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Microstructure-properties correlations in dynamically vulcanized nanocomposite thermoplastic elastomers based on PP/EPDM.


Thermoplastic elastomers (TPEs) constitute a unique class of polymers that offers many of the positive attributes of vulcanized rubbers, such as low compression set and high flexibility, while being processed using conventional thermoplastic fabrication techniques, without the time-consuming cure step, usually required with vulcanizates. A special type of TPEs is obtained when the rubber is dynamically vulcanized during its blending process to give rise to thermoplastic vulcanizates (TPVs). The microstructure of TPVs consists of a thermoplastic matrix (continuous phase) that contains a dispersed cured rubber phase [1, 2]. Many commercial TPVs have been developed for various applications in automotive parts, cable insulation, footwear, packaging, and medical industries because of their excellent weather-ability, low density, and relatively low cost [3-5]. The addition of fillers, such as glass fiber, carbon black, talc, and calcium carbonate to the TPV blends at high loading levels improve properties like stiffness, heat distortion temperature, and dimensional stability. However, these fillers also increase the weight of the TPV blends, rending them less attractive for automotive and aerospace applications [6, 7]. On the other hand, similar benefits but without the weight penalty could potentially be obtained by using a much smaller amount of so-called nanoclays [8], owing to the nanoscale dimension of silicate layers, which brings stronger interfacial interactions between the dispersed solids and the polymer matrix [9].

Montmorillonite ([Na.sup.+]-MMT) is a layered mica-type clay mineral whose crystal lattice consists of a central octahedral sheet of alumina/magnesium sandwiched between two fused tetrahedral sheets. Montmorillonite has negative charges on the interlayer gallery walls and the galleries are normally occupied by cations such as [Na.sup.+], [Ca.sup.2+], and [Mg.sup.2+]. The silicate layers are often modified to organomontmorillonite (OMMT) by an aliphatic alkylammonium or alkyl phosphonium ion-exchange reaction in order to improve its compatibility with nonpolar polymers and thus facilitate its dispersion [10, 11]. The structure of nanocomposites depends on the structure of dispersed clay particles in the polymer matrix, which are usually classified as intercalated or exfoliated. In the intercalated state, the clay is dispersed as lamellar structures and polymer chains penetrate between the silicate layers of clay. The exfoliated structure characterizes itself by a complete delamination of clay particles and individual silicate layers are dispersed throughout the polymer matrix.

Nanocomposites based on polypropylene [12-16] and thermoplastic elastomers based on PP/EPDM [17-19] have been prepared using a traditional melt mixing process and their properties and microstructures have been published. These studies revealed that clay intercalation or exfoliation can only be achieved when clay/matrix compatibility and suitable processing conditions are found. Lee and Goettler [20] studied TPV/clay nanocomposites prepared by melt mixing a commercial TPV with several types of clay, but without the use of a compatibilizing agent for the polyolefins. They reported that the mechanical properties of the TPV nanocomposites increased when the miscibility of the clay and polymer matrix increased. Although there have been a few other studies on the preparation and properties of thermoplastic elastomer (TPE) nanocomposites [21], the area of dynamically vulcanized nanocomposite TPEs prepared using polypropylene of different viscosities, maleic anhydride-grafted polypropylene, and sulfur as curing agent has remained mostly uncharted. The main objective of this project was to obtain dynamically-vulcanized nanocomposite thermoplastic elastomers based on PP/EPDM and to study the parameters affecting their microstructure and properties. For that purpose, we studied dynamically-vulcanized nanocomposite TPEs while investigating the effect of PP/EPDM viscosity ratio and clay concentration on their microstructure and mechanical properties. The evolution of the nanocomposite microstructure during the mixing process was also monitored.



The basic specifications of EPDM and the three grades of polypropylene (Basell Co.) employed in this study are reported in Table 1. The EPDM rubber (Buna EP T 6470) was based on ethylene norbornene (ENB) as a termonomer, and was supplied by Bayer Co. The clay used was Cloisite 15A from Southern Clay Products, a natural montmorillonite modified with a dimethyl dehydrogenated tallow quaternary ammonium having a cation-exchange capacity (CEC) of 125 meq/100 g and a specific gravity of 1.66, for an average particle (bulk agglomerate) diameter of 8 [micro]m. The maleic anhydride-modified polypropylene (PPMA) Epolene G3015 (acid number = 15 mg of KOH/g PPMA, MW = 47,000 g mol, MA = 1.3%) was obtained from Eastman Chemical. The curing of TPV nanocomposites was achieved with benzothiazyl disulfide (MBTS), tetramethyl thiuram disulfide (TMTD), and sulfur (S) obtained from Bayer Company.

Sample Preparation

TPV nanocomposite samples were prepared by a three-step melt mixing process in a Haake Rhemix 600 internal mixer. In the first step of the process, the clay powder and PPMA pellets were dry-mixed in a bag to ensure that the materials were well dispersed at the macroscale level. The mixture was then melt-mixed at 180[degrees]C under nitrogen to obtain a master batch using the mixer at a rotor speed of 150 rpm for 10 min. For all master batches, the PPMA/clay weight ratio was kept constant at 3:1. In the second step, the master batch was then dry-blended with polypropylene to give the desired composition (as shown in Table 2). and fed to the mixer by applying the same previous processing conditions. In the third and final step, PP nanocomposite (from step 2), EPDM, zinc oxide (ZnO), and stearic acid (St. Ac.) were mixed at 180[degrees]C and 100 rpm Three to four minutes after the melting of the PP nanocomposite, the TMTD and MBTS curing agents were added to the mixer and after 60 s of mixing, the sulfur (S) was introduced and followed by a continuous mixing for 4 min. Prepared compounds were compression-molded at 200 C for 10 min to obtain suitable samples for testing. The ratio of EPDM/PP was kept constant at 60:40 (wt%) and the vulcanizing system was based on 100 phr EPDM, 5 phr ZnO, 1 phr stearic acid, 1 phr TMTD, 0.3 phr MBTS, and 1 phr S. For comparison purposes, unfilled but otherwise similar TPVs were also compounded as reference materials. In all instances, polypropylenes and clays had been dried at 80[degrees]C for 24 h before processing.


To evaluate the clay dispersion in the polymer matrix, X-ray diffraction (XRD) analysis were performed at room temperature using a X-ray diffractomer (Philips model X'Pert) in the low angle of 2[theta] (SAXD). The X-ray beam was a Cu K[alpha] radiation ([lambda] = 1.540598 [Angstrom]) obtained from a 50 kV voltage generator and a 40 mA current. The basal spacing of silicates was estimated from the position of the plane peak in the SAXD intensity profile using the Bragg's law, d = [lambda]/(2 sin[[theta].sub.max]). Specimens for X-ray diffraction were taken from compression-molded sheets of 2 mm in thickness. The nanostructure of the clay was observed by a transmission electron microscopy (JEM-2100F, JEOL) with an accelerator voltage of 200 kV. The surface of the samples was first coated with gold and then a thin section of each specimen was prepared by using a focused ion beam (FB-2000A, Hitachi) and also a cryogenic ultramicrotome at -100[degrees]C. To study the morphology of the TPV nanocomposite samples, cryogenically fractured surfaces of the samples were etched by hot xylene. Treated samples were then coated with gold and viewed with a JEOL scanning electron microscope JSM-840. The rheological characterization of polypropylenes and EPDM were carried out using a stress-controlled rheometer (SR 5000). The experiments were performed in 25 mm parallel-plate geometry under a nitrogen atmosphere at a temperature of 220[degrees]C and in the frequency range of 0.01-80 Hz. The stress-strain properties of the composites were determined in accordance with the test procedures set forth in ASTM D412 using an Instron model 4201 and a crosshead speed of 50 mm/min. Crystallization was studied using a Pyrist 1 (Perkin Elmer) differential scanning calorimeter under nitrogen atmosphere. Samples were heated from 50 to 200[degrees]C at a rate of 10[degrees]C/min and then they were cooled from 200 to 50[degrees]C at the same rate after holding at 200[degrees]C for 5 min to erase any thermal history effects. Oxygen permeability was measured by using an OX-TRAN model 2/21, Mocon Co. The 1-mm thick sample specimens were prepared by pressing the TPVs in standard mold at 200[degrees]C. The permeability was evaluated as the volume of oxygen permeated per unit area of TPV sheet for 24 h.



Figure 1 shows the elastic moduli (G') and complex viscosities ([eta]*) of EPDM and polypropylenes with different viscosities (P1, P2, and P3) as a function of angular frequency. Because it was observed that the pure materials rheology follows the Cox-Merz relation, ([eta]*([omega] = [gamma]) = [eta]([gamma])) [22], the simplified Carreau-Yasuda model was used to fit the complex viscosity data of the pure polymers from oscillatory shear measurements. The Carreau-Yasuda (C.Y.) model is given by [23]:

[eta]*([omega]) = [[eta].sub.0]/[(1 + ([omega][tau])[.sup.a])[.sup.1-n/a]]. (1)

Where [eta]* is complex viscosity, [[eta].sub.0] is the zero shear viscosity, a is a dimensionless parameter which specifies the transition from Newtonian to shear thinning behavior, [tau] is in general close to the average of the melt relaxation time, and n is the so-called power law index. The zero shear viscosity of the three different PPs, as reported in Table 1, was estimated by fitting [E.sub.q]. 1 to experimental data. The evolution of torque during the compounding of a multiphase reactive formulation often brings evidences of the time dependent mixture state. Figure 2 shows torque traces for an unfilled TPV based on PP21/EPDM (40/60 w/w) and a TPV nanocomposite prepared from NP21/EPDM (40/60 w/w). These traces were obtained during steps 2 and 3 of the sample preparation process described in the earlier section. As shown in this figure, a higher steady-state value of the mixing torque is obtained for a TPV nanocomposite prepared from NP21/EPDM, in comparison with that of the unfilled TPV. We have noticed that for formulations with similar clay contents (not shown in Fig. 2), a simple indicator of the extent of clay exfoliation within the TPV matrix is the steady-state torque recorded during melt mixing. An increased torque value is an indirect evidence of stronger interactions between the TPV matrix and the clay [24].



The structure and morphology of the nanocomposites were characterized by using X-ray diffraction analysis and TEM imaging of samples taken after the observed maximum in the mixing torque, as discussed previously. In preliminary experiments assessing the effect of PPMA content, low viscosity polypropylene (P1) was used to prepare 2% (w/w) clay nanocomposites with PPMA/clay weight ratios of 0, 1, 2, and 3. The X-ray diffraction results of corresponding samples are shown in Fig. 3. It can be seen from curves (a) and (b) that the diffraction peak corresponding to the (001) plane of clay appears at 2.9[degrees], while the diffraction peak of the uncompatibilized system (PP/PPMA/clay, 98/0/2) appears at a higher angle (2[theta] = 3[degrees], d = 29.42 [Angstrom]). These results reveal that propylene cannot intercalate between the layers of clay even modified by a dimethyl dehydrogenated tallow quaternary ammonium [25]. On the contrary, the XRD peaks of the samples prepared using master batch with different ratios of PPMA/clay show an increase in interlayer spacing as a result of the intercalation of polypropylene. The interlayer distance of the clay (d = 33.31 [Angstrom]) in the nanocomposite prepared with a PPMA/clay ratio of 2 (PP/PPMA/clay, 94/4/2) is higher than that of the sample with PP/PPMA/clay (96/2/2), (d = 31.3 [Angstrom]) as shown in curves (c,) and (d). From the results in (e), it is shown that the position of the clay characteristic XRD peak of the sample with PPMA/clay = 3 remains relatively constant. However, an important difference is observed in the decrease of the peak intensity, in comparison with the similar nanocomposite based on a PPMA/clay ratio of 2. These results are in consistent with the recent work of Vergnes [25]. The PPMA/clay ratio was therefore maintained constant at 3 for this work, but the role of PPMA concentration on nanocomposites morphology would need to be further investigated in the future. The concentration range and type of PPMA used for this work were similar to the conditions reported in Refs. 25 and 26 and therefore, miscibility of the polymers was assumed.

The X-ray diffraction results of clay itself, PPMA/clay = 3 master batch, and nanocomposites based on 2 wt% of clay, respectively, prepared with polypropylene of three different viscosities (P1, P2, and P3) are shown in Fig. 4. It is observed that the diffraction peak corresponding to the (001) plane of clay appears at 2.9[degrees], while the diffraction peak of the master batch appears at a lower angle (2[theta] = 2.75[degrees], d = 32.10 [Angstrom]) than that of clay, which indicates an increase in the interlayer distance of the silicates. It has been reported that PPMA can penetrate into the silicate layers of clay during melt mixing and induces an expansion of the gallery distance [27]. The interlayer distance of the nanocomposite based on P1 with 2 wt% of clay is about 33.65 [Angstrom], larger than the original interlayer spacing of Cloisite 15A and the master batch, which implies that the polypropylene can intercalate into the clay. The interlayer distance of P3 nanocomposite for the 2 wt% clay formulation (d = 32.25 [Angstrom]) is less than that of the similar nanocomposite based on P1 with the same clay content. It should also be noted that the intensity of the diffraction peak in the nanocomposite based on P1 decreases and also the peak is broader than its counterpart in nanocomposites based on P2 and P3 of identical clay content. The decrease in intensity and the broadening of peaks indicate that the stacks of layered silicates become more intercalated or partially exfoliated [24]; hence, the nanocomposite obtained from low viscosity polypropylene yields a more disordered structure (better clay dispersion). This is explained by the fact that polypropylene molecules with high MFI (low viscosity) need a lesser level of attractive Flory-Huggins interactions to penetrate between the silicate layers in comparison with high viscosity polypropylene products [16]. This entropic effect is found to be more important than the increased stress level imparted to the clay particles by polypropylenes of higher viscosity [16]. In Fig. 5, we present the XRD patterns of PP nanocomposites based on 5 wt% of clay, respectively, prepared with polypropylene of three different viscosities (P1, P2, and P3). It can be seen that the intensities of the peaks increased when compared with the nanocomposites prepared with 2 wt% of clay. This is because of the influence of the packing density, rending more difficult the penetration of the polymer chains between the silicate layers. X-ray diffraction patterns of the clay and TPV nanocomposite prepared based on NP15/EPDM (40/60 w/w) are presented in Fig. 6 at different mixing stages (A, B, C as shown in Fig. 2). The diffraction peak corresponding to the (001) plane of the nanocomposite with activators (stage A) appears at 2.5[degrees] (d = 35.35 [Angstrom]), while the nanocomposite with activators and accelerators (stage B) shows a peak at a lower angle (2[theta] = 2.2[degrees], d = 40.12 [Angstrom]) as shown in Fig. 6b and c, respectively. The peak [d.sub.001] is further shifted toward lower angles (2[theta] = 1.9[degrees], d = 46.31 [Angstrom]) after the addition of the sulfur curing system (stage C). It is also seen that the intensities of the peaks decreased. Similar trends were observed for TPV nanocomposites based on NP25/EPDM and NP35/EPDM (40/60 w/w). This can be explained by the fact that the dynamic vulcanization of EPDM phase increased the viscosity of the blend and hence the shear stress imposed by the matrix during the mixing is also increased, which facilitates the break-up process of clay agglomerates.





Figure 7 shows the effect of postprocessing (compression molding) on the X-ray diffraction patterns of TPV nanocomposites based on NP15/EPDM, NP25/EPDM, and NP35/EPDM (40/60). It can be seen that the first characteristic peak of the clay disappeared for all TPV nanocomposites (PP3, PP2, and PP1), respectively, when postprocessing is applied to these samples as shown in Fig. 7b-d. The absence of the diffraction peak suggests that the compression molding process could contribute to the further exfoliation of silicate layers in the TPV nanocomposites. It should be noted that X-ray diffraction pattern of the clay shows two peaks, the second peak appeared at 2[theta] = 7.1[degrees], which may be because of the diffraction from a second silicate layer ([d.sub.002]), because 2[theta] of the second peak is about twice the [d.sub.001] peak of the clay. The second peak of clay may also come from the portion of the clay that is not modified [28]. On the other hand, for the X-ray diffraction patterns of the TPV nanocomposites, the second peak of the clay appears at lower 2[theta] (6.7[degrees]).

To get a better insight of the clay dispersion, TEM images were taken from these PP nanocomposites and TPVs as illustrated in Figs. 8-10. The TEM image of the PP nanocomposites prepared by using a low viscosity polypropylene (P1) shows silicates agglomeration at 2 and 5 wt% of clay as shown in Fig. 8a and b, respectively. It can be seen that the size of the stacks in creased at higher clay content. We also observed the agglomeration of the clay for the nanocomposite prepared from high viscosity polypropylene (P3) at 5 wt% of clay as shown in Fig. 8c. A TEM image of TPV nanocomposite prepared from low viscosity polypropylene (P1) before the curing stage (stage A) is shown in Fig. 9a. It indicates that the silicate layers are dispersed through the PP matrix and cannot penetrate into the EPDM phase during the mixing process. The rubber phase is white in color, and silicate layers appear as black lines in the darker polypropylene phase. Also at stage A of the mixing process, in TPV nanocomposites based on high viscosity PP, the clays were found to be dispersed in the continuous phase but partly agglomerated at the interfacial boundary between the polypropylene and rubber phases, as shown in Fig. 9b. TEM images of TPV nanocomposites prepared from NP15/EPDM (40/60), and NP35/EPDM (40/60) are presented in Fig. 10a and b, respectively. For both compositions, XRD results suggested a nearly completely exfoliated structure. For the low viscosity polypropylene formulation, it can be seen that a significant fraction of the silicate layers are exfoliated in the PP matrix, while a slight amount remained in a more clustered state. The latter is more pronounced for the high viscosity PP-based TPV (Fig. 10b). The existence of the small peak around 2[theta] = 7[degrees] in the X-ray diffraction patterns should therefore be associated to these unexfoliated silicates.



Further insights on the morphology of these materials are found in Fig. 11, where SEM images of cryogenically fractured and etched TPV samples based on P2/EPDM with and without clay are reported. Despite the higher concentration of the rubber phase, it is clearly seen that the rubber particles are dispersed throughout the polypropylene in the form of aggregates and the size of rubber particles is less than 2 [mu] in the unfilled TPV sample of Fig. 11a. This behavior is attributed to the viscoelasticity of the rubber phase, as reported by Wu [29]. It should be noted that the crosslinked rubber particles are covered by a layer of polypropylene and attached together via a joint shell mechanism [18]. As shown in Fig. 11d, the solvent has not been able to wash out these shells, but they developed a porous structure after the extraction process. It can also be observed in Fig. 11b-d that the size of rubber particles has been increased by the introduction of the clay. The increase in the rubber phase of the TPV nanocomposites is explained by the modified rheology of the nanocomposites. It is accepted that the size of the dispersed rubber phase in unfilled PP/EPDM TPV materials based on PP/EPDM (without clay) depends on the viscosity ratio and interfacial interactions between two phases (Goharpey et al. 2003, [18]). In the case of the TPV nanocomposites, the clay plays an important role to determine the morphology of the samples. Initially, the nanoclay cannot penetrate into the rubber phase, but after adding curing accelerators such as TMTD, the rubber phase become more polar [30]. Therefore, it is possible that some nanoclay goes to EPDM phase before the completion of the curing cycle. This changes the viscosity ratio between the two phases, and consequently, the size of the rubber phase increases. Figure 10a and b also confirm the presence of some nanoclay in EPDM phase.

In a further characterization of these materials, mechanical properties and oxygen transmission rate for the TPVs and TPV nanocomposites have been obtained and are reported in Table 3. It is observed that the tensile strength and modulus of the TPV reference samples without clay increases when the viscosity of polypropylene phase increases. This is due to the fact that the decrease in the viscosity difference between the two phases leads to the formation of smaller rubber particles with higher specific surface area, therefore, the level of interactions between two phases (PP and EPDM) improves [18]. In TPV nanocomposites, the tensile modulus of samples increased from 20 to 90% depending on the clay content and viscosity ratio of PP/EPDM. For example, the exfoliation of 2 wt% clay in a low viscosity PP-based TPV nanocomposite (NP15/EPDM, 40/60) increased the tension modulus by 92%, while the similar addition of 2 wt% clay to a TPV based on high viscosity PP (NP35/EPDM, 40/60) resulted in only a 23% increase in tension modulus over unfilled TPV. This would tend to confirm that the clay dispersion and exfoliation is better achieved in the former composition. It is also observed that the tensile strength of the TPV nanocomposites decreases with increasing nanofiller content as reported in Table 3. The improvement in the mechanical properties of a filled system not only depends on the filler type and extent, but also on the crystallinity of polymer matrix [24]. Figure 12 shows the crystallinity and crystallization temperature of polypropylene nanocomposites prepared with different melt flow index and different nanofiller loading levels. It is seen that the degree of crystallinity of the nanocomposites decreases with the increasing clay content as shown in Fig. 12a. This decrease is because of the interaction between the polymer matrix and the clay, which reduces the mobility of crystallizable chain segments [31]. We noted that the reduction in crystallinity of the nanocomposites correlates well with the decrease in tensile strength. As can be seen in Fig. 12b, the peak crystallization temperatures ([T.sub.c]) of all compositions increased with the clay content because the nanoscale particles acted as nucleating agents, also causing a reduction of the size of the crystalline domains [32]. Another important feature of the TPV nanocomposites is their barrier properties such as oxygen permeability. The oxygen permeability results of the pristine TPVs and TPV nanocomposites are reported in Table 3. The oxygen permeability of all TPV nanocomposites decreased with the introduction of the clay, indicating an improved gas-barrier property of the TPV nanocomposites. In particular, the TPV nanocomposites prepared with only 2% (w/w) of clay, the oxygen permeability decreased by as much as 20% when compared with the pristine TPV.





TPV nanocomposites have been prepared in a melt compounding process by dynamically vulcanizing thermoplastic elastomers based on PP/EPDM. The compounding process itself was found to be instrumental in the final dispersion state of the layered silicates. The interlayer distance is first increased by the incorporation of the filler in the PPMA to obtain a master batch. This preliminary blend is then compounded with the PP, where the PPMA fraction acts as a compatibilizer for the nonpolar PP. Further intercalation and partial exfoliation are finally achieved by the shear stress developed during the vulcanization of EPDM due to the increased viscosity. The microstructure of the TPV nanocomposites was found to be sensitive to the viscosity ratio of PP/EPDM and clay content. In the TPV nanocomposites prepared with low viscosity polypropylene, an almost complete exfoliation and random distribution of clay in the thermoplastic phase was observed by X-ray diffraction analysis and TEM images. In all TPVs, the rubber particles were dispersed through the polypropylene in the form of agglomerates and their size increased with the introduction of clay as evidenced by the morphological features of the dynamically vulcanized nanocomposite samples. In the TPV nanocomposite based on low viscosity polypropylene, tensile modulus increased by 90% at 2 wt% of clay. The exfoliated structure resulted in a reduction of the degree of crystallinity and an increase of the polymer crystallization temperature because the dispersed clay silicates acted as nucleating agents. Moreover, the oxygen permeability of the TPV nanocomposite samples prepared from polypropylene of different viscosities decreased in comparison with the pristine TPVs.


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Ghasem Naderi, Pierre G. Lafleur, Charles Dubois

Center for Applied Research on Polymers and Composites, CREPEC, Department of Chemical Engineering, Ecole Polytechnique de Montreal, Montreal, Quebec, Canada

Correspondence to: Charles Dubois; e-mail:
TABLE 1. PP and EPDM specifications.

 MFI 14 g/10 min
 Density 0.91 g/[cm.sup.3]
 [T.sub.m] 165[degrees]C
 [[eta].sub.0] 1818 Pa s
PP (P2)
 MFI 4 g/10 min
 [[eta].sub.0] (Pa s) 5000 Pa s
PP (P3)
 MFI 0.5 g/10 min
 [[eta].sub.0] (Pa s) 27,100 Pa s
 Mooney viscosity ML (1 + 4), 125[degrees]C 55
 Ethylene content 68%
 Termonomer content 4.5% ENB
 Density 0.86 g/[cm.sup.3]

TABLE 2. Samples names and compositions.

Sample code PP PPMA Nanoclay

NP11 P1 = 96 3 1
NP12 P1 = 92 6 2
NP13 P1 = 88 9 3
NP15 P1 = 80 15 5
NP21 P2 = 96 3 1
NP22 P2 = 92 6 2
NP23 P2 = 88 9 3
NP25 P2 = 80 15 5
NP31 P3 = 96 3 1
NP32 P3 = 92 6 2
NP33 P3 = 88 9 3
NP35 P3 = 80 15 5
PP21 P2 = 97 3 -

Details about P1, P2, and P3 are found in Table 1.

TABLE 3. Mechanical properties and gas permeability of TPV

 Tensile Elongation Tensile Oxygen transmission rate
 modulus at break strength ([cm.sup.3]/[m.sup.2]
Sample code (MPa) (%) (MPa) day)

Pristine TPV P1 19.10 740.3 12.2 540
TPV NP11 25.15 694.3 11.1
TPV NP12 30.82 623.5 10.2
TPV NP13 34.12 532.7 9.3
TPV NP15 36.70 459.9 8.5 428
Pristine TPV P2 21.93 790.5 13.5 514
TPV NP21 25.77 760.2 12.8
TPV NP22 28.54 700.3 12.1
TPV NP23 31.14 645.9 11.4
TPV NP25 37.32 560.2 10.0 415
Pristine TPV P3 27.51 880.2 17.1 495
TPV NP31 27.91 817.1 16.7
TPV NP32 29.54 711.9 15.9
TPV NP33 30.98 660.1 14.8
TPV NP35 33.75 621.6 14.1 405
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Author:Naderi, Ghasem; Lafleur, Pierre G.; Dubois, Charles
Publication:Polymer Engineering and Science
Geographic Code:1USA
Date:Mar 1, 2007
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