the mechanical properties of styrene-butadiene-styrene (sbs) triblock copolymer blends with...

The Mechanical Properties of Styrene-Butadiene- Styrene (SBS) Triblock Copolymer Blends with

Polystyrene (PS) and Styrene-Butadiene Copolymer

JORAM DIAMANT, D A V I D SOONG, and M I C H A E L C. WILLIAMS

Department of Chemical Engineering University of California, Berkeley

Berkeley, California 94720

Measurements were made of linear viscoelastic properties and nonlinear stress-strain properties of phase-separated styrene-butacliene-styretie (SBS) copolymers and their blends with several homopolymer polystyrenes (PS) and one random copolymer (SBR). Torsion pendulum testing yielded shear moduli G', G", and Rheovihron experiments produced tensile moduli E ' , E " , over a 220°K range of temperature, both at low frequencies. For pure copolymers and their PS blends, G' and E' correlated quite well with the total PS content, but G" andE" were more sensitive to how the additive \vas c1istril)rited. Re- sults suggest that a PS additive whose molecular weight (A[) is less than that ofthe copolymer PS-block (hf,y) causes expansion of hoth the interphase and the homogeneous PS-rich phase, while an additive with M > M s mixes less well with these phases (probably forming separate domains of pure PS) and is less effective i n enhancing the linear moduli. The Idending with SBK produced reduction in G' but a broad midrange peak i n G", suggesting that SHK was localized almost entirely within an expanded interphase. Tensile stress-strain data were oh- tainetl with an Material Testing System at room temperature. For PS blends, properties did not correlate well with the total PS coritent, the Idends being always weaker than the SRS ofthe same overall composition. The amount of set also increased with PS content in the blends. Cyclic tests to increasing strain showed progressive structural alterations (as for the host SBS), with blend behavior resembling host properties more closely

le. When SBR was the additive, amounts a s small a s 1 percent reduced the curves by 15 percent. The yield stress was eliminated entirely with an addition of 10 percent SBR, but for all cases the set was the same. Results are dis- cussed in terms of interphase force barriers to chain flow.

INTRODUCTION s a result of microphase separation, block copoly- A mers exhibit a wide range of complicated rheologi-

cal behavior. These polymeric systems often possess mechanical properties superior to either of the homopolymer components alone, rendering them desirable materials for engineering applications (I). Hence, research activities on these systems flourished in the last decade (2-4). A number of thermodynamic models were designed to explain the phase-separated nature of these polymers (5-14). Parallel to this development, mechanical models were being proposed (15-21) to describe the elastic and viscoelastic properties of two-phase composite solids in general. The con-

cept of microphase separation has gradually evolved from that of a two-phase morphology with a sharp interface to one involving a diffuse boundary layer with a continuous composition profile varying betwecn the major phases. Recognition of this structural feature forces substantial modifications of earlier mechanical models. Recent progress in this direction resulted in a structural model (22-23) that accommodated detailed microstructural features, such as volume fractions of the homogeneous regions and the diffuse interphase, as well as the average compositions of these regions and the interphase profile. Success was achieved in fitting the G' ( T ) and G"jT) data over -120°C to 120°C for several SBS samples cast from different solvents (22-24).

POLYMER ENGIN€ERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, No. I I 673

Joruni Diutnunt, Uauicl Sootig, und &fichuel C. Williums

A number of important findings emerged from the above modeling work (22-24). The interphases were found to possess complex and often asymmetric composition profiles, which were influenced by the casting solvents. Hence, kinetic as well as thermodynamic fac- tors were active in shaping the final sample morphology. Next, the polystyrene phase sometimes contained an appreciable amount of the polybutadiene blocks, which acted as polymeric plasticizers to lower the glass transition temperature of the rigid phase. T.he degree of this mixing depended on the molecular weights of the blocks and thc casting solvents. Engineering properties (24), such as stress-strain behavior, were consistent with microstructures deduced from fitting G’(T) and G”(T).

Similar to solids, liquid-phase SBS block copolymers exhibit rheological properties reminiscent of microphase separation. For example, within a certain temperature range, the zero-shear-rate melt viscosity of SBS was found to exceed that of pure polystyre.ne (25). This observation is consistent with thermodyna.mic models (11, 12,26), which predict that block copolymers can ex- ist in a multiphase liquid state at temperatures below the microphase separation temperature, T,, as long as T > T,. While phase-separated diblock systems can flow like ordinary suspensions, triblock molecules endow the phases with a connectedness that enhances viscosity greatly and-depending on morphology-:an produce a yield stress (25, 27). The major physical origin of the anomalous rheological behavior associated with triblock systems lies in the force barrier experienced in pulling PS blocks across the interphase to mingle with PB (28).

Highly complex transitions within these block copolymers were also observed in the past. A strain- induced mechanical transition (29) of solid SBS was reported in which a plastic-like sample transformed into a rubber upon a large deformation. After annealing for a certain period, the specimen reverted to the original plastic state. Time-dependent phenomena suggestive of progressive structural transition appear to cause thixotropy in certain block copolymer melts (30, 31) and solutions (32).

In the present work, the role of microphase separation in affecting the viscoelastic properties of SBS triblock copolymers is further scrutinized t)y observing changes incurred when polystyrene (PS) homopolymer or SBR is blended with the copolymers. Through the microstructural mechanical model established previously for linear properties, we ascertain in which region of the blend system the added components lie and un- derstand how the linear viscoelastic properties compare with those of the parent SBS. Stress-strain data of large deformations are used as further evidence to demon- strate the disproportionate effects introducd by even small amounts of the additives. Hysteresis loops are obtained, indicating the different patterns of structural degradation of the blend systems, in contrast to the original block copolymers.

This work impinges on two areas of potential com- mercial interests in utilizing SBS. First, if the SBS-PS blends perform as well in engineering applications as

pure SBS copolymers having the same PS content, the latter can be replaced by the former, less costly blends. However, as will be shown later, appreciable differences between the two are observed in large deformations despite similar linear properties, suggesting cau- tion in substituting one for the other. Next, SBS-SBR blends display substantial decreases in (or even complete elimination of) the yield stress in solids, when only small amounts of SBR are present. This may have potential implication in melt processing, as the addition of a small amount of SBR may greatly reduce the viscosity of the SBS.

EXPERIMENTS Three SBS triblock copolymers with different com-

positions and block lengths (Table 1) were obtained from the Shell Development Co. Here, they will be designated as 0.48 PS, 0.29 PS and 0.27 PS to reflect their polystyrene content. Note that the 0.29 PS and 0.27 PS samples have comparable PS content but much different molecular weights. Four homopolymer PS and one random copolymer SBR (75 percent butadiene: 25 percent styrene) were obtained from the Pressure Chemical Co. The former samples are glassy solids with weight-average molecular weight, Gw, of 4,000; 4,800; 10,300; and 20,400; whereas the SBR is a liquid at room temperature with Gu, = 3,600.

Sheet samples of both SBS and its blends were prepared by spin-casting from a tetrahydrofuran (THF) and methyl ethyl ketone (MEK) mixed solvent in a 9:l volume ratio. This is an excellent solvent system for PS (8soIup711 = tjPS = 9. 1) but only moderately good for PB (aPB = 8.3). The tendency for this preferential solvation promotes PS phase continuity and increases the relative amount of PS in the interphase region (22-24).

Samples were dried in a vacuum oven at 60°C until constant weight was achieved (approximately two weeks) and then were annealed at 100°C for 24 h. The annealing temperature was higher than the T, of PS and, thus, sufficient to soften the PS phase, relieving locked-in stress and allowing partial improvement of microstructural regularity. However, 100°C is lower than T,v in the 0.48 PS and 0.29 PS block copolymers [T,

242°C and 462”C, respectively (26), with strong M-dependence], so complete equilibration of morphology was probably not achieved for these samples. In the 0.27 PS copolymer, with T , 25”C, microstructures may have eqnilibrated during annealing.

Table 1. Experimental SBS Block Copolymers*

TR 41 TR 41 TR 41 Designation -1649 -1648 -1647

Molecular Weight Ms, 14 16 7

. 10-3 M, 30 78 36 M~~ 14 16 6 M 58 110 49

Polystyrene Weight Fraction 0.482 0.293 0.268

* Prepared by sequentlal anionic polymerlzation at Shell Development Co. The polybutadlene block composition Is approxlmately 4w6 CIS 1.4; 50% trans 1.4; and 10% Of 1.2.

674 POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, No. 1 1

The Mechanical Properties of Styrelie-Biitadicne-Styrene ( S B S ) l’rihlock Copolymrr Blrncls

The annealed samples were tested in a torsion pendulum at 0.5 Hz to obtain shear moduli G’(T) and G”(T) over a temperature range of - 120°C to 120°C. The amplitude ofoscillation was limited to the range giving rise to linear viscoelastic responses only. This was ensured by noting that material properties were independent of the amplitude of the oscillation in the damping curve used for data analysis. Similar dynamic mechanical tests were done on two SBS-PS blends using a Rheovibron over the same temperature range, giving tensile moduli E’JT) and E”[T).

Stress-strain curves were obtained at room temperature using a Material Testing System (MTS).

RESULTS AND DISCUSSION Linear Properties

The linear properties, by definition, reflect the microstructure of samples in their virgin state. This makes it possible for basic morphological information to be inferred and provides insight about subsequent structural alterations that occur when nonlinear properties arise for large deformations.

Our previous work with SBS copolymers (22-24) led to the development of a conceptual model for these multiphase systems and a rather straightforward method for calculating their linear mechanical properties when certain microstructural features are known. Conversely, fitting the model quantitatively to linear properties permits inferences about the microstructure. These structural parameters are so general that the model can be applied to many types of microphase systems, including the blends of SBS with homopolymer PS and random copolymer SBR. This will prove to be useful in the data analysis that follows.

Blends with PS. The experimental shear and tensile moduli ofthrce blends are shown in Figs. 1 and2. Table 2 summarizes the blend compositions. In each case, a base SBS copolymer is mixed with a PS homopolymer to raise the overall PS content. We are interested in both the practical consequences of doing this and information about the microstructural location of the extra PS.

Figure I displays the result of “upgrading” an SBS polymer having 27 percent PS by adding enough homopolymer to achieve the level of 48 percent PS. The major effect is to raise G’ and G “ throughout their plateau regions, which include the normal temperatures of use. Thus, the blend exhibits properties that are closer to plastics than were those of the rubbery pure SBS. For

I I 1 I I I I I I

Torsion Pendulum, 0.5 Hz

Blend with PS (4,000)

-120 -90 -60 -30 0 30 60 90 120 Temperoture ( “ C )

(0) Fig. 1 ( ( 1 ) Storuge shecir ~ i t o d u l u s G ’ f o r three SBS .systrmv: p i r e copol!/irtc~rs 0.27 PSjO) untl 0.48 P S (L j), und ( I blend of the 0.27 PS copolymer with PS homopolymer (a) to cichietic ( in oixrull PS content of 48%. See 7’(ibles 1 und2forfiirther churcicterizu- t i on . The I > l ~ n c l dutu (ire fitted by the mrrltiplzcise i)i(idel O J P

timu11;y f-) uritl c r l s o less well h y cis.siinting (1 liric~ur i t i t c r - phase composition profile {-----) und a sharp interfuce i......). Frequency i s 0.5 H z . ( k ) L0.w modulus 6” f o r the .sumr .system.

comparison, the properties of pure SBS containing 48 percent PS are also represented in F i g . I ; these moduli are found to be only marginally different than those of the blend, with a slightly higher G’ and a slightly lower G” in the plateau (thus, also room temperature) regions. This result implies that the blend properties over a 220°C range are governed primarily by the total PS content, as long as the phase-separated nature of SBS microstructure is maintained. [The very different na-

Table 2. Blends with SBS Copolymers

Base SBS Copolymer Additive Blend

Styrene Styrene - Additive Styrene Content Label Content M, Type Content M, Concentration

TR41-1647 26.8% -7,000 PS 100% 4,000 29.3% 48.3% TR41-1647 26.8% -7,000 PS 100% 4,800 30.0% 48.6% TR41-1647 26.8% -7,000 PS 100% 4,800 44.0% 59.0°/o TR41-1648 29.3% 16,000 PS 100% 10,300 27.2% 48.4% TR41-1649 48.2% 14,000 PS 100% 4.800 21.2% 59.0% TR41-1649 48.2% 14,000 PS 100% 20,400 21.2% 59. oo/o TR41-1649 48.2% 14,000 SBR 25% 3,600 1 .O% 48.0% TR41-1649 48.2% 14,000 SBR 25% 3,600 10.0~/0 45.9%

POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, NO. 1 1 675

Joruni Diumunt , Daoid Soong, und Michael C . Wil l iams

A S B S (0.482 PS) Blend with ASBS (0.482 PSI Blend wi th OSBS (0 .482 PS)

z 5

PS 14,800) PS (20,400)

t -I

I00 24 ( ( 1 )

SBS (0.48 PS) Blend with PS (4,8001 SBS (0.48 PS) Blend with PS (20,400) SBS (0.48 PS)

tures of the homogeneous parent polymers, PS and PB, are well-known and displayed elsewhere (22-24)].

The added PS has a molecular weight of 4,000 in this case, substantially smaller than the PS blocks ( M = 7,000) of the SBS. Thus, a single molecule of the additive would be capable of fitting entirely within a styrene domain formed from the host SBS; these domains have a minimum dimension roughly twice that of the S-block random coil size. If only a small amount of PS were added, we could reasonably expect that it would reside entirely within the host styrene domains. In the case ofFig. 1 , however, the added PS accounts for 29.3 percent of the total blend volume, contributing an amount of PS comparable to that originally existing in the SBS. This suggests that other, more subtle, structural alterations may have occurred to accommodate the additive. In particular, the elevation of the G" plateau is known to be associated with growth of the volume fraction of an interphase region (22-24). Fur- ther, the elevation of G' is associated with morphological changes in a direction leading from dispersed PS spherical domains toward dispersed PS cylinders, interconnected cylinders and, ultimately, to alternating PS/interphase/PB lamella (22, 23).

These considerations are qualitative, but may be evaluated quantitatively by application of the microstructural model cited above. The model is reviewed briefly in Appendix A, but a few general remarks about it will be made here. In its previous application to pure SBS copolymers (22, 23), the G'JT) and G"(T) curves given in Fig . 1 were fitted very well by the following procedure:

Employ as input the corresponding C' and G" curves for the component homopolymers, PS and PB. Characterize the microstructure in terms of model-independent parameters that are tied together by material balances. These include the volume fractions -4,k ofthe three mqior regions (k = PS-rich homogeneous phase. PB-rich homogeneous phase, and interphase of smoothly varying composition), the corresponding average compositions in terms of local PS volume fraction d k , and the interphase composition profile <b,(x). Introduce the one model parameter, @lnc,s,

describing the maximum packing fraction for the dispersed-phase configuration. (This is tied to the G' plateau level.) Allow these parameters to vary in a systematic fashion until an optimal curve-fit is obtained.

In cxtending this structural model to apply to blends involving SBS triblock copolymers, the same basic procedure is followed. The overall material balance, however, is now dictated by the composition of the blends. For example, the styrene content in the two homogeneous phases and the interphase should sum up to the overall styrene content existing in the blend. Next, a decision is made (for the first trial) to identify the phase(s) in which the added polymer reside(s). Here, intuition, as well as thermodynamic compatibility, chain length, and additive molecular weight, all play major roles in a priori prediction of the location of the added component in the blend. Compatibility and molecular weight arguments cited previously suggested a trial wherein all the PS additive was assumed to rcside in the homogeneous PS-rich phase, and this proved to be reasonably successful. However, as expected from arguments based on the large volume of additive, further trials assuming some intrusion into the interphase led to even better curve-fits. The most successful result led to the solid lines drawn through the blend data in Fig. 1 .

The microstructures corresponding to the optimal model curve-fits for the blend and the two SBS copolymers are displayed in Fig. 3. All parameters listed in Fig . 3--4, values and <b values- show that the blend is intermediate between the two pure SRS polymers, but always closer to the 0.48 PS one than the 0.27 PS one. The model places a dividing line between the PS-rich phase and the interphase, which is somewhat artificial, since the operational boundary between them is actually where <b,(x) drops abruptly. Use of this operational definition shows that the blend has a PS-rich phase of as 34 percent, as compared to 28 percent for the 0.48 PS copolymer and 25 percent for the 0.27 PS host. In

676 POLYMER ENGINEERING AND SCIENCF, MID-AUGUST, 1982, Vol. 22, No. 1 1

The Mechunicnl Properties of Styrene-Butudiene-Styrene ( S B S ) Triblock Copolymer Blends

SBS (0.268 PSI Qmax = 0.75

I r i

01

SBS (0.268 PSI Blend with PS (4000) Omax = 0.88

I I I

‘p

0

I

9

0

I I % = O I I I I I I L 1 I J 3

(hi

SBS (0.482 PSI @‘mnw = 0.93

I

I

comparison with the host, then, the blend has both a larger PS-rich phase and a larger interphase. The pecu- liar wiggle in the blend +&) corresponds to the broad G”(T) pcak in the midrange plateau (see Fig. l b ) and probably reflects a nonequilibrium distribution.

The +/(N) shape is critical to successful curve-fits, as demonstrated in Fig . 1 by the two broken lines, which represent model predictions for two unrealistic +/(x) cases. One case is 0, = 0, so that +I is a step function (“sharp interface”), and this fails dismally for G”, espe- cially; the need to assume a substantial interphase region is readily apparent. The second case assumes that +/(x) is linear, which must be incorrect at the bounda- ries and seems a physically implausible curvature, be- sides. Still, a linear +/(.) produces major improvement in the G ” leuel. Its failure in fitting thc G”(T) curvature indicates the sensitivity of both G“ and the model to subtle composition variations in the interphase.

The G’(T) data are fitted quite well even for the two unrealistic +&) cases, demonstrating that C’ is insensitive to interphase composition. Instead, it proves to be sensitive to the microstructural configuration through amas. For the pure 27 percent-PS SBS host, am,, = 0.75 emerged from the optimal model curve-fit (23) and this corresponds closely to dispersed spheres on a hexagonal lattice. To achieve the result shown for the blend in Fig. 1 , the value Qmns = 0.88 is required. This corresponds closely to maximum packing for cylinders on a hexagonal lattice (0.91) and could also reflect the presence of imperfect or disordered lamellar structures (1.00 for perfect order). The pure 48 percent-PS SBS shown in Fig. 1 is optimally fitted (23) by Qnlrrs = 0.93 and is known to have a basically lamellar microstructure. Thus, the fact that C’ curves are similar for the blend and the pure copolymer containing 48 percent PS is only indirectly related to their PS content; more likely, they share similar microstructures, which is a consequence of their PS content. This viewpoint is reinforced by the microstructural details of purcl 48 percent-PS copolymer, shown in Fig. 3c, which are remarkably like those for the blend in Fig . 3h.

In Fig . 2, the tensile moduli E ’ ( T ) and E”(T) are given for the pure 48 percent PS copolymer and for two blends that upgrade total PS content to 59 percent while contributing 21.2 wt. % of the total volume. For the E‘ data, the plateau regions are elevated, as expected, in both blends, but by slightly different amounts. We believe that the difference is explainable by the PS additives having_molecular weights that differ by a factor of about 4 ( M I , . = 20,400 us. 4,800) and, perhaqs more significantly, represent one hl lower and one M higher than the SBS copolymer PS block ( M = 14,000). The PS additive with M,? = 4,800 has the higher plateau level, presumably because the lamellar microstructure of the host copolymer can accommodate it more completely within the preexisting PS layers. For the additive with = 20,400, the larger chains have a more disruptive effect on the morphology (amnx is less).

A specialized interpretation of the morphological dis- ruption is that some or all of the a,. = 20,400 additive has formed a separate pure phase, spherical islands of PS within the other phases. This has been seen by trans- mission electron microscopy (1) when the PS additive R.I exceeds M s of the host. Such PS islands would be less effective than lamella in stiffening the sample and would have a lower characteristic Qmc,z (as spheres), which would bring down the overall average Qmrrs; both these conjectures are consistent with the E‘ data. Since the PS islands would be unplasticized with PB (unlike the block copolymer PS-rich phases; see Fig. 3) , their stiffening effect as fillers would persist to higher temperatures than would the plasticized lamellar phases. This could, indeed, explain the crossover of G‘(T) data for the two blends above 80°C, where the system with the higher-M additive becomes the stiffer.

The E”(T) plateaus for these systems show considera- ble similarity among themselves, something of a sur- prise in view of the appreciable increases in G”(T)

POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, No. 1 7 677

shown in Fig. 1 to arise from PS addition to the base SBS copolymer. The invariance of these plateau leuels means, according to the mechanical model (23), that @ I

is not changed substantially from the base case. If so, then, for the case of the M , = 4,800 additive, there must be a redistribution of the PB and PS blocks of the SBS copolymer between all three regions in order to maintain this condition. Indeed, this is also suggested by the minor differences in the E"(T) curve shapes. The M,. = 4,800 additive causes the curve to tilt up at high T (near the T, of PS) and tilt down at low T (near the T, of PB), the fulcrum being about -20°C. Although @ I re- mains almost unchanged, the interphase becomes relatively enriched in PS, so that +,(x) is even more skewed than before. The alp = 20,400 additive causes even less change in the E" plateau, relative to the base case, as was true also for the E' plateau. There is some elevation in the vicinity of -20" to 40", and also above 80°C, while no changes are evident for -90" to -20" and 40" to 80°C. The midrange "bulge" (-20" to 40") implies a local inflection in the Cbl(x) profile, similar to that seen in Fig. 3b, which characterizes segmental re distribution from the pure SBS (Fig . 3c). However, the major feature ofE"(T) for this additive is its virtual duplication of the curve for the host SBS between -85" and 80"C, which reinforces the concept that mingling of the two polymers is minimal. The slight enhancement of @, probably is a response of the SBS to its own ~ I ( x )

redistribution. At both ends of the temperature span, there is other

microstructural information from E". For the a,. = 4,800 case, the elevated T , peak for PS and reduced T , peak for PB means that DS has increased and QB de- creased, thus confirming the earlier speculation (from E' data) that the additive has migrated primarily to join the host copolymer PS-rich phase. For the iTI,. = 20,400 additive, there is now additional evidence to support the suggestion that separate island domains may form: the T, peak for PS is extended to lOO"C, characteristic of unplasticized PS of G,c > 15,000. From the fact that E"(T) for the blend is coincident with that for the host in the 65" to 80°C range, we infer that virtually none of the additive has gone into the PS-rich phase ofthe host; otherwise, this region of the blend data would have been somewhat elevated.

Blends with SBR. Another way of probing the role of the interphase in determining viscoelastic properties is to add a random copolymer whose average composition would lead to its being attracted preferentially to the interphase. Thus, one can investigate the relative importance of interphase, homogeneous phases, and phase continuity.

F i g u r e s 4 a and4b give G'(T) and G"(T) for the pure 48 percent-PS copolymer and for itsblend with 10 percent by weight of a short-chain SBR (Mlc = 3,600). The SBR has an average composition of 75 percent butadiene and 25 percent styrene, but the exact cornposition distribution along the chain is unknown and this composition may vary greatly over the molecules in the sample. The primary features to note in Fig. 4 are the drop-off of G'(T) for the blend relative to the SBS copolymer, com-

1 0 ~ ~ I I I I I I I 1

Torsion Pendulum, 0.5 Hz 5:

0 2 - -

V SBS (0.48 PS) Blend with 10% SBR (3,600) 0 SBS (0.48 PS)

2 I I I I I I I I

iu)

Jorum Diumunt, D m i d Soong, und Michuel C . U'illiunzs

678 POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, NO. 1 1

-120 -90 -60 -30 0 30 60 90 120

Temperature ("C) ib)

Fig. 4 . ( ( 1 ) Storage sheur nioclrrlus G ' f o r ttuo SBS systems: pure copolynier0.48 P S (5) und (I 1 0 % blerrd w i t h S B R (r). See Tuble 2-for f 11 rther charcrcterizutioir . ( b ) Loss inoc1ulu.s G " f o r the ,su rne s!ystems.

mencing at -4O"C, together with the broad peak in G"(T) commencing at the same temperature and, subse- quently, declining gradually until dropping below the pure copolymer at +30°C and reaching a minimum at +70"C.

Obviously, the added SBR is responsible for both the G' and G" alterations, but this qualitative observation alone cannot identify the microstructural location of the additive. The three major possibilities are that the SBR may reside primarily within the interphase (as suggested previously), or that it is mixed uniformly with either of the two homogeneous phases, or that it exists as a separate phase (i.e., dispersed islands) within the two homogeneous phases. The latter possibility can be tested by computing T , for the SBR, since the broad G" peak should be centered close to this T , if the SBR formed a separate phase. However, as shown in Appen- dix 2, the various plausible ways of computing T:BR for an SBR of this composition fall in the range -67" to -47°C. This is entirely below the observed range of the broad G" peak in F i g . 4b, so this microstructural possibility may be discarded.

If the added SBR is uniformly mixed with one of the homogeneous phases, this would change T, of that phase and shift the corresponding G" peak either to

lower temperatures (for the PS peak at 90°C) or to higher temperatures (for the PB peak at -93°C). How- ever, F i g . 4b shows T, for the PS phase to be identical for the host copolymer and its blend with SBR, so the SBR cannot be residing there (improbable, anyway, for an SBR with 75 percent PB). The PB-rich phase shows a T, shift upward by about 3"C, too slight to be attributed to the full amount of the additive but consistent with a minor amount of the SBR being located there.

By default, then, the explanation of the broad G" peak and the G' dropoff must be that the major portion (and perhaps all) of the SBR resides in the interphase. This is also consistent with the host copolymer having a large interphase (@/ = 66 percent) and the SBR having a low Hut. A further confirmation comes from estimates of T,, for a hypothetical interphase that is uniform in composition(24), the value = 0.515 PS) corresponding to complete incorporation of the SBR in the preexisting SBS interphase. As shown in Appendix 2, these estimates put this T, in the range -30" to + 1"C, which is well within the temperature span of the GI' peak in Fig. 4b. And, if the SBR is attracted into the interphase, we expect @, to increase and both QB and QS to decrease correspondingly. These latter expectations are indeed fulfilled by the G" evidence showing the PB and PS peaks both to drop in magnitude and, although the plateau level for the blend is difficult to evaluate (inas- much as it is not close to horizontal), it seems clear that a reasonably-defined average value would show an increase-and, hence, an increase in @,--for the blend relative to the host copolymer.

c v, I I I I I I I

L i i -

6 SBS (0.48 PSI --- SBS (0.27 PS) Blend with

-........ SBS (0.27 PSI Blend w i t h 44% PS (4,800)

/ 30% PS (4,800) /'

/' 4 --SBS (0.27 PSI - Nonlinear Properties

Blends with PS. Figure 1 showed that the addition of PS homopolymer to 0.27 PS block copolymer elevated both G' and G" in the plateau region, resulting in dynamic mechanical properties similar to those reported previously for the 0.48 PS triblock copolymer. This suggested that PS from either the added homopolymer or the triblock copolymer is equally effective in determining linear properties dictated by small, localized deformations.

Similar conclusions cannot be drawn, however, when large deformations are involved. Figures 5a and 5b show the stress-strain behavior of the first and second loading cycles for two SBS (0.48 PS and 0.27 PS) block copolymers (24) and their blends with homopolymer PS. In the first cycle, the 0.48 PS copolymer exhibits a pronounced yielding phenomenon, while the 0.27 PS copolymer lacks the yielding behavior. Adding enough PS to the latter to change the overall composition to 0.48 PS creates the yielding phenomenon, yet the stress-strain curve is consistently lower than the corrc- sponding 0.48 PS SBS copolymer. This is actually consistent with the G'(T) data at room temperature, which showed the blend to be weaker than the host by about the same ratio as for the nonlinear data. Thus, the added PS is not as effective in supporting tensile stress as the PS block in the copolymer, even though the added PS is mostly dissolved in the PS-rich phase of the system. This observation reveals the importance of the

21.2% PS (10,300) 14

The hfechanical Proper t ies of Styrene-Butadiene-Styrene ( S B S ) Trihlock C o p o l y m e r Blends

POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, No. I 1 679

h

0 a 5 Y

Second Cycle

Strain (%I

F i g . 5. ( a ) Stress-straitt curves of severul SBS copol!jmers and SBS-PS blends in the first loading cycle, at room tentperuture. The 0.48 P S copolymer exhibits a definite yielding phenonze- non, whereus the0.27 PS copolymer lucks i f . Blending the latter with P S homopolymer produces the yielding behavior, yet all of the stress curves ure consistently lower thun for the 0.48 P S copolymer. Adding P S homopolymer to 0.48 P S copolymer greatly increases its low-strain modulus and yield point, hut emhrit- tlenient occurs, and the sample fails at u low elongation, Nom- inal strain rate is 0.5 mirtr'. (b) Thecorre,sponding curaes for the second loading cycle. All samples exhibit rubbery behavior, the yield points having been lost. The blertds possess larger "set," wi th mugnitude increasing wi th P S homopolywier content, than the pure block copolymers.

Joram Diamant, David Soong, und bfichael C . Williams

chemical linkage between the PS and PB blocks of a block copolymer in contributing to its desirable engineering properties. Further increase of the PS content of the 0.27 PS blend to 0.59 PS by blending in more of the homopolymer elevates the stress at a given strain, but the resulting property is still inferior to the 0.48 PS copolymer alone. Addition of PS to 0.48 PS copolymer, which already has a co-continuous lamelEar morphology (29), rapidly increases the yield stress. However, the sample is embrittled, and the ductile behavior disappears.

In the second loading cycle, all samples exhibit rubbery behavior. Note that blends with overall PS con- tents significantly higher (0.59, 0.48) than the parent SBS (0.27) only give stress-strain curves comparable to the pure SBS alone. This indicates that the PS homopolymer, partially effective in supporting stress when incorporated with interconnected styrene domains in the copolymer, becomes totally deficient in this task after the destruction of domain connectedness during the first cycle. Another pronounced distinction for the blends lies in the “set” behavior acquired after the first stretching cycle. If the second stretching cycle is initi- ated immediately upon completion of the first cycle, both 0.48 PS and 0.27 PS block copolymers exhibit a relatively small “set” (strain beyond which stress becomes nonzero). The blend samples, however, display much greater sets, whose magnitudes increase with the PS homopolymer content. This demonstrates that the PS homopolymer component does not possess as strong a memory effect as the SBS copolymer .which, in the blend, is being diluted. The added PS also flows more easily under stress, as it is not chemically attached to the rest of the system.

Further experiments were performed to elucidate the morphology destruction pattern and the role of PS in the blends. The samples were stretched repeatedly, each successive cycle terminating at a higher strain level than the preceding cycle by an increment of 100 percent strain. Figures 6 and 7 show stress-strain curves for two SBS-PS blends, each compared with envelopes of the cyclic maxima for two relevant pure SBS copolymers (the host-either 27 percent or 29 percent PS-and the 48 percent-PS copolymer, which has the same PS content as the blends). Detailed curves for all three pure copolymers have appeared elsewhere (24) and will not be reproduced here; suffice it to say that each cycle brought ahout significant structural varia- tion, with concomitant large hysteresis and increasing set. The upper bounds ofthese curves (see Figs . 6 and 7 ) form envelopes that in all respects are almost identical to the corresponding stress-strain curves in Fig . 5.

The basic hysteresis loops remain in the test results of an SBS (0.27 PS) and PS (M,( = 4,800) blend with overall composition of 0.48 PS (F ig . 6). However, the set in each cycle is appreciably larger than the ones observed for the pure triblock copolymers; after the sixth cycle, the host SBS set is 120 percent (24), while the blend in Fig. 6 has 360 percent. Furthermore, the peak stresses exhibit a downward trend after approximately three cycles, a behavior unique to this blend, as the stress en-

2 ~--&>~/+~/,~ / I ’ /

0 0 100 200 300 400 500 600 700

_-- - -- -

-

-

-

-

’ ’ SBS(O.29PS)

0 100 200 300 400 500 600 700

St ra in (%)

Fig . 7. C! /c / i c strrs.c.-strcciii behcicior (!f ( i bleticl cotttciittirig 0.29 PS copol!/nier with .uufficiritt I i o i i i o ~ ~ ~ ~ / y i i t c r PS ( :V,c = ZO,.POO) to cittciiii u?i o v e r d P S cotiteilt of48%. Cotit1itioii.r. (1.5 i t i F i g . 6. Entielopes ure giverifor t h e c!/clic t i t u x i n i i c (22, 241, fo r the hos t 0.29 PS ccnd 0.48 PS block copolyniers.

velope seems to be approaching that of the 0.27 PS host copolymer rather than the 0.48 PS pure copolymer. This indicates that the blend PS-phase continuity ere- ated by the PS additive (Fig. 1) is ineffective after that continuity is ruptured by large deformations. The load is being borne primarily by the rubbery PB phase of the host copolymer, and this tendency increases as each succeeding cycle further degrades the PS-phase connectivity.

In Fig . 7, the 0.29 PS-SBS copolymer, blended with PS (ai(. = 10,300) to reach an overall PS content of 48 percent, gives stress-strain loops similar to those of the pure block copolymer systems. The set is far smaller than in the previous case: only 160 percent after six cycles, scarcely more than the 0.29 PS copolymer host,


The Mechunicul Properties of Styrene-Butadiene-Styrene (SBS) Triblock Copolymer Blends

which suffered 130 percent (24). The stress envelope observed here is between those of the 0.29 PS and 0.48 PS block copolymer samples and shows the steady increase typical of the latter. Apparently, the added PS is more effective in supporting stress and undergoes less flow than the previous case. This may be due, in part, to the larger molecular weight (MI,. = 10,300) of the PS homopolymer than that used before, even though the added PS is expected to be again dispersed primarily in the PS-rich phase of the block copolymer (PS block M.7

= 16,000). A more likely explanation for the major difference

between the two blend systems in a cycling strain lies in how the additive affects morphology. The highest stress envelope in Figs. 6 and 7 belongs to pure 0.48 PS copolymer, which continues to be strong after rupture of PS-phase continuity because its lamellar microstructure (22, 24) breaks into platelets that have maximum strengthening ability as filler particles. The much. lower envelope for the pure 0.29 PS host in Fig. 7 corresponds to two fiictors, the lower volume fraction of filler and the rodlike nature of the particles, known to be less effective than platelcts. When this host receives extra PS, the rodlike elements are expanded in diameter to the point where they contact and overlap each other, effectively establishing an imperfect lamellar microstructure. This suhsequently is broken up by cycling strain, but the stress envelope is higher than that of the host because of now having 48 percent PS rather than 29 percent and resembles the 0.48 P S copolymer in its strain-dependence because the broken microstructure contains imperfect platelets rather than rods.

The results in Fig. 6 can be explained by analogous arguments. The 0.27 PS host (with smaller M than the 0.29 PS host) is usually characterized as having spherical PS domains. When the PS additive joins these domains, diametral growth is less than in the case of rodlike domains, and the PS-phase continuity, subse- quently, consists of point-to-point contacts in three di- mensions. This is sufficient to produce the yield stress that is seen, but strain-induced degradation causes PS breakup into chunky or spherical particles that have minimum strengthening effect as a filler. Thus, in Fig . 6, the blend envelope is always above that of the host (48 percent vs. 27 percent PS) but exhibits similar strain-dependence because they contain particles having basically spherical shapes.

The blends shown in Figs. 6 and 7 clearly illustrate that overall PS content alone is an almost useless quan- titative indicator of nonlinear properties, even if linear properties seem to suggest otherwise.

Blends with SBR. Stress-strain characteristics of two SBK blends with the same (0.48 PS) host are shown in Fig. 8 , together with the pure SBS properties. Addition of even 1 percent SBK has a remarkably large effect in reducing strength properties. In the first cycle, the yield stress drops by 16 percent (with only 0.3 percent reduction in overall PS content!) and the high-strain parts of the curve by 19 percent. Even in the second cycle, blend properties are reduced by 13 to 15 percent. This constitutes powerful evidence that the SBR is

- SBS ( 0 . 4 8 PS) -- SBS ( 0 . 4 8 PSI Blend with

....... SBS (0.48 PSI Blend with / I% SBR (3,600)

10% SBR (3,600) 6 -

; -0 40 80 120 160 200 240 280

St ra in (%) Fig . 8. S/re.ss-sfrciiri ct i rwr . ofuti SBS copolymer c i t d i ts hlent1.r. with S B R i t i the first arid S B C O ~ K ~ locidirig cyclcs ( i t roortt t cw- perufure , with riomiwil strciiti rute = 0.5 t n i l i - ' . The piire Idock cvpolynier, 0.48 P S , has the highest sfre.r.s ut (I given .strc/iti. The strays decreci.se.7 with iricreasirig citnouiits of added SBR. At 10% b!y i~cij iht o f S B R , the yield behcioior i t i tliefirst loutl i t ig cycle is conlpletel!/ .sli~Jpres.Y"'l.

strongly localized in a critical region of the microstructure. This is in qualitative agreement with inferences drawn previously, from the linear data of Fig. 4 , to the effect that the additive resides almost entirely in the interphase of the host SBS.

Addition of 10 percent SBR reduces strength properties even further, from 26 to 33 percent for various parts of the first- and second-cycle curves. Also, the yield stress vanishes entirely for this case, signifying that PS connectivity has been disrupted by the growth of the interphase; indeed, the SBR additive contributes more volume of the interphase than the nominal volume of the homogeneous PS-rich regions in the host (a.7

= 8 percent; see Fig . 3c) . These results are consistent, also, with the lower G', seen for the blend in Fig . 4a, although the latter represents an even greater reduction (about 56 percent) at room temperature.

The major lowering of the stress-strain curves with only minor amounts of SBK can be explained with the concept of an interphase force barrier (28). Such a barrier, related to gradients of free energy experienced by a segment of one SBS block being drawn through the interphase into an incompatible phase, can be clearly related to +,(x) and its gradients. The presence of SBR concentrated in one region of the interphase could have disproportionate effects on this force barrier. For exam-

POLYMER ENGlNEERlNG AND SCIENCE, MID-AUGUST, 1982, Vol. 22, No. 1 1 68 1

joruni Diumunt, David Soong, und Michael C . Williums

ple, Fig. 3 shows (for a different SBS) that a steep gradi- ent in &(x) exists at 0.25-the SBR average composition-so that localization of more material at that point would stretch the x-coordinate there, reducing dd,,Idx and, hence, the force barrier.

These arguments only apply if actual flow occurs, so that block chains are pulled out of their preferred domains and the reduced force harrier is manifested by a corresponding stress reduction. Indeed, all systems studied here exhibited some flow, as represented by the set. Interestingly, the set in all three samples in Fig. 8 is the same, unlike the PS blends of Fig. 5 . These two results are consistent with each other if the set is caused .by withdrawal of styrene segments from PS phases. Whcn SBR is the additive, it seeks out the interphase and does not affect deformation of the PS phase. When homopolymer PS is the additive, it is incorporated primarily in the PS phase and, therefore, each different blend has a different set.

Although the properties presented in Fig. 8 are those of the solids, the interpretation about SBR reducing stress levels has definite implications in melt processing. SBS block copolymer melts are known to possess excessively high viscosities below T',s, explainable (in part) by the force barrier concept (28). By adding a small amount of SBR polymcric plasticizer, we can expect the melt viscosity to decrease significantly if the above argument for solids holds also for melts. This prediction awaits experimental verification.

ACKNOWLEDGEMENT

Financial support was provided by the National Sci- ence Foundation Polymer Program, Grant DMR 76-83679, during development of the model. Most of the experimental work was supported earlier by ONR Contract P-N 14-75C-0955 under Professor Mitchel Shen (deceased). We are grateful to have collaborated with Professor Shen and acknowledge our sorrow at his passing. The authors would also like to thank Dr. David Hansen for supplying us with the SBS copolymers.

Appendix A

Summary of the Microstructural Model

The microstructural model developed recently (22, 23) to describe linear viscoelastic properties is based on the Nielsen model (33), originally designed to predict the elastic shear modulus of a two-phase composite containing inclusions occupying a volume fraction, Qi, sus- pended uniformly in a continuous matrix. In the present adaptation (i.e., rubbery inclusions in a glassy matrix (PB in PS), the relevant mixing equation is

(1) G - 1 - A + R @ i --

G , 1 + R@i

while A + 1 is the Einstein coefficient for the inclusions and the composition dependence is accommodated by

in which Qmas is the limiting packing fraction for the

inclusions. In addition, G , - Gi

AG, + Gi R = (3)

where G , and G, are the shear moduli of the matrix and inclusions, respectively. For microstructures en- countered in materials used here, realistic modeling re- quires values ofA sufficiently high for the model to be- come entirely insensitive to it. Thus, the only significant model parameter is @mas.

The diffuse interphase is treated as a collection of thin slices, each an SBR, with composition conforming to the local composition profile d,[(x). These slices are then lumped with the PB phase as inclusions in E 9 1 to 3; thus; Qt = @, + @ B . The material properties of each slice are computed by a simple volume-weighting scheme (22-24) from those of PS and PB. Together with properties associated with PS and PB in the major phases, these derived quantities are used to calculate the overall properties of the phase-separated system. Readers are referred to the original development (22, 23) for detailed discussion of model manipulation and systematic strategies for fitting data.

Appendix B Calculation of T , for SBR and Its Blends

All computations below will employ for the homopolymers T,'" = 363°K and T,PB = 180"K, taken from the corresponding SRS G" peaks for PS and PB in Fig. 4b. From these component values, T,, for the SBR random copolymer of styrene fraction d, = 0.25 is expected to lie within a range bracketed by two well-known mixing rules:

Gordon-Taylor linear rule (34),

T f H K = T,Psd, + TLB(l - d,) = 226°K = -47°C

Fox inverse rule {35),

Experimental values (36) for an SBR with d, = 0.235 lie in the range -59" to -64"C, so the calculations above for d, = 0.25 seem to give a plausible bracket.

Preparatory for the subsequent estimates of TL (the interphase average Tg for the SBWSBS blend), it is nec- essary to compute the interphase average for the pure host SBS copolymer, T L h . Since this copolymer has been characterized (23) as having an interphase average composition of 6 = 0.56, use of the two mixing rules gives estimates of Ti,, = 282°K and 251°K.

Finally, we assume that all the SBR additive is contained in the now-expanded interphase, where its fraction is $i.BR = 0.144. Applying the two mixing rules to the interphase only, we obtain T', = 274°K (all Gordon- Taylor values) and 243°K (all Fox values). This range of -30°C to +1"C brackets the midrange Gl'maximum in Fig. 4b almost perfectly, with both values being close enough to the peak center to support the basic conten- tion (see text) within the numerous approximations being employed here.


l h e Mechtrniccil Proper t ies of St~jreii”-Butncliene-St!jr“rie ( S B S ) l r ib lock Copolyriier B1ctid.s

REFERENCES 1. M. Shen and H . Kawui, AIChE J . , 24, 1, 1978. 2. “Slulticompon~tit Polymer Systems,” R. F. Gould, ed.,

Adt-. Claeni. Series, Vol. 99, ACS, Washington DC (1971). 3. “Copolymers, Polyblcnds, and Composites,” N. A. J.

Platzer, ed.,Ado. Chern. Series, Vol. 142, ACS, Washington, DC (1975).

4. “Multiphase Polymers,” S. I,. Cooper and G. M. Estes, eds., Arlo. Cheni. Series, Vol. 176, ACS, Washington, DC (1979).

5. 1). J . hleier, J. Polyt,i. Sci. C., 26, 81 (1969). 6. S . Krause, J . Polyni . Sci. A-2, 7,249( 1969); ~~acronro/ec i i les ,

3, 84 (1970). 5. V. Bianchi, E. Pedeinonte, antl A. Torturro, J . P o l y n i . Sci.,

Part B , 7, 785 (1969). 8. L. Marker, Polym. Prepr., Ani. Chut,i. Soc., Diu. Pol!ytn.

Cheni., 10, 523 (1969). 9. U. G . LeGr;ind, ihitl., 11, 434 (1970).

10. T. Inoue, T. Soen, T. Hashimoto, and H. Kawai, J . Polyni . Sci., Port A-2, 7, 1283 (1970); hiucroniolecules, 3,87 (1970).

11. D. F. Learyantl hf. C. Willianis,J. P o l y r n . Sci., Port B , 8,335 (1970).

12. D. F. Leary and X I . C. Williams,J. Pol!yrn. Sci., Polyni. Phy.~. Ed. , 11, 345 (1973).

13. E. Helfand and Z. R. Wassernian, Pol!yrti. E t i g . Sci., 17, 582 (1977).

14. D. J . hleier, Proc. Polym. Colloq., Kyoto, Japan(Sept. 1977). 15. E. € I . Kerner, Proc. Phys. Soc., B69, 808 (1956). 16. J. C. Halpiii, J. Compos. Mater., 3, 732 (1969). 17. J. E. Ashton, J . C. Halpin, and P. H . Petit, “Primer on Corn-

posite Analysis,” Chap. 5, Technoniic Publishing Co., Staniford, Connecticnt (1969).

18. S. W. Tsai, “Porii~ulas for the Elastic Properties of Fiber- Reinforced Coniposites,” U.S. Dept. of Coiiinrerce Report AD834851 (1968).

20. R. A. Dickie, J . A p p l . Pol!yni. Sci., 17, 45 (195%‘3). 21. R. Ivl. Christensen, J. Mech. Ph1.r. S o l i d s , 17, 23 (1968). 22. J. Diamant, P1i.D. Thesis, Cnivel-sity of California, Brrk-

eley (1982). 23. J . Diamant, D. S. Soong, and M. C . Williams, to he p u b

lished in “Contemporary Topics in Polymer Science,” Vol. 4, W. J. Bailey and T. Tsuruta, eds., Plenum Press, Kew York (1982).

24. J. IXamant, 1). S. Sooiig, and hl. C. Williams, .4NTEC Pro- ceedings, hlerting of Soc. Plastics Engineers, Bostoii (\la!. 1981). Maiirtscript to follow.

25. G. Holden, E. T. Bishop, and N. H. Legge, J. Pol!yvi. Sci., 26C, 37 (1969).

26. 1). F. Leary and hl. C. Williams,]. Polyni. Sci., Ph1.r.. E d . , 12, 265 (1974).

27. D. F. Sleier, J . Appl . Po/!yni. Sci., 14, 427 (1970). 28. C. Hendersoti and hl. c. \villianls, J . P o l y i r i . sci., POl! /J l i .

Lett . Ed., 17, 257 (1979). 29. T. Hashimoto, 51. Fujimura, K . Saijo, H. Kawai, J. Dimxuit,

and hl. Shen, in “Mlultiphase Polymers,” S. I,. Cooper and G. h.1. Estes, eds., A&. Chem. Series, Vol. 176, Am. CIieni. Soc. (1979).

30. P. F. Erhardt, J . J. O’hl;dley, and R. G. Crystal, in “Block Co- polymers,” ed., S. L. Aggarwal, Plennni, New York (1970).

31. G . V. Viiiogradov, V. E. Dreval, A. Y. hlalkin, Y. G. Yanovsky, V. V. Harancheeva, E. K. Borisenkova, M. P. Zahugina, E. P. Plotnikova, and 0. Y. Sal)sai, Rheol . Actu, 17,258 (1978).

32. N . Nemota, K. Okaww, antl H. Odani, Rrt i l . Z t i s t . Chui)i. Rcs. , 51, 118 (1973).

33. I , . E. Nielsen, Rheol. Actu., 13, 86 (1974). 34. hl . Gordoti and J . S. TaylorJ. A p p l . Chein., 2, 493 (1952). 35. T. C . Fox, Bul l . Am. Ph!/,s. Soc., 1, 123 (1956).

19. M. ‘I‘akayanagi, Proc. 4th Internat. Congress Hheol., Part 1, 36. “Pol) nier Hantll;ook,” J . Bratidrup antl I . Immergut, eds., p. Interscieiice, New York, 1965, p. 161. V-8, Wile>-, N e w York (1955).

Correction In the paper “Mechanisms of Flow Induced Crystallization” by A. J. McHugh, which was published in the January 1982 issue of POLYMER ENGINEERING AND SCIENCE (VOL. 22, NO. l), the following correc- tions should be noted:

Page 18 Equation (7) ,K =

0

Page 19 The first line of the first paragraph, “The constant K takes on the value of 4 or 3 . . .” should read “The constant K takes on the value of 4 or 2 . , _”

1 - 4 4 I‘ [ (Af + Afex) kT Page 20 Equation (19) RL = Go exp

Page 21 The second paragraph line 12 “range of 200s-’ to 500s-’ and tube entrance region” should read “range of 200s-’ to 1500s-’ and tube entrance region”

The second paragraph line 14 “the range of 200s-I to ~ O O O S - ~ . Significantly different” should read “the range of 20s ’ to ~OOS-’. Significantly different” Page 22 The second line following equation (24) “symmetry plane, 7 = d 2 , the flow is Z-D shear and K“ should read “symmetry plane, = d2 , the flow is 2-D shear and _K” --

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POLYMER ENGINEERING AND SCIENCE, MID-AUGUST, 1982, Vol. 22, NO. 1 1 683

the mechanical properties of styrene-butadiene-styrene (sbs) triblock copolymer blends with...

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