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Effect of organoclay content on the rheology, morphology, and physical properties of polyolefin elastomers and their blends with polypropylene.


Blending of polypropylene with elastomers has become common practice, yielding compounds with diverse properties, ranging from rigid to soft, depending on the PP/elastomer ratio. The most commonly used elastomers in the production of these Thermoplastic Olefin (TPO) blends are ethylene-propylene rubber (EPR) [1, 2], ethylene-propylene-diene rubber (EPDM) [3, 4], and more recently, metallocene or single-site catalyzed ethylene-[alpha]-olefin copolymers [5-9]. Although the majority of preceding research deals with rigid TPOs, wherein polypropylene comprises the matrix phase, soft TPOs are also gaining importance in recent years, particularly because of their applications in automotive interiors [10].

The use of fillers, such as talc, calcium carbonate, and mica in TPO formulations is widespread, aiming at reducing production cost and improving properties such as stiffness, heat distortion temperature, and dimensional stability. Recent reports have highlighted the effects of addition of organoclay-type nanofillers in TPOs and their thermoplastic vulcanizate counterparts (TPVs) [11-17].

Lee and Goettler [11] studied a commercially available TPV-based nanocomposite and a TPO nanocomposite consisting of a physical blend of PP with EPDM. PP/polyethylene-octene copolymer elastomer nanocomposites were studied by Ma et al. and Maiti et al. [15-17]. In these studies the blend components were not functionalized. Other publications have focused on the addition of organoclays on commercially available TPOs, using maleated PP as a compatibilizer. Mehta et al. [14] highlighted the effect of organoclay on the morphology and mechanical properties of a commercially available TPO. They observed that clay resided in the PP matrix and affected the morphology of the blend. Mishra et al. [13] presented a study on soft TPOs, where the elastomer comprised the major phase. In a recent paper, Maiti et al. [17] reported that the improvements in mechanical properties were optimum when nanoclays were added in the elastomeric phase.

In these studies, the reported effectiveness of organoclay addition varies depending on the composition, type of elastomer, presence of compatibilizer, and method of preparation of the base TPO (through in-reactor polymerization or physical blending). Final properties seem to be influenced significantly by the location of the clay in one or both phases, as well as by the extent of clay exfoliation and/or intercalation.

In this paper, we investigate the effect of organoclay addition in maleated elastomer/PP thermoplastic olefin blends, wherein the maleated elastomer comprises the matrix phase. To provide a systematic approach in the analysis of the results, the structure and properties of the maleated elastomer/nanocomposite matrix have been fully characterized prior to analyzing the structure and physical properties of the TPO nanocomposites.



A montmorillonite clay, Nanomer[R] I.44PA, ion-exchanged with dimethyl-dialkylammonium halide (70% C18, 26% C16 and 4% C14), was supplied by Nanocor Inc., Arlington Heights, Illinois. Maleated ethylene-co-1-octene copolymer (EOC-g-MAn, Exxelor MDEX 95-2) containing approximately 0.35 wt% of pendant succinic anhydride groups and polypropylene (PP, Escorene PP1042) were supplied by ExxonMobil Chemical. A maleated ethylene-co-propylene copolymer (EPR-g-MAn, Fusabond[R] MF-416D) containing 0.5-1.0 wt% maleic anhydride was obtained from E.I. DuPont Canada. The properties of all materials are summarized in Table 1. Irganox B225 antioxidant was obtained from Ciba-Geigy. All materials were used as received.

Composite Preparation

Mixtures of polymer, clay (2-20 wt%) and Irganox B225 (500 ppm) were prepared by dry blending in a tumble mixer and were introduced subsequently to a Haake PolyLab torque rheometer, connected to a Rheomix 610p mixing chamber equipped with roller rotors. The instrument was operated at 60 rpm for 8 min, using a fill factor of 70%, and compounding temperatures of 170[degrees]C and 200[degrees]C for the elastomers and elastomer/PP blends, respectively. Blends containing elastomer and PP at a ratio of 70/30 were prepared by compounding the polymers and 2-20 wt% clay simultaneously. Unfilled polymers were similarly compounded as control materials for comparative evaluations of nanocomposite properties.

X-Ray Diffraction

The basal spacing of the clays in the composites was evaluated by X-ray diffraction (XRD) using a Scintag Model X1 powder X-ray diffractometer (Cu K[alpha] radiation, [lambda] = 1.5406 [Angstrom], generator voltage = 45 kV, current = 40 [micro]A). Films approximately 400 nm thick were prepared by compression molding and scanned in a 2[theta] range from 1[degrees] to 50[degrees] at a rate of 1[degrees]/min. Measurements were recorded at every 0.03[degrees] interval.

Transmission Electron Microscopy

A FEI Tecnai 20 transmission electron microscope (TEM), operated at 200 kV, was used to characterize the structure of all composites. Samples were sectioned with a diamond knife on a Leica ultracut cryo-ultramicrotome to approximately 70 nm thickness at -100[degrees]C and placed on formvar-coated copper grids. Microtomed sections from the nanocomposite blends were collected on carbon-coated copper grids and held in 0.5% Ru[O.sub.4] vapor for 1 h to enhance compositional contrast between phases. Images of the samples were obtained using a Gatan Dualview digital camera.

Scanning Electron Miscroscopy

Surfaces for SEM were prepared by fracturing samples under liquid nitrogen. The specimens were sputtered with gold and observed using a JEOL 840 scanning electron microscope. Image analysis was performed to determine the average diameter of the dispersed particles, using the Sigma-Scan Pro image analysis software.

Rheological Characterization

The elastic modulus (G'), loss modulus (G"), and complex viscosity ([eta]*) were measured as functions of angular frequency ([omega]) using a Reologica ViscoTech oscillatory rheometer. The rheometer was operated in the dynamic oscillatory mode with parallel plate fixtures 20 mm in diameter and at a gap of 1.5 mm. All tests were carried out under nitrogen blanket to limit the extent of polymer degradation and moisture absorption. Strain sweeps were performed at a constant frequency of 0.1 Hz to investigate the strain-dependent behavior of the composites. The testing temperatures were 170[degrees]C and 200[degrees]C for the elastomers and their blends respectively.

Steady-shear viscosity measurements were done using a Rosand RH2000 dual bore capillary rheometer by Bohlin instruments, at a temperature of 200[degrees]C. The Rabinowitch and Bagley corrections were applied to calculate steady-shear viscosity as a function of shear rate.

Thermal Properties

Melting, crystallization, and glass transition temperatures, as well as degrees of crystallinity were determined using a TA Instruments Q100 Series differential scanning calorimeter (DSC). Samples were heated from -90 to 200[degrees]C at a rate of 10[degrees]C/min under nitrogen. In all cases, a preliminary heating sequence was performed to remove thermal history effects.

A TA Instruments Q500 Series Thermogravimetric Analyzer (TGA) was used for thermogravimetric analysis. Samples were heated from room temperature to 900[degrees]C at a rate of 10[degrees]C/min under a nitrogen atmosphere. The onset of degradation was taken as the temperature where the rate of change of weight loss with respect to time increased beyond zero.

Mechanical Properties

An Instron Model 3369 Universal Tester, equipped with a 5-kN load cell, wedge grips, and Series IX software, with a crosshead speed of 50 mm/min at room temperature was used to determine tensile properties according to ASTM D638. Samples were prepared by compression molding sheets 1.5 mm thick at 170[degrees]C and 200[degrees]C for the elastomer and TPO samples respectively, from which test specimens were cut using a Type V die. Six measurements per sample were taken to optimize the precision of reported data.

Evaluation of Bound Polymer Content

Chopped EPR-g-MAn/clay composite (1 g) was dissolved in toluene (40 ml), at 80[degrees]C. After cooling, the solution was subjected to centrifugal separation at 4000 rpm for 1 h. The supernatant was decanted and the solids were shaken with 30 ml of fresh toluene before being left to stand for 2 h. Centrifugal separation was repeated and the remaining solids, containing clay and bound polymer were dried under vacuum at 60[degrees]C.

Bound polymer determinations were made using a TA Instruments Q500 Series thermogravimetric analysis instrument to heat samples under a nitrogen atmosphere from 25[degrees]C to 650[degrees]C at a rate of 10[degrees]C/min. Percent weight losses were recorded for clay, unextracted samples of EPR-g-MAn nanocomposites containing 5 and 20 wt% clay, and the solids extracted from this material. These values provided the information needed to calculate estimate the weight-percent of bound polymer within the extracted composite sample.

It should be noted that these tests were only done on EPR-g-MAn samples, because this polymer dissolves completely in toluene.


Polyolefin Elastomer Nanocomposites

Structure. Representative XRD traces recorded for the organoclay and the polyolefin elastomer-based composites are presented in Fig. 1. Note that prior to XRD analysis, the pure clay sample was heated to 170[degrees]C for 30 min and stored under nitrogen in order to reproduce the heat history and moisture content of the filler that was compounded into the various composites. The XRD profiles of the heat-treated and untreated clays were virtually identical, which suggests that intercalant degradation did not result from this process. Thermal stability studies on the clay through TGA analysis corroborated this finding.

The onium-modified clay exhibited a strong peak at a diffraction angle of 2[theta] = 2.84[degrees], which corresponds to an interlayer spacing of 32 [Angstrom] (3.2 nm) for the modified silicate. The absence of a peak for the EPR-g-MAn/organoclay composites at low clay loadings (Fig. 1a) suggests that the ordered clay structure has been virtually destroyed. At higher clay loadings, reemergence of the peak signifies the presence of substantial amounts of intercalated clay. For the EOC-g-Man-based composites (Fig. 1b), the characteristic diffraction peak of the mineral appeared reduced in intensity and shifted to lower angle of observance. Given that such a shift corresponds to an increase in interlayer distance between clay platelets to 40 [Angstrom] (4 nm), its occurrence is consistent with an intercalated clay composite structure.

Examination of the elastomer-based nanocomposites by TEM revealed the presence of delaminated clay platelets, alongside intercalated tactoids, suggesting the existence of hybrid structures (Figs. 2a-2c). This confirms that polyolefin elastomers have a high capacity for clay exfoliation, in agreement with previous reports on EPDM [18, 19] and polyethylene-co-octene nanocomposites [20].

Rheological Properties. In the presence of clay platelets, the onset of the nonlinear viscoelastic region shifted to lower strains, as the clay content increases (see Fig. 3). This phenomenon is the well-known Payne effect [21] in filled polymer systems and can be attributed to the presence of strain-sensitive rigid networklike structures that are more likely to break down at lower strains, resulting in the narrowing down of the region of linear viscoelasticity [22].

Several researchers have suggested that the networklike effects observed in the melt state in polymer nanocomposites may be due to the presence of interconnected nanoclay platelets or tactoids [23]. Filler/polymer interactions may also be responsible for such observations. Schmidt et al. [24] suggested that at high strains there is a change in the kinetics of the adsorption/desorption equilibrium between polymer and clay particles in viscoelastic clay/polymer solutions.


Although the topic of polymer/filler interactions has not received a great deal of attention in polymer/organoclay-based systems, it has been addressed extensively in elastomer filled systems, containing silica and carbon black. An excellent recent review of the current state-of-the-art understanding of these interactions is provided by Zhu et al. [25].

Upon examining the TEM images of the nanocomposites containing 5 wt% organoclay (see Fig. 2), we could not find evidence of extensive filler/filler contacts that would result in network formation. This is corroborated by the rheological properties shown in Fig. 4, where at clay loadings up to 5 wt% we saw typical viscoelastic-liquid behavior, accompanied by the presence of a Newtonian plateau and a tendency to approach the terminal flow region.



On the other hand bound polymer measurements in the EPR-g-Man-based composites revealed that in a composite containing 5 wt% organoclay, 6 wt% of the polymer was bound to the filler, meaning that the mass ratio of adsorbed (bound) polymer to organoclay content was 1.2. This can be converted into a volume ratio of 1.94, using densities of 870 kg/[m.sup.3] and 1400 kg/[m.sup.3] for the polymer and organoclay respectively, as reported by the suppliers. By assuming an idealized composite structure consisting entirely of exfoliated platelets 300 nm in diameter and 1 nm in thickness, and considering that, based on TGA analysis, a composite containing 5 wt% organoclay contains 3.5 wt% inorganic clay, the volume ratio of 1.94 corresponds to an interfacial layer of bound polymer, having an approximate thickness of 2.8 nm. This means that the clay platelets may be entirely surrounded by a layer of adsorbed polymer. In comparison, Fragiadakis et al. [26] reported that the thickness of the interfacial layer of bound polymer in PDMS/silica nanocomposites is between 0.8 and 5 nm.


Given the evidence of presence of adsorbed polymer, we suggest that the narrowing of the region of linear viscoelasticity seen in the stress-sweeps at low clay loadings may arise from filler/polymer interactions. Aranguren et al. [22] suggested that in PDMS/silica-filled systems, possible filler-polymer-filler interactions may include direct bridging of a single chain adsorbed on two separate filler particles, and primary entanglements, between two chains adsorbed on separate particles. Direct filler-filler contact through a thin layer of adsorbed polymer is also possible especially at higher filler loadings [25]. In the case of polymer-organoclay nanocomposites, the interactions can be even more complex, involving directly bound polymer onto the silicate layers, as well as interactions between the organic modifier and polymer [27].

At high loadings, when platelets actually have a high probability of coming into direct contact with each other, the nature of the polymer/filler interactions may change, with a higher probability for the platelets to come into direct contact with each other. Indeed, as the clay loading increased, the ratio of adsorbed polymer to filler decreased (mass ratio of 0.7, or volume ratio of 1.2 for the composite containing 20 wt% clay). In Fig. 4a, at loadings above 10 wt% there was a complete loss of the Newtonian plateau, accompanied with power-law dependence with frequency of the complex viscosity, G' and G" with respect to frequency (Figs. 4a-4c). The elastic and loss modulus curves were almost parallel and actually G' became higher than G", resulting in tan [delta] values that were lower than unity and almost independent of frequency (Fig. 4d). The EOC-g-MAn-based composites exhibited identical behavior. These observations are akin to those reported for systems undergoing chemical or physical crosslinking, at the onset of the gel point, where a network structure begins to form [28]. For the present systems, these findings may signify the presence of a percolation threshold, above which filler/filler interactions are established.

Mechanical Properties and Thermal Stability. Substantial increases in the Young's moduli were noted as the organoclay content increased, as demonstrated in Table 2. Although this was accompanied with a loss of ductility in the EPR-based nanocomposites, the EOC nanocomposites maintained a ductile behavior to a large extent. Our previous work on maleated polyethylene-based nanocomposites [29] has shown that the reinforcement capability of the clays is moderated by a decrease in crystallinity. Apparently this is not the case in the inherently low-crystallinity systems that are under consideration in the present work (see also Table 1), since our DSC studies revealed that the crystallinities of both elastomers remained virtually unaffected in the presence of clay.

The thermal stability of the dimethyldialkyammonium-exchanged clay was investigated by TGA for 30 min at 200[degrees]C, which corresponded to the temperature used to prepare the composites. No substantial weight loss was observed, which suggests that the dimethyldialkyammonium cation associated with the mineral surface was stable under compounding conditions. In agreement with previous reports on similar polyolefin elastomers [20], addition of clay increased the onset temperature for thermal degradation with respect to the pure polymer (Table 3). A relatively large increase in onset temperature was gained with just 2 wt% of clay, beyond which further increases were only marginal.

Polyolefin Elastomer/PP Nanocomposites

Structure and Morphology. Addition of PP did not affect the XRD traces of the nanocomposite blends, which were very similar to those recorded for the elastomer nanocomposites (see Fig. 5). It should be noted that PP/organoclay composites exhibited a strong diffraction peak (included for comparison in Fig. 5), which is evidence that the nonfunctionalized PP was incapable of exfoliating the organoclay, under the conditions employed in this study.

TEM images, wherein the elastomer phase had been preferentially stained with Ru[O.sub.4], revealed that the clay partitioned preferentially within the functionalized elastomer matrix in a primarily exfoliated state, whereas little or no clay existed in the PP phase (see Fig. 6). A large disparity in the viscosities of the components has been cited as a possible reason for preferential partition of fillers, with the fillers having a tendency to localize in the less viscous phase [30]. However according to Fig. 7, which shows the shear viscosities of the PP, EPR-g-MAn and EOC-g-MAn components, the viscosities were not significantly different at a shear rate range of 50-100 [s.sup.-1], which is relevant to shear rates encountered during compounding in the batch mixer. The filler always partitioned in the elastomer phase, regardless of whether its viscosity was slightly higher (EOC-g-MAn), or lower (EPR-g-MAn) than that of the PP phase.

The large difference in melting points of the elastomer and PP components (see also Table 1) could be another possible reason for the observed partitioning, because the lower melting point elastomers would tend to wet and incorporate the clay before the PP phase melts. However this possibility was excluded by compounding the PP and elastomer prior to the addition of clay. TEM imaging confirmed that once again the organoclay was located into the elastomer phase. The most plausible reason for the preferential localization of clay in the elastomer phase would then seem to be its affinity to the maleated EPR and EOC components, which may have higher surface energy due to the presence of the succinic anhydride functionality.


Both blend systems displayed droplet matrix morphology, as seen by SEM (Figs. 8a and 8c), with an average droplet size of 2.3 [micro]m for the EPR-g-MAn/PP blends and a much finer dispersion with an average droplet size of around 1 [micro]m for the EOC-g-MAn/PP blends.



In the presence of clay the size of the droplets decreased to the point that the dispersed phase was not easily discernible (Figs. 8b and 8d). The reduction in particle size was observed in both systems under investigation and became more pronounced as the amount of clay increased, as summarized in Fig. 9. This observation is in agreement with previous observations by Mehta et al. [14] on TPO systems, as well as numerous reports on other polymer blends [see for example Ref. 31, 32]. Among the various explanations offered in the literature, the most plausible ones are that organoclays alter the interfacial tension between the two components, thus generating a compatibilizing effect [31], or that dispersed platelets in the matrix act as physical barriers, preventing droplet coalescence [31, 32]. The latter explanation is more plausible in this situation, where the majority of the clay resides in the elastomer matrix; this is the subject of ongoing investigation.


For both blends the particle size leveled off at around 0.4 [micro]m and no further reductions in particle size were noted beyond 10 wt% clay content. This may be associated with the limited capability of this system to exfoliate additional clay at loadings above 10 wt%, as evidenced also by XRD analysis (see Fig. 5).

Rheology, Thermal and Mechanical Properties. The rheological properties displayed features identical to those reported previously for the elastomer matrices, including loss of Newtonian plateau and power-law type response above 10 wt% clay loading. This suggests that the viscoelastic properties of these blends are largely dominated by the properties of the elastomer matrix. However, given the complexity of these multiphase systems, the rheological properties of these blends are the topic of ongoing investigation.



Introduction of 30 wt% of PP in the elastomer increased their stiffness, while, as expected, reducing their elongation at break (Table 2). The Young's modulus, stress, and elongation at break of the unfilled EOC-g-MAn/PP blends were substantially higher than those of the EPR-g-MAn/PP ones, which is obviously attributed to the better properties of the base EOC-g-MAn material. Addition of clay to the blends resulted in further enhancements of the Young's modulus and, correspondingly, a reduction in elongation at break. In the EPR-g-MAn-based systems, the tensile stress and Young's modulus seemed to level off above 10 wt% clay. The slight reduction seen beyond this loading may be due to inhomogeneity in the dispersion of clay platelets and tactoids.

The substantial improvements in stiffness may be associated to some extent to changes in the crystallization behavior of the polypropylene component in the presence of clay. As shown in Fig. 10, in the unfilled blends the PP component crystallized at different degrees of supercooling, with homogeneous nucleation being favored for the EOC-g-MAn/PP blends as evidenced by a shoulder at the right of the EOC peak around 80[degrees]C. In the unfilled EPR-g-MAn/PP blends PP nucleation took place around 115[degrees]C, which is the usual region of activity of heterogeneous nuclei. However, it seems that most of the heterogeneous nuclei became active at lower temperatures than 115[degrees]C, possibly because of smaller size or lower degree of perfection [33].

In both TPO nanocomposites the peak at 115[degrees]C, corresponding to heterogeneous nucleation, was more pronounced as the clay content increased, signifying that this was the preferred mode of nucleation in the presence of clay. Normally, we would attribute this to clay platelets serving as nucleating agents, favoring PP crystallization by heterogeneous nucleation [34]. However, such nucleation effects in immiscible blends would be expected to occur only if the filler resided inside the PP domain [35], which is clearly not the case in our blends. The change in nucleation behavior in PP may therefore be caused by either a few platelets that reside at the interface of the two components, or could be attributed to complex self-nucleated crystallization phenomena [36, 37] that are associated with changes in the blend morphology. Further research is required to clarify these findings.



Figure 11 shows representative XRD traces from the PP crystal region for the EPR-g-MAn/PP blend and its nanocomposite containing 5 wt% organoclay. In spite of the differences in crystallization behavior noted above, no differences in the type of PP crystal form can be seen in this figure, with the monoclinic [alpha]-form dominating in all cases.

Upon addition of clay in the blends the degradation onset temperature increased in a similar fashion as that reported for the elastomers (Table 3). Figure 12, which depicts representative TGA traces obtained for the EOC-g-MAn-based system, shows that blends containing PP have generally lower degradation temperatures, reflecting the inherent instability of PP relative to the ethylene-rich elastomer.


A hybrid nanocomposite structure consisting of mixtures of exfoliated and intercalated clay was created by compounding onium-ion exchanged montmorillonite clay into blends of maleated polyolefin elastomers with polypropylene. The clay resided exclusively in the maleated elastomer phase. The blend morphology and crystallization behavior were altered in the presence of clay, with the blends containing clay having very fine domain sizes.

Solid-state analysis of these composites has revealed substantial improvements in Young's modulus of the elastomers, with the EOC-g-MAn-based systems having a more favorable overall balance of stiffness and ductility. The onset of thermal degradation was retarded by the presence of clay platelets in all systems.


Materials were kindly supplied by E.I. DuPont Canada Company and ExxonMobil. The authors thank Prof. J.S. Parent of the Department of Chemical Engineering, Queen's University for his valuable input, Dr. Y. Liu for performing the bound rubber content testing, and Mr. Doug Holmyard of the Mt. Sinai Hospital for his assistance with TEM imaging.


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Jeremy R. Austin, Marianna Kontopoulou

Department of Chemical Engineering, Queen's University, Kingston, Ontario K7L 3N6, Canada

Correspondence to: M. Kontopoulou; e-mail:

Contract grant sponsors: Materials and Manufacturing Ontario (MMO); Natural Sciences and Engineering Research Council (NSERC).
TABLE 1. Material properties.

Material EPR-g-MAn EOC-g-MAn

Trade Name Fusabond[R] MF-416D Exxelor[R] MDEX 95-2
MFI/MFR (g/10 min) 23 (280[degrees]C) 7.5 (190[degrees]C)
Density (kg/[m.sup.3]) 870 830
Graft content (wt%) 0.50-1.0 0.35
Melting Point ([degrees]C)* 24.8 73.4
Crystallization Point 13.3 60.5
% Crystallinity* 7 15

Material PP

Trade Name Escorene[R] PP1042
MFI/MFR (g/10 min) 1.9 (230[degrees]C)
Density (kg/[m.sup.3]) 900
Graft content (wt%) --
Melting Point ([degrees]C)* 163
Crystallization Point 115
% Crystallinity* 57

* Determined by DSC.

TABLE 2. Mechanical properties of the elastomer and elastomer/PP

Content Young's Modulus % Elongation Stress at Break
(wt%) (MPa) at break (MPa)

 0 3.70 [+ or -] 1.26 3231 [+ or -] 658 13.82 [+ or -] 3.20
 2 5.20 [+ or -] 0.69 3021 [+ or -] 285 14.38 [+ or -] 1.46
 5 6.71 [+ or -] 0.53 2906 [+ or -] 100 15.53 [+ or -] 1.13
10 9.64 [+ or -] 1.44 2410 [+ or -] 534 13.26 [+ or -] 6.18
20 13.08 [+ or -] 0.81 2393 [+ or -] 193 14.35 [+ or -] 7.05

EOC-g-MAn/PP 70/30
 0 18.76 [+ or -] 2.80 2628 [+ or -] 183 13.60 [+ or -] 6.90
 2 30.54 [+ or -] 4.41 1786 [+ or -] 729 13.81 [+ or -] 5.60
 5 37.76 [+ or -] 3.03 1144 [+ or -] 613 10.22 [+ or -] 5.46
10 34.90 [+ or -] 2.14 1344 [+ or -] 688 9.88 [+ or -] 5.25
20 42.67 [+ or -] 3.11 703 [+ or -] 749 7.22 [+ or -] 5.58

 0 1.51 [+ or -] 0.27 1372 [+ or -] 217 1.06 [+ or -] 0.06
 2 1.51 [+ or -] 0.30 1057 [+ or -] 221 1.17 [+ or -] 0.04
 5 1.84 [+ or -] 0.24 633 [+ or -] 154 1.29 [+ or -] 0.26
10 3.16 [+ or -] 0.15 433 [+ or -] 52 1.56 [+ or -] 0.26
20 3.99 [+ or -] 0.56 322 [+ or -] 61 2.03 [+ or -] 0.14

EPR-g-MAn/PP 70/30
 0 3.59 [+ or -] 0.27 297 [+ or -] 30 0.81 [+ or -] 0.14
 2 11.43 [+ or -] 1.29 140 [+ or -] 30 1.79 [+ or -] 0.22
 5 21.61 [+ or -] 3.10 122 [+ or -] 51 2.59 [+ or -] 0.51
10 23.30 [+ or -] 3.70 75 [+ or -] 26 3.45 [+ or -] 0.68
20 21.52 [+ or -] 2.11 114 [+ or -] 27 3.25 [+ or -] 0.39

TABLE 3. Onset temperature for degradation as a function of clay

Clay Degradation Onset Temperature ([degrees]C)
Content EOC-g-MAn/ EPR-g-MAn/
(wt%) EOC-g-MAn PP 70/30 EPR-g-MAn PP 70/30

 0 399 386 415 389
 2 433 415 455 429
 5 436 420 450 420
10 440 426 454 424
20 441 433 458 433
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Author:Austin, Jeremy R.; Kontopoulou, Marianna
Publication:Polymer Engineering and Science
Date:Nov 1, 2006
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