the thermodynamic (energetic) basis for the excitation of a material [15], it is then logical to seek the strain time correspondence. In tensile strain stress measurements, the tensile modulus, E, presented as the stress, a, based on the deformed cross-sectional area divided by the strain, e, which is in general a function of strain and time, t, may be expressed as [14]; aðt; eþ=e ¼ Eðt; eþ ð3þ with many gum rubbers, an application of strain time shift makes the tensile modulus as a function of reduced time only. This leads to a nonlinear expression, recognizing a similarity between the effect of thermal activation and that of mechanical activation on the material time [15]; aðt; eþ=e ¼ Eðt; a e Þ where a e has the same meaning as a e in the Schapery s equation for G n (t), and is a strain shift factor, analogous to the temperature shift factor, a T. This equation is a simple, special case of Schapery s representation, where h e = h 1 = h = 1. The applicability of the above equation has been demonstrated with gum rubbers such as NBR, SBR, polyethylacrylates, and epichlorohydrine rubber [11,17,18]. Also, the results show that the strain shift factor is a simple measure of strain and is equal to the reciprocal of the extension ratio; a e ¼ 1=a ð5þ ð4þ where a = l/l (l and l are the stretched and original length of the specimen) hence, rðt; eþ=e ¼ EðatÞ This reduction scheme is analogous to that based on the time temperature superposition and has been accordingly termed the strain time correspondence (STC) principle [19]. It is noteworthy that in the former, the shift factor (a T ) must be found from the experimental data whereas in the latter, the shift factor (a e ) is the extension ratio (a) and non-adjustable parameter (universal parameter). When the above equation is applicable, plots of loge(t,e) against logt at fixed values of e may be shifted along the time-axis by the magnitude of a, resulting in a master curve. Subsequently, critical examination has been made on the applicability of the STC with a variety of gum rubbers and filled rubbers [1, 3]. Generally speaking, this reduction scheme is shown to be obeyed by the structurally homogeneous gum rubbers, which are linear (gel-free) polymer, lightly branched polymer, and polymer containing crosslinked latex particles (i.e., microgel). However, it has been shown that the rubbers containing the filler ð6þ

3 J.W. Chung et al. / European Polymer Journal 45 (9) (i.e., carbon black) or the significant fractions of the extensively branched gel molecules (i.e., macrogel) do not obey the scheme. This may be because some structural heterogeneity to constrain the rubber molecule is introduced by the carbon black or the macrogel. This constraint structure is affected during the large deformation and brings about the modulus change. If the time and strain effect on modulus are separable, the modulus may be reduced. In such a case, in addition to the strain shift factor (a), a modulus shift factor (C(a)) is necessary [1,]; rðt; eþ=e ¼ EðatÞCðaÞ ð7þ where C(a) is a modulus shift factor that is decided by the empirical work and can be expressed as CH CH CH CH CH CH CN a C 18 H N + 3 C CH 3 C 18 C 18 b CðaÞ ¼a n ð8þ H 3 C N + CH 3 Hence, the degree of the modulus shift at a specific a is actually decided by the n value, which can be considered to be a measure of the structural heterogeneity [1,]; in order to calculate the n value, empirically obtained the C(a) is logarithmically plotted against a and then the slope of this plot is measured as a n value, because C(a) is related to the a in the form of logc(a) = nloga. 3. Experimental 3.1. Materials Uncrosslinked acrylonitrile butadiene copolymer (NBR) containing approximately 34 wt% acrylonitrile was obtained from Korea Kumho Petrochemical Co., Ltd. Its glass transition temperature, T g, is about C. The commercialized organophilic clays (OCs) composed of sodium montmorillonite (MMT) and various intercalants were supplied from Southern Clay Products, Inc. (USA) under the trade names Closite Ò A (OCA), Closite Ò 5A (OC5A), and Closite Ò 3B (OC3B). These OCs have a nonpolar intercalant with a long chain length, MHT, a nonpolar intercalant with a short chain length, MHTL8, and a polar intercalant, MTEtOH, respectively. Fig. 1 is the molecular structures of NBR and intercalants. The cationic exchange capacities (CECs) of the OCs were 95, 95, and 9 mequiv/1 g, respectively, and all OCs were used after vacuum drying at 8 C for 1 day. The intergallery heights of the OCs were calculated from their XRD patterns to be 1.57, 1.34, and.94 nm, respectively. 3.. Preparation of NBR nanocomposites NBR/organoclay nanocomposites were fabricated by the melt mixing process. The mixing was performed in a batch, non-intermeshing, counter-rotating Haake internal mixer with Banbury blades. The rotor speed was 1 rpm and the fill factor was about.7. NBR was first masticated at 65 C for 3 s before mixing and then mixed with OC at 9 C for 6 s. After mixing, the nanocomposites were annealed for 1 min at 9 C. Each nanocomposite was prepared with silicate contents of 1, 3, 5, 7, 8, and 1 wt% and the nanocomposites with OCA, OC5A, and O3B are referred to herein as NBROCAX, NBROC5AX, HOH CH C and NBROC3BX, respectively (where X is the silicate contents) Characterization CH CH(CH ) 3 CH 3 CH CH 3 c C 18 N + CH 3 d CH CH OH Fig. 1. Chemical structures of neat NBR and intercalants used in this study: (a) neat NBR, (b) dimethyl dehydrogenated tallow ammonium (MHT), (c) dimethyl hydrogenated tallow -ethylhexyl (MHTL8), and (d) methyl tallow bis(-hydroxyethyl) (MTEtOH). XRD measurements were performed on a MAC Science MXP 18A-HF X-ray diffractometer with CuK a radiation (k =.154 nm) at 4 kv and 1 ma. Diffraction spectra were obtained over a h range of 1 with a scanning rate of 3 min 1. XRD specimens were prepared by pressing at 1 C for min, followed by quenching to room temperature. The tensile measurements were carried out at strain rates of.16,.975, and.187 s 1 under ambient conditions using a universal testing machine (UTM) LLOYD LR1K. Force was measured using a 1 N load cell, and specimens were prepared from the ASTM D41-9T (dumbbell shape with a 15.5 mm gauge length). High-resolution transmission electron microscopic (HR-TEM) analysis was performed with a Jeol JEM 31CX at 3 kv. For the HR-TEM observation, ultra-thin cross-sections of the specimens were carried out by using a Leica Ultracut UCT ultracryomicrotome equipped with a diamond knife. The thickness of the HR-TEM specimens was controlled as approximately 8 nm, and the temperature in specimens was kept at 1 C using liquid nitrogen throughout the experiment. These cryotomed sections were then transferred to the copper grid.

4 8 J.W. Chung et al. / European Polymer Journal 45 (9) Results and discussion 4.1. WXRD The WXRD patterns of the OCs and NBR nanocomposites with different silicate loadings are presented in Fig.. As shown in Fig., {1} peaks for NBROCA, NBROC5A, and NBROC3B were observed at h =.3 (3.8 nm),.6 (3.4 nm), and. (4. nm), respectively, while {1} peaks for OCA, OC5A, and OC3B were observed at h = 3.6 (.5 nm), 4.3 (. nm), and 4.7 (1.9 nm), respectively. These shift of {1} peaks to the lower diffraction-angle side reveals that gallery heights of all nanocomposites finitely increase by the insertion of a θ =.3 ο (3.8 nm) {1} θ = 4.7 ο (1.9 nm) {} b θ =.6 ο (3.4 nm) {1} θ = 5.3 ο (1.7 nm) {} 1 wt% 1 wt% 8 wt% 8 wt% Intensity (a.u.) θ (degrees) c θ =. ο (4. nm) 7 wt% 5 wt% 3 wt% 1 wt% OCA Intensity (a.u.) 1 wt% 8 wt% θ (degrees) 7 wt% 5 wt% 3 wt% 1 wt% OC5A Intensity (a.u.) 7 wt% 5 wt% 3 wt% 1 wt% OC3B θ (degrees) Fig.. X-ray diffraction profiles for (a) OCA and NBROCAs, (b) OC5A and NBROC5A s, and (c) OC3B and NBROC3Bs.

5 J.W. Chung et al. / European Polymer Journal 45 (9) a.16 s s s -1 b.16 s s s -1 NBROCA1 NBROC5A1 Modulus, Modulus/Γ(α) NBROCA8 NBROCA7 NBROCA5 Modulus, Modulus/Γ(α) NBROC5A8 NBROC5A7 NBROC5A5 NBROCA3 NBROC5A3 NBROCA1 NBR Reduced Time, α t c Modulus, Modulus/Γ(α).16 s s s -1 NBROC3B1 NBROC3B8 NBROC3B7 NBROC3B5 NBROC5A1 NBR Reduced Time, α t NBROC3B3 NBROC3B1 NBR Reduced Time, α t Fig. 3. Modulus master curves applying STC to the tensile test data of (a) NBROCAs, (b) NBROC5As, and (c) NBROC3Bs.

6 84 J.W. Chung et al. / European Polymer Journal 45 (9) NBR chain into gallery spacing with maintaining the regular order of silicate layer, indicating the intercalated structure. In particular, the considerable increase of the gallery height occurred in NBROC3B compared to other nanocomposites; the gallery heights of the NBROCA, NBRO- C5A, and NBROC3B can be seen to increase by 1.3, 1.4, and.1 nm with respect to those of OCA, OC5A, and OC3B, respectively. Besides, a relative increase in full-width-at-half-maximum (fwhm) of the {1} reflection was observed for NBROC3B compared to NBROCA and NBROC5A, which demonstrates a decrease in the degree of coherent silicate layer stacking [4], and {} peaks (h = 4.7 and 5.3, respectively) as well as the {1} peaks were observed for NBROCA and NBROC5A, manifesting the preservation of long-range order that commonly appears in ordered intercalation system [5,4 6]. Therefore, it is believed that silicate layer of NBROC3B was more dispersed than that of other NBR nanocomposites, attributed to the intermolecular interaction between the polar groups of NBR and intercalant. However, the positions of the {1} peak for all nanocomposites were essentially unaffected by the increase in silicate loading, indicating that the interlayer gallery height for all nanocomposites were independent of the silicate loading. This means that further information about structural changes in the nanocomposites as a function of silicate content cannot be obtained from WXRD. Table 1 n Values for neat NBR and NBR nanocomposites. Silicate contents (wt%) Over 8 Neat NBR NBROCA Not superposed NBROC5A Not superposed NBROC3B Not superposed a Stress (MPa) (7) (6) (5) (4) (3) (1) neat NBR () NBROCA1 (3) NBROCA3 (4) NBROCA5 (5) NBROCA7 (6) NBROCA8 (7) NBROCA Strain c 7 () (1) 4.. Strain time correspondence (STC) Fig. 3 shows the results applying the tensile test data of neat NBR and NBR nanocomposites to STC. As shown in Fig. 3, it was confirmed that master curves of neat NBR and the NBR nanocomposites can be generated by applying tensile data acquired from three different strain rates to STC, excepting the NBR nanocomposites with silicate loading of 8 wt% and 1 wt%. For neat NBR, the master curve b Stress (MPa) (7) (6) (5) (4) (3) (1) neat NBR () NBROC5A1 (3) NBROC5A3 (4) NBROC5A5 (5) NBROC5A7 (6) NBROC5A8 (7) NBROC5A1 () (1) Strain Stress (MPa) (7) (6) (5) (4) (1) neat NBR () NBROC3B1 (3) NBROC3B3 (4) NBROC3B5 (5) NBROC3B7 (6) NBROC3B8 (7) NBROC3B1 1 (3) () (1) Strain Fig. 4. Tensile strain stress curves of (a) NBROCA s, (b) NBROC5A s, and (c) NBROC3B s obtained at a strain rate of.16 s 1.

7 J.W. Chung et al. / European Polymer Journal 45 (9) can be constructed by using the horizontal shift factor (extension ratio, a) only, indicating that the material is structurally homogeneous. However, for all of the nanocomposites (excepting the NBR nanocomposites with silicate loadings over 8 wt%), the master curves were formed by using the vertical shift factor (modulus shift factor, C(a)) as well as the horizontal shift factor. This reveals that the nanocomposites are structurally heterogeneous, which is due to the presence of domains such as the silicate tactoids by the physical jamming between silicate layers [6]. In particular, C(a) is known as a simple function that is dependent upon a: C(a)=a n [1,]. Hence, the degree of the modulus shift at a specific a is actually expressed by the n value, which can be considered to be a measure of the structural heterogeneity. Table 1 shows the n values obtained by applying the modulus shift factor to the initially generated double logarithmic plots of tensile modulus against horizontally reduced time for neat NBR and NBR nanocomposites. As shown in Table 1, n value for neat NBR is zero, because a STC master curve was constructed by only the horizontal shift factor. However, all of the NBR nanocomposites show the various n values with silicate loading and intercalant, excepting for the NBR nanocomposites with silicate loading over 8 wt%. In this case, it was found that NBROC3B containing a polar intercalant had the lowest n value in the NBR nanocomposites, indicating relatively lower structural heterogeneity than other NBR nanocomposite. This may be because the gallery height for NBROC3B more increased than other nanocomposites by the polar group of intercalant in NBROC3B, as shown in WXRD data; the increase of gallery height may somewhat loosen the silicate tactoid structure and weaken the heterogeneous structure of the nanocomposite formed by the structural boundary between matrix polymer and silicate tactoid. In addition, Giannelis et al. reported that the relaxation behavior of the polymer chains for the exfoliated nanocomposite are not altered by the presence of the silicate layers and displays similar behavior as the pure polymer, at least within the sensitivity of the dynamic mechanical measurements [7]. Besides, recognizing the modulus shift (b T ) as the measures of structural heterogeneity of material in time temperature superposition (TTS) [8] and the similarity between the effect of thermal activation and that of mechanical activation on the material time (i.e., a similarity between the TTS and STC) [15], it can be regarded that the silicate layer in the nanocomposite with an n value close to that of neat NBR is more dispersed. Therefore, it is considered that the silicate in NBROC3B with the n value close to that of neat NBR is relatively more dispersed than that in NBROCA and NBRO- C5A, which was consistent with WXRD results. Interestingly, for silicate loadings over 8 wt%, the nanocomposite did not fall onto a master curve irrespective of the type of organoclay, although both horizontal and modulus shift factors were used. In addition, the strengths and toughness of all of the nanocomposites with 8 wt% silicate loading show a greater strength enhancement than would be expected based on the strain stress data in Fig. 4. This may indicate that the silicate in the nanocomposites undergo an abrupt structural change over a silicate loading of 8 wt%. According to Lee et al. [9], such remarkable enhancement of mechanical properties can arise to the nanocomposites over a certain critical silicate loading due to the percolation of silicate. Krishnamoorti et al. [3] and Hyun et al. [31] also reported a similar observation and they explained it as the formation of three-dimensional network structure resulting from percolation between silicate tactoids over critical silicate loadings. Thus, it is considered that network-like percolated structure may be developed in the NBR nanocomposites with silicate loadings over 8 wt% (percolation threshold = 8 wt%) as shown in Fig HR-TEM To verify the change of silicate structure predicted from applying of STC, we carried out the morphological analysis by means of HR-TEM. Fig. 6 is HR-TEM images of (a) NBROC3B5 and (b) NBROC3B1. From Fig. 6, it was confirmed that both had the intercalated structure, corresponding to WXRD results. Moreover, it was found that network structure composed of silicate factoids was formed in NBROC3B1, whereas no network-like structure was observed in NBROC3B5. This is consistent with Fig. 5. Schematic model of network-like percolated silicate structure.

8 86 J.W. Chung et al. / European Polymer Journal 45 (9) Fig. 6. HR-TEM images of (a) NBROC3B5 and (b) NBROC3B1. the results suggested from STC principle and the schematic model in Fig. 5. Therefore, it is considered that STC is the useful and new tools to characterize the silicate structure in NBR nanocomposites. 5. Conclusions In this study, NBR nanocomposites containing OCs with the polar intercalant, nonpolar intercalant with short chain length, and nonpolar intercalant with long chain length were prepared with various silicate loadings by the melt mixing process. From XRD analysis, we found that the NBR chains were successfully melt intercalated within the silicate layers and the increasing rate of gallery height of the NBR nanocomposites increased in the order NBRO- CA < NBROC5A < NBROC3B. However, the interlayer gallery spacing of NBR nanocomposites was independent of the silicate loading. Hence, tensile strain stress measurements were carried out at various strain rates, and then the results of tensile measurement applied to the STC principle in order to further analyze the silicate structure in NBR nanocomposites. In this case, the modulus master curve for the NBR nanocomposites were constructed by the application of horizontal and vertical shift factor, indicating the heterogeneous structure ascribed to the presence of silicate tactoids. From comparison of n value for the nanocomposites obtained from vertical shift, it was found that the NBROC3B had the lowest n value. This implies that NBROC3B has more dispersed silicate structure than other nanocomposite, being consistent with WXRD. Interestingly, the STC principle was not applicable to the nanocomposites with silicate loadings over 8 wt% in spite of applying both the horizontal and vertical shift factors. Moreover, their strength enhancements were superior to expected strength enhancement. These results can be explained as an abrupt structural change of silicate structure such as a network-like percolation of silicate tactoids over a silicate loading of 8 wt%, which was in good agreement with HR-TEM results. Thus, it is thought that STC is a useful tool to characterize the silicate structure in NBR nanocomposites and can provide the clue for investigating the silicate structure of other rubber-based nanocomposites. Acknowledgement The authors would like to thank the Korea Science and Engineering Foundation (KOSEF) for sponsoring this research through the SRC/ERC Program of MOST/KOSEF (R ). References [1] Ray SS, Okamoto K, Okamoto M. Structure property relationship in biodegradable poly(butylene succinate)/layered silicate nanocomposites. Macromolecules 3;36(7): [] Akelah A, Salah El-Deen N, Hiltner A, Baer E, Moet A. Organophilic rubber-montmorillonite nanocomposites. Mater Lett 1995;(1 ): [3] Wang MS, Pinnavaia TJ. Clay-reinforced epoxy nanocomposites. Chem Mater 1994;6(1):16 9. [4] Pinnavaia TJ. Intercalated clay catalysts. Science 1983;(4595): [5] Giannelis EP, Krishnamoorti R, Manias E. Polymer silicate nanocomposites: model systems for confined polymers and polymer brushes. Adv Polym Sci 1999;138: [6] Jinnai H, Smalley MV, Hasimoto T, Koizumi S. Neutron scattering study of vermiculite poly(vinyl methyl ether) mixtures. Langmuir 1996;1(5): [7] Nakajima N, Collins EA. Deformational behavior of rubber in a variable speed Mooney rheometer. Rubber Chem Technol 1974;47(): [8] Nakajima N, Collins EA, Kumler PR. Dynamic viscoelastic properties of raw butadiene acrylonitrile elastomers. Rubber Chem Technol 1974;47(4): [9] Nakajima N, Collins EA. A master curve for amorphous elastomers derived from small and large deformations using various instruments. Rubber Chem Technol 1975;48(1): [1] Nakajima N. Quantitative characterization of gels in elastomers through temperature dependence and strain dependence of

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