Rheological and mechanical behavior of blends of styrene-butadiene rubber with polypropylene.
There is currently a lot of interest in "soft touch" materials--materials that are overmolded on hard or stiff structural forms to give a soft, comfortable, often textured, grip or surface for hand tools, domestic appliances, writing instruments, cellular phones, cameras, etc. . Such soft touch components are usually formed from thermoplastic elastomers (TPE) . TPEs combine the moldability of conventional thermoplastic polymers in the melt state with the elasticity, hardness, fracture resistance, and surface durability of cross-linked elastomers.
This study considers the processing-structure-properties linkages of blends comprising polypropylene (PP) and styrene-butadiene rubber (SBR). Such blends form the basis for TPE or thermoplastic vulcanizate (TPV) fabrication, prior to the addition of oil extenders and elastomer cross-linking. The use of SBR in TPE or TPV blends has been shown to improve properties by decreasing melt shear viscosity and solid modulus, compressive set. friction coefficient, and oil absorption [3-5]. The blends are formulated so as to generate continuous thermoplastic phases [6-9], thereby generating composites that have rheological properties dominated by the low viscosity PP in the melt state and mechanical properties dominated by SBR in the solid state. The goal is to explore the dependence of the resulting microstructure on the components (especially weight fraction of the constituents) and the dependence of the properties (especially melt viscosity and solid modulus) on microstructure. In contrast to previous studies [3-5], the SBR was not cross-linked in order to explore the microstructural dependencies of the base components.
Table 1 shows the constituent polymers with their characterization. Five different SBR compositions were used: one commercial product (Goodyear, Solflex[R]) and four custom synthesized materials. The custom synthesized materials consisted of two each of low (about 200 kg [mol.sup.-1]) and high (about 400 kg [mol.sup.-1]) molecular weight materials in both linear and branched chain configurations. Molecular weight was determined by gel permeation chromatography and a combination of light scattering and refractive index detectors. Long chain branching was accomplished by proprietary synthesis technology of Goodyear and was verified by rheological methods. The low molecular weight-linear configuration ("LL") material was similar to the commercial material (identified as SX). The rubbers were in crumb form, contained approximately 5 wt% talc, and consisted of similar styrene and vinyl contents. Mooney viscosity decreased with decreasing molecular weight and on branching.
A single, commercial PP material (Atofina, 7823M) is the same material as used in previous studies [3-5] and was obtained as pellets. The glass-transition temperatures of the SBR materials were similar (about -60[degrees]C), considerably less than that of the PP (0[degrees]C). Table 2 shows the elastomer blends formed: 1) a composition sweep over 20-45 wt% PP with the commercial SBR (identified as percentage PP-SX), and 2) a structure sweep over the custom SBR materials (identified as LL, LH, BL, BH) at 35 wt% PP. The intent of the composition sweep was to test the dependence of properties and morphology on the phase fractions. The intent of the structure sweep was to test the dependence of properties on SBR structure above the PP continuity boundary.
Blends of PP-SBR were compounded using a 16-mm co-rotating twin-screw extruder shown schematically in Fig. 1a (Prism). PP pellets were tumbled with the SBR crumb. The mix was then fed with a constant rate feeder to the extruder, which was held at 200[degrees]C. Two different screw speeds were used, 70 rpm for compounding blends to study the effects of PP content and 130 rpm to study the effects of SBR structure. As the extrudate exited the die of the extruder it was fed into an ice bath. Quenching the blends in this way served two purposes: it trapped the non-equilibrium morphology of the blends and allowed the blends to be easily pelletized.
A ram injection molder (Morgan-Press Model G-55T) was used to form tensile specimens (50 mm long X 10 mm wide X 3 mm thick, with 25-mm gauge). The molding operation is shown schematically in Fig. 1b. Molding was performed using barrel and nozzle temperatures of 200[degrees]C and a mold temperature of 52[degrees]C. To prevent variations in the thermal history of molded specimens, the barrel was loaded with just enough pellets to make one specimen. After loading the barrel, 3 min were allowed to melt the pellets (after the first 90 s, the partially molten polymer was compressed for 5 s to remove air), [10.sup.5] Pa was then applied for 1 min to pack the mold. Pressure was then released, the sample was allowed to cool for 2 min and then removed from the mold. After each injection, the barrel was purged and then reloaded to make another specimen, ensuring that each of the blends and specimens had the same residence time in the injector barrel and the same thermal history. No subsequent treatment was used to cross-link the SBR.
To study the effects of annealing, some full barrel injection molding experiments were carried out. The shaded bands in Fig. 1b indicate the increasing amounts of (barrel + mold) time material in the injection molder was exposed to 200[degrees]C in initial multi-specimen full barrel loading experiments. It was found that mechanical properties depended significantly on thermal history, presumably due to coarsening of the morphology: Increasing the residence time in the barrel from 4 to 15 min decreased elastic modulus and tensile yield stress of a 35% PP blend by about 50% and 15%, respectively and increased failure strain by about 30%. Other small-scale specimen fabrication methods--microinjection molding and compression molding--were also examined. Compression molding was not able to fabricate specimens of all compositions and could not fabricate specimens fast enough to be comparable to industrial processes . Similarly, the low pressures used during microinjection molding (DACA) were not comparable to industrial processes and led to long thermal exposures. Ram injection molding required small quantities of each blend (~5 g per sample) yet used high pressures, allowing for relatively fast production of specimens with a consistent thermal history, comparable to industrial manufacturing processes.
[FIGURE 1 OMITTED]
Surfaces for scanning probe microscopy were formed by cryo-microtoming sections, using a diamond-knife, from the tabs of molded tensile specimens, as indicated in Fig. 1b. The specimen sections were then scanned using a Si cantilever probe in non-contact, constant-amplitude "tapping mode" (Digital Instruments Multimode NanoScope IIIa), recording the phase of the probe motion relative to that imposed on the entire cantilever as a function of position over 30 [micro]m X 30 [micro]m squares. Images were formed by conversion of phase into intensity as a function of position [10, 11]. Typical images are shown in Fig. 2.
The area fraction of PP present in the sections was determined digitally by converting sections of the images into black and white, i.e., via a binary threshold, such that the observed smallest isolated SBR feature remained isolated after rendering, and then counting the pixels above the threshold. Five or six 15 [micro]m X 15 [micro]m sections within the 30 [micro]m X 30 [micro]m squares were analyzed for each material. The scale of the PP and SBR structures in the sections was also determined digitally through similar rendering of the entire image followed by counting of pixels on the thresholded boundaries to give the interface length per area . For a random section of a uniform microstructure, as intended by the diagonal cut through the blind end of the tensile specimen mold (Fig. 1b), the area fraction is identical to the volume fraction (in this case the PP volume fraction) and the reciprocal of the interface length per area is the mean interfacial intercept length (in this case the mean distance between PP-SBR interfaces) .
[FIGURE 2 OMITTED]
Rheological measurements were performed on melts of pelletized blends using a capillary rheometer (Gottfert 1500). Measurements were made at 200[degrees]C, using a 12 mm diameter barrel and a 1 mm diameter, 20 mm long capillary die, over a shear rate range of 10-5000 [s.sup.-1]. Rheological analysis was performed on three samples of each blend, along with the base PP and commercial SBR materials. Apparent viscosity vs. apparent shear rate are reported. No shear rate or entrance corrections were made.
Quasi-static elastic, plastic, and failure properties of solid, molded materials were measured by uniaxial tensile testing of molded specimens using a servohydraulic test frame and wedge-action grips (MTS). The tensile tests were performed under displacement control using a strain rate of [10.sup.-3] [s.sup.-1], measuring load and displacement during the test and converting these to engineering stress and strain via the specimen gauge dimensions. Three specimens of each composition were tested. Average values are reported.
To assess the resistance of the blends to localized sharp contact, conventional micro-hardness measurements were performed. A Vickers diamond square pyramid indenter was dead-weight loaded onto specimen surfaces to a peak load of 2.98 N for 5 s using a gravity-controlled hardness tester (Zwick). The diagonals of the resulting residual square contact impressions (e.g., Fig. 3) were then measured by optical microscopy within 5 min of indentation to determine the contact area. (Time-dependent recovery of the impression seemed negligible. The impressions did not alter in size 24 hr after indentation consistent with control of the plastic deformation by the PP; if the SBR was cross-linked, recovery would occur in less than 5 min.) Hardness of the materials, the mean supported contact stress, was then determined from the ratio of the contact load and residual contact area . For the permanent deformation here, the hardness is a measure of the resistance to plastic deformation, in contrast to the resistance to reversible elastic deformation sensed in often-used blunt probe rubber penetrometer tests, e.g., Shore A [3-5, 14].
[FIGURE 3 OMITTED]
Representative force microscope images of four of the microtomed blends are shown in Fig. 2; PP appears as the lighter phase. Figure 2a and b illustrate the effects of increasing PP wt% for the 20-80 and 45-55 blends containing commercial SBR. Figure 2c and d illustrate the effects that result from changing the structure of the custom SBR for the (35-65) LH and BH materials. The microstructures appear well-mixed, although full interpenetration or cocontinuity of the phases is not obvious. In fact, the blends appear to consist of a continuous PP phase containing extended, but isolated, SBR components. This microstructure was observed in all the blends examined and is similar to that observed in a TPV blend of Nylon-6 and ethylene-propylene-diene terpolymer (EPDM) .
Figure 4 shows the PP area fraction and the mean interfacial intercept length in the blends. The solid lines in Fig. 4a and b indicate the anticipated volume (and thus area) fraction microstructure trends with the processing parameter of PP wt% (making the small correction from wt% to vol% accounting for the slight difference in density between PP, 0.94 g [cm.sup.-3] and SBR, 0.93 g [cm.sup.-3] ). Nothwithstanding the scatter, indicative of the microstructural heterogeneity visible in Fig. 2, the observed increase in volume fraction in the PP sweep and invariance in the SBR sweep are in agreement with the predicted values, suggesting negligible entrained air or mutual dissolution.
Figure 4c and d shows that the characteristic length scale for these composite microstructures is about 1 [micro]m. On increasing the proportion of PP in the blend, the scale of the composite microstructure increased somewhat more rapidly than volume fraction PP, shown in Fig. 4c. As intercept length is the reciprocal of interfacial area per volume, the implication is that as the fraction of PP was increased, the microstructures became increasingly clustered into larger, single-phase SBR and PP domains, minimizing the SBR-PP interfacial area. In the custom SBR blends, the microstructural scale increased significantly on increasing the SBR molecular weight but appeared relatively insensitive to chain configuration, as shown in Fig. 4d. The implication is that the microstructural scale increased with viscosity of the SBR phase.
The melt flow characteristics of the blends determined by capillary rheology are shown in Fig. 5 as viscosity vs. shear rate, along with the characteristics of the constituent PP and commercial SBR. All materials were shear thinning with the PP exhibiting the smallest viscosity and the SBR the greatest viscosity over the shear rate range. As perhaps anticipated, the viscosity of the blends in the PP content sweep lies between the characteristics of PP and the SBR homopolymer, Fig. 5a. Particularly at small shear rates, there is a large decrease in viscosity of the composites from the base SBR response for small (~20%) PP additions with smaller, but systematic, decreases on further PP addition. The large viscosity decrease and Fig. 2c both indicate that PP is a continuous phase, even at 20%. It has been found  that
[[[phi].sub.1]/[[phi].sub.2]]([[eta].sub.2]/[[eta].sub.1])[.sup.0.5] = 1 (1)
can be used to indicate "the onset" of continuity for such large viscosity ratios (where [[phi].sub.i] and [[eta].sub.i] are the fraction and viscosity for each phase). We compare the viscosities at 30 [s.sup.-1], which is the average shear rate in our twin-screw extruder . Using these viscosities for PP and SBR from Fig. 5a, Eq. 1 predicts that PP will be continuous for [[phi].sub.1] > 0.17. If the viscosities at 200 [s.sup.-1] (near the greatest shear rate in the extruder) are used in Eq. 1, [[phi].sub.1] = 0.2.
The viscosities of the composites containing the custom synthesized SBR materials were comparable to those containing the commercial product, Fig. 5b. However, the viscosity appeared to be more sensitive to changes in SBR chain configuration than to changes in PP fraction. Decreasing SBR molecular weight or using a branched rather than a linear morphology decreased the viscosity significantly, consistent with the Mooney viscosity (Table 1). Average viscosity values at a shear rate of 30 [s.sup.-1], [[eta].sub.30], are given in Table 2.
[FIGURE 4 OMITTED]
Solid Deformation and Failure
The solid deformation properties of the composite blends, determined by uniaxial tensile testing, are given in Fig. 6 as the engineering stress-strain responses; the median trace for each composition is shown. Additional testing showed that all materials exhibited reversible linear elastic responses up to approximately 2% strain. This was followed by a gradual yield and the onset of irreversible deformation. Both of these responses are barely discernible in the traces of Fig. 6. For post-yield strains less than about 20%, the deformation in the gauge appeared homogeneous, with a uniform narrowing of the cross-section reaching about 10-15%, consistent with constant volume deformation. Strain recovery (1-5%) was observed on unloading in this region. For strains greater than this, the stress derivative with strain significantly decreased, even becoming transiently negative in some cases (e.g., 45-55, 40-60, and LH), followed by increased strain with a small work-hardening rate towards a peak stress. Some materials exhibited a weak change in the stress derivative and very little subsequent strain prior to the peak stress (e.g., 20-80). The location of the large stress-derivative decrease was strain-rate sensitive, appearing at greater stresses with greater strain rates, as illustrated in Fig. 7.
A localized strain or neck region was observed to form at the point at which the stress derivative decreased; the subsequent, almost stress-free, strain was associated with the increase in length of this necked region along the gauge at almost constant neck width, a further 15-20% decrease in section. The strain prior to the peak stress increased and the stress level decreased with decreased strain rate (Fig. 7). At the peak stress, an extended failure process began that consisted of consecutive failure of discrete outer fibrils of material (about 0.1-0.5 mm in diameter) with stress drops and plateaus until complete failure, leaving the specimen in a "shredded" feather-like state. Very little recovery was observed after the onset of the failure process. PP fraction appeared to alter the stress-strain characteristics more than SBR character.
All deformation and failure properties were evaluated from the tensile data obtained at a strain rate of [10.sup.-3] [s.sup.-1]. Young's modulus was estimated from linear fits to data up to 1% strain. Yield stress was taken to be the 2% offset flow stress. Modulus and yield stress are shown as a function of PP content and SBR configuration in Fig. 8; symbols represent measurements on individual specimens. Both modulus and yield stress increased with increasing PP content, although modulus did so more rapidly, indicating the decrease in the strain at which yielding occurred with increasing PP content. Neither modulus nor yield stress varied strongly with SBR type, although the LH specimens exhibited the largest value for both properties. Also included in Fig. 8 are the hardness values for the materials. Hardness was strongly correlated with yield stress for the PP series materials, varying approximately as hardness/yield stress [equivalent]2.5, consistent with the behavior of some metals, glasses and glassy polymers [13, 18]. The correlation was less strong for the SBR series materials, with the high molecular weight materials (LH, BH) exhibiting much greater hardness than the low molecular weight materials (LL, BL); a resistance to plastic deformation that was not reflected in the yield stress values. Average modulus and yield stress values are given in Table 2.
[FIGURE 5 OMITTED]
[FIGURE 6 OMITTED]
The failure properties, work of rupture (integrated area under full stress-strain curve), ultimate tensile strength, and failure strain, as a function of PP content and SBR structure are shown in Fig. 9. As with the pre-failure deformation properties, all the failure properties increased with increasing PP content. Similarly, the failure properties did not vary strongly with SBR structure, although once again, the LH material was distinguished by a greater strength and smaller failure strain. Average failure strain values are given in Table 2.
Consideration of Figs. 4, 5, 8, and 9 shows that there was no abrupt change in the dependence of any microstructural or mechanical property on PP content; all displayed almost linear dependencies on the wt% PP over the full composition range. This suggests that the form of the microstructure was invariant with composition and that all the materials were composed of continuous PP and isolated SBR phases with the properties controlled by PP. This is somewhat surprising for the smaller PP wt% materials but does suggest greater than expected engineering flexibility in composition and thus properties selection. The large decrease in the viscosity effected by small additions of PP to the SBR (Fig. 5) and the lack of porosity in the microstructures (Figs. 2 and 4a and b) suggest that during the blending process low-viscosity molten PP was able to interpenetrate between SBR fragments, essentially coating the SBR fragments and leading to continuous PP structures. This is similar to that observed in the Nylon-EPDM system  in which the Nylon was continuous. However, the sensitivity of resulting microstructures to processing conditions is highlighted by the fact that the EPDM remained dispersed as isolated spheroids for fractions of 50-60% with a slight coarsening of the structure with increasing EPDM content.
[FIGURE 7 OMITTED]
[FIGURE 8 OMITTED]
Figure 10 illustrates that the microstructures generated did indeed meet the engineering goal of combining the melt characteristics of the PP (important for processing via injection molding) with the solid characteristics of SBR (important for soft touch properties). Plots of solid elastic modulus, E, (from Fig. 8a) and melt viscosity, [[eta].sub.30], (from Fig. 5a) as a function of composite PP volume fraction (taking converted weight fraction to represent volume fraction), are given in Fig. 10a and b, respectively. The solid lines in both plots represent bounds  on the properties assuming equal strain or strain rate distribution (upper bounds, parallel-like properties, q = [[phi].sub.1][q.sub.1] + [[phi].sub.2][q.sub.2], where [q.sub.i] is a constituent property) or equal stress distribution (lower bounds, series-like properties, 1/q = [[phi].sub.1]/[q.sub.1] + [[phi].sub.2]/[q.sub.2]) in the composites, using E(PP) = 1.2 GPa, E(SBR) = 10 MPa,  and [[eta].sub.30] (PP) = 379 Pa s, [[eta].sub.30](SBR) = 9070 Pa s (Fig. 4). In both cases, the data follow the predicted trend and lie much closer to the lower bounds. This suggests that in the melt the properties are dominated by the low-viscosity PP phase, which carries most of the flow strain. After processing, in the solid state, the elastic properties are dominated by the low-modulus SBR phase that carries most of the elastic strain. As indentation of SBR does not lead to a residual contact impression, the resistance to plastic deformation must be controlled by the more-easily yielding PP phase; the increase in hardness with increasing PP fraction suggests that another factor must be controlling both this and the failure properties.
[FIGURE 9 OMITTED]
[FIGURE 10 OMITTED]
The variation of modulus with PP fraction is similar to that observed for PP-styrene/(ethylene-butylene) (SEBS) blends : moduli close to the lower bound for small volume fractions, increasing towards the upper bound for fractions greater than ~40%. Although other properties increased in similar fashions, there was no "synergism"  observed: no property was greater than the equi-strain upper bound. In contrast, PP-sulphonated butyl ionomer (IIR) blends  exhibited ultimate tensile strengths greater than the linear upper bound, suggesting that the mutual constraint arising from full cocontinuity of an interpenetrating phase network or composite [8, 9] was present. The mechanical observations here, consistent with the microstructural observations (Fig. 2) suggest that the PP phase was continuous while the SBR phase was not.
Figures 8 and 9 also illustrate the dominance of PP volume fraction in determining properties: all solid deformation and failure properties increased monotonically with PP fraction in the PP series, and, on average, the properties of the custom 35-65 SBR materials were very similar to those of the 35-65 material (blended using the commercial SBR) in the PP series. The secondary effects of the SBR configuration (with volume fraction held constant) were most evident in the elastic modulus and hardness, as well as in the viscosity. Figure 11 demonstrates this; plotting these properties for all five 35-65 materials as a function of interfacial intercept length. All increase with intercept length, independent of the SBR type. An implication is that SBR configuration does not influence properties directly but, by its rheological behavior, determines the microstruetural scale and it is the scale that is the dominant influence on properties. The increased resistance to deformation of all sorts with increasing microstructural scale suggests that interfacial interactions are not dominating the deformation behavior, but rather that it is the constraint on the deformation of the constituent phase elements that depends on length scale . For these materials, the constraint increases with microstructural scale.
[FIGURE 11 OMITTED]
PP and SBR were blended and injection molded to form precursor TPE composites, varying both the PP fraction and the SBR type. Rheological measurements in the melt state, and microstructural characterization and mechanical measurements in the solid state, exhibited systematic variations in properties, suggesting that the PP and (uncross-linked) SBR phases were not cocontinuous but consisted of isolated SBR domains in a continuous PP matrix. The scale of the interpenetrating phase composite microstructures was approximately 1 [micro]m. Capillary rheometry measurements showed the viscosity of the composites to be intermediate to that of the two phases but dominated by PP, decreasing with PP fraction and increasing with SBR molecular weight and chain length. All materials were shear thinning. Uniaxial tensile tests showed that modulus, yield stress, ultimate tensile strength, and failure strain all increased with PP fraction. The stress-strain responses were strain-rate sensitive, exhibiting decreasing yield stress and increasing failure strain with decreasing strain rate. Hardness tests showed that the materials plastically deformed under sharp (Vickers indentation) contact. Hardness followed yield stress and increased with increasing PP content. Combining the microstructural analysis with the property measurements suggested that strain was discontinuous in the structures during deformation: in the melt state, the PP carried most of the strain and dominated the viscous response; in the solid state, the SBR carried most of the strain and dominated the elastic response. Microstructural analysis also suggested that the major effect of changing SBR type was to alter the microstructural length scale, which in turn controlled the melt and solid deformation properties, independent of SBR type.
The authors thank Mike Malamaci for obtaining the force microscope images; Joel Bell for creating the injection molding image; Steve Henning, Scott Christian, and Mike Kerns for custom SBR synthesis; and Jeff Galloway for helpful discussions.
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Robert F. Cook, Kurt J. Koester, Christopher W. Macosko
Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455
Chemical Divison, The Goodyear Tire & Rubber Company, Akron, Ohio 44306
Correspondence to: R. Cook; e-mail: firstname.lastname@example.org
Kurt J. Koester's current address is Department of Materials Science and Engineering, University of California, Berkeley, CA.
Contract grant sponsor: IPRIME. The Industrial Partnership for Research in Interfacial and Materials Engineering, University of Minnesota.
TABLE 1. Characteristics of SBR and PP components. Source Structure Styrene (%) 1.2 PDB vinyl (%) SBR characteristics Commercial, SX Linear, low Mw 25 15 Goodyear, LL Linear, low Mw 23.7 10.4 Goodyear, LH Linear, high Mw 23.8 12 Goodyear, BL Branched, low Mw 25.4 14.5 Goodyear, BH Branched, high Mw 24.7 11 PP characteristics PP Linear - - Glass transition temperature [T.sub.g] Source Mw (g/mol) Mooney viscosity ([degrees]C) SBR characteristics Commercial, SX 272,000 50 -56 Goodyear, LL 202,200 34.3 -58.2 Goodyear, LH 370,500 115.5 -58.3 Goodyear, BL 208,600 13.4 -59.1 Goodyear, BH 428,000 44.1 -61.6 PP characteristics PP ~50,000 Melt flow index, [T.sub.m] = 145 MFI = 30 TABLE 2. Viscous, elastic, plastic, and failure properties of SBR-PP blends. PP content Melt viscosity, Solid modulus, SBR structure (wt%) [[eta].sub.30] (Pa s) E (MPa) Commercial, SX 20 1780 38 25 1550 61 30 1520 74 35 1330 130 40 1130 155 45 1160 169 Goodyear, LL 35 1350 92 Goodyear, LH 35 1980 152 Goodyear, BL 35 970 75 Goodyear, BH 35 1710 70 Yield stress, Failure SBR structure [[sigma].sub.Y] (MPa) strain, [[epsilon].sub.f] (%) Commercial, SX 2.01 2.07 2.53 3.66 2.88 3.73 3.67 5.31 4.24 6.54 4.70 6.63 Goodyear, LL 2.82 5.89 Goodyear, LH 4.44 3.70 Goodyear, BL 2.69 5.50 Goodyear, BH 2.42 6.16
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|Author:||Cook, Robert F.; Koester, Kurt J.; Macosko, Christopher W.; Ajbani, Manoj|
|Publication:||Polymer Engineering and Science|
|Date:||Nov 1, 2005|
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