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31 Mechanisms of Reinforcement of

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Elastomers by Polymeric Fillers MAURICE MORTON Institute of Polymer Science, The University of Akron, Akron, Ohio 44304

Two systems have recently been studied involving elastomeric networks reinforced by finely divided glassy polymeric fillers. One of these was prepared by blending latices of polystyrene, or other glassy polymers, with SBR, and vulcanizing the resulting elastomer. It was found that the tensile strength of the SBR increased with higher filler content, lower particle size, and increase in rigidity of the filler and was virtually unaffected by filler-elastomer bonds. The other system studied consisted of the "thermoplastic elastomers" prepared from styrene-diene-styrene triblock polymers. Here the polystyrene forms small "domains" (~200 A) which act both as network junctions and reinforcing filler. These glassy particles yield at high stresses, thus relieving high stress concentrations and resulting in high tensile strength but large inelastic deformations. T o u r i n g the last few years new light has been thrown on the mechanism of reinforcement of elastomers by particulate fillers, based on studies of the effect of polymeric fillers on elastomers. These studies have i n volved two different systems, i.e., (1) the use of model polymeric fillers (8,9), such as polystyrene prepared by emulsion polymerization, i n S B R vulcanizates, and (2) the "thermoplastic elastomers" obtained from A B A block copolymers (10, 11) where A is a polystyrene block and B is a polydiene block. The morphology of these two systems is shown i n Figures 1 and 2, respectively. Figure 1 shows electron photomicrographs of fracture replicas of S B R vulcanizates containing polystyrene fillers of two different particles sizes, and the existence of the individual polystyrene particles is easily confirmed. Figure 2 shows a schematic of the morphology of a styrene—diene-styrene block copolymers, i n which the formation of a 490 In Multicomponent Polymer Systems; Platzer, N.; Advances in Chemistry; American Chemical Society: Washington, DC, 1971.


31.

MORTON

Figure 1.

491

Reinforcement of Elastomers

Morphology of polystyrene-filled SBR vulcanizates (X 10,000)

Left: R-polystyrene (2200 A) vulcanizate (Vt = 0.25) Right: SBR-yolystyrene (6500 A) vulcanizate (Vt = 0.25) POLYSTYRENE 10 - 1 5 , 0 0 0 M.W.

POLYBUTADIENE

5 0 " 7 0 , 0 0 0 M.W.

POLYSTYRENE 10- 15,000

M.W.

These have actually been observed by electron microscopy and are quite small (10), approximately 300 A i n diameter. Hence, these polystyrene domains virtually constitute a finely divided filler dispersed within an elastic medium, similar, i n principle, to the model polystyrene fillers of

In Multicomponent Polymer Systems; Platzer, N.; Advances in Chemistry; American Chemical Society: Washington, DC, 1971.

492

MULTICOMPONENT POLYMER SYSTEMS


the first system described above. The only distinction between the two systems is that the "network" i n the case of the block copolymers does not involve chemical crosslinking of the elastomer chains, the polystyrene domains themselves acting as network junctures for the elastic chains. The striking similarity i n the morphology of the two systems makes it interesting to compare their mechanical properties which have already been studied extensively (8, 9,10, 11). Model Filler Studies As previously reported ( 8 ) , the effect of incorporating a finely d i vided polystyrene filler into an S B R vulcanizate is to raise the tensile strength b y increasing the stiffness or modulus of the material. This is shown clearly i n Figures 3 and 4 where it is seen that the increase i n tensile strength, at any given temperature, depends directly on the amount of polystyrene filler. The polystyrene fillers i n these figures are presumably not bonded chemically to the rubber. However Figure 4 shows the effect of a chemically bonded filler (SB-10, a styrene-butadiene copolymer) on the tensile strength, and it appears to have a somewhat smaller reinforcing effect than the equivalent polystyrene. The interesting conclusion derived from this work is that the effect of the filler on the vulcanizate strength can still be explained i n terms of the over-all viscoelastic response of the material. This is demonstrated in Figure 5 where it can be seen that a l l the vulcanizates, filled and un-

Figure 3.

Effect of 485-A polystyrene filler on tensile strength of SBR (strain rate 20 inches/min)



31.


MORTON

I

1

I

-2.0

0

i 2.0 (t /a )

Log

Figure 4.

b

493

i 4.0

i

II 8.0

6.0

T

Effect of fillers on tensile strength of SBR. Filled points refer to 25° C and 20 inches/min.

0.2

I

Log

Figure 5.

0.6

0.4

€

0.8

b

Failure envelope for filled SBR vulcanizates. Filled points refer to 25° C and 20 inches/min.

filled, can be plotted on the same "failure envelope," i n accordance with the treatment proposed by Smith (12). Since all the vulcanizates were prepared with the same crosslink density, the effect of the fillers is to increase the viscous component of the network—i.e., to move the failure


494



Table I. Filler

Type

TFE SB-10 DCB PS DC DCA PAB PA

Teflon Styrene/Butadiene (90-10) 2,6-Dichlorostyrene/Butadiene (78-22) Styrene 2,6-Dichlorosty rene 2,6-Dichlorostyrene/Ethyl acrylate (83-17) Acenaphthylene/Butadiene (81-19) Acenaphthylene

M

Properties n

X 10~

b

3 Gel 2.8 1.7 3.0 Gel 0.9

points up the envelope. This effect is, of course, analogous to that of decreasing the temperature and/or increasing the rate of the test. Hence, the presence of these fillers simply shifts the time-temperature response of these vulcanizates to higher stress values. The apparent effect of chemical bonding between filler and rubber, as demonstrated by the behavior of the SB-10 filler i n Figures 4 and 5, is to decrease slightly the tensile strength. However, since this chemical bonding was accomplished by introducing some 10% of butadiene by copolymerization with the styrene, the resultant SB-10 filler had a substantially lower glass transition temperature than polystyrene ( 6 9 ° vs. 1 0 5 ° C ) . T. L . Smith brought to our attention the fact that such a drop in T might actually decrease the rigidity of the filler and thus possibly alter the viscoelastic response of the filled vulcanizate—i.e., decrease the viscous component and hence the strength. g

Hence, it was of interest to investigate any possible effects of the T and modulus of polymeric fillers on the tensile strength of vulcanizates. For this purpose a series of polymeric fillers was prepared by emulsion polymerization, using monomers or mixtures of monomers designed to yield polymers or copolymers of varying T and modulus. The experimental details of the polymerization w i l l be described i n a forthcoming publication. The characteristics of these polymeric fillers are given i n Table I. y

g

The original objective i n preparing emulsion polymers from the 2,6dichlorostyrene and acenaphthylene was to obtain polymeric fillers of higher T than that of polystyrene. It was also presumed that these fillers would not be bonded chemically to the rubber during vulcanization and that the copolymers with butadiene would enable such bonding to be effected. Actually, the polydichlorostyrene and polyacenaphthylene d i d become bonded to the rubber, as indicated by the inability to extract most of the filler by solvents. The final result was that the copolymers with butadiene served merely as fillers of lower T than the above homog

g


31.


MORTON

495

of Polymeric Fillers Diameter D ,A

T„ °C

Young's Modulus, kg cm- X 10-*

1700 530 460 570 560 470 550 380

70 100 105 165 120 155 250

0.33 1.2 1.6 3.5 4.0 4.0 3.4 >10


n

2

polymers. The ethyl acrylate copolymer also served the analogous purpose. In this way a series of fillers was prepared having T values varying from 70° to 250°C. The Teflon filler was obtained as a latex kindly supplied by the E . I. du Pont de Nemours C o . It had a much larger particle size than the fillers prepared in this laboratory, the latter having been designed to fall in the particle size range 400-600 A . However, it was still of interest to study the effect of the Teflon filler in view of the low rubber-filler interfacial adhesion that could be expected. The modulus of the fillers was determined by tensile measurements on films cast from suitable solvents, while the T values were determined by differential scan calorimetry. The T and modulus values agreed well with available published values. There is no direct correlation between T and Young's modulus of these glassy polymers. The latter appear to fall into three ranges of modulus and can be grouped as follows: (1) SB-10 and D C B , (2) PS, D C , D C A , and P A B , and (3) P A . The modulus of the polyacenaphthylene ( P A ) could not be determined accurately because of the extreme brittleness of the film, but it appeared to be greater than 10 kg cm" . The other two groups contained polymers having similar modulus values. The small difference in modulus between the polystyrene and the poly-2,6-dichlorostyrene was quite unexpected in view of the wide dispersity in T values. The characterization of the SBR vulcanizates containing these fillers is shown in Table II. The low values of sol content indicate the presence g

g

g

g

5

2

g

Table II.

Network Density of Filled SBR Vulcanizates

Filler, 10 vol %

None

PS

DC

DCA

DCB

PA

PAB

% Sol, benzene

3.8

7.5

6.6

7.6

5.4

7.1

6.8

Swelling Ratio, q. Benzene Decane

5.4 2.5

5.73 2.44

5.54 2.35

5.35 2.30

5.39 2.35

5.70 2.42

5.85 2.44



496


of chemical bonding between filler and rubber in all cases. The swelling measurements in benzene cannot be taken as a reliable index of the network density owing to the swelling of the bound fillers in this solvent. However, the values in n-decane are valid since the latter is not a solvent for the fillers. O n this basis, the crosslink densities can be taken as reasonably constant for all the vulcanizates. The comparative effect of the polystyrene and poly-2,6-dichlorostyrene fillers on the tensile strength of a polybutadiene vulcanizate is shown in Figure 6. Despite the large difference in T values for these fillers, there is no difference in their effect on the vulcanizate. This is illustrated further by the failure envelope plot shown in Figure 7, where the data points for the two fillers, at equal volume fraction, appear to coincide quite well. The fact that all the points fall on the same envelope is a good indication of the constant crosslink density for these vulcanizates. Thus, the similarity in effect of these two fillers appears to be more related to their similar modulus values. Figure 6 is also useful in demonstrating the difference in viscoelastic response of polybutadiene and SBR vulcanizates. The higher values of tensile strength of the latter, at any given temperature, can obviously be ascribed to the substantially higher T of the SBR since the crosslink densities of the two vulcanizates are similar. g

y

Figure 6.

Tensile strength of polybutadiene with poly-2,6-dichlorostyrene (DC) filler. Strain rate, 2 inches/min.



31.

MOHTON

497


H3i2

Figure 7.

1

5*

1

ClS

1

d5

Failure envelope for PBD with poly-2,6dichlorostyrene filler

The behavior of the SBR vulcanizates containing the series of fillers shown in Table I can be seen in Figure 8. Thus, it appears that the three groups of fillers having different modulus values lead to three different tensile strength curves and that the tensile strengths are directly related to the magnitude of the modulus. Hence the original results (8) reported for the SB-10 b o n d e d " filler can now be explained adequately by the lower modulus of this filler, and this effect also appears to be corroborated by the higher tensile strength of the polyacenaphthylene-filled SBR. In other words, the rigidity of the filler appears to be sufficiently important to affect the viscoelasticity of the matrix and hence the tensile strength. These findings must be considered tentative, pending corroboration by more extensive studies. The Teflon-filled vulcanizates have not been included until now since this filler must be considered as a special case, involving poor adhesion at the filler-rubber interface. The marked difference between Teflon and the other fillers is seen in Figure 9, which shows that the Teflon filler exerts only a slight effect on the tensile strength of the polybutadiene vulcanizate. As a matter of fact, although this filler does increase the strength slightly at temperatures above 0°C, it actually appears to de-


498


TENSILE STRENGTH OF FILLED SBR VULCANIZATES


njM ( y a w i t » tom4 • POtVACCNAmmm.INC XOKtCM.-* V (MCKLOflOf TYRINI ^ O rOCV (ACCNA^NTNYLCNt-CO-MITAOlCK€,«0/tO V 14-4.01*01.-* Q MtYtTVUCNI J Q KLYU,«-MCMI4ftOfTYimt-CO-MTAD1!Nt > L4-l.« MUCM"*

Figure 8.

Figure 9.

Effect of various fillers on tensile strength of SBR

Tensile strength of Teflon-filled PBD. Strain rate, 2 inches/ min.



31.


MORTON

499

crease the strength, to some extent, at lower temperatures. N o explanation of this effect can be advanced at this time. However, electron photomibrographs of this filler, i n latex form, show that the particles are generally nonspherical and have axial ratios greater than 1. Furthermore, it is assumed that this filler does not reinforce the rubber because its poor adhesion leads to early dewetting of the particle surface. Thus, it is difficult to predict what effect low temperature w i l l have on the dewetting and subsequent stress concentration i n the vicinity of such nonspherical particles. The nonreinforcing character of the Teflon filler is also shown i n Figure 10, where it is seen that the points for the filled and gum vulcanizates coincide at any given temperature or strain rate. This filler shows no reinforcement of the strength of the vulcanizate, and this is most readily ascribed to the poor filler-elastomer adhesion, which prevents the filler particle from exerting viscous, energy dissipating forces on the vulcanizate in its vicinity. However, this filler also has a significantly lower Youngs modulus than any of the others, so it might be expected to show the lowest reinforcing effect. Even so, its modulus of 3 X 10 k g c m ' is several orders of magnitude greater than that of the polybutadiene vulcanizate at ambient temperatures (— 10 kg cm" ), and it would hardly 3

2

2

28.

o 1.2

filled points r e f e r t o 5 cm/mirv a t 0 & - 1 5 C . #

0.2

0.4

LOG

0.6

C

0.8

b

Figure 10. Failure envelope for Teflon-filed PBD. Filled points refer to 2 inches/min at 0° and — I5°C.


500


be expected that such a difference i n rigidity of the two phases would not exert some effect on the rupture strength if the adhesion were adequate.


Strength Mechanisms in "Thermoplastic Elastomers" (ABA Block Copolymers) The synthesis and characterization of these A B A block copolymers of styrene and dienes have been described elsewhere (10, 11). Since the polystyrene end blocks aggregate into glassy domains which act as network junctions, the elastic center blocks must virtually represent the "network chains." The polystyrene domains should also act as a finely divided filler. Hence it might be expected that the mechanical properties of these materials could depend on the two basic parameters: polystyrene content and length of center block ("molecular weight between crosslinks"). Some interesting results have already been obtained (10,11) on these polymers, where the effect of the above molecular parameters on the mechanical properties has been studied. Thus, Figure 11 shows the effect of variations i n block length and styrene content on the stress-strain behavior of styrene—butadiene-styrene (SBS) polymers. As expected, the stress levels increase with increasing styrene ("filler") content but are independent of the block lengths. In other words, the center block size does not exert the same influence as the "molecular weight between cross-

EXTENSION

Figure 11.

RATIO

X

Stress-strain properties of SBS elastomers



31.

MORTON


501

links ( M , . ) " i n a normal vulcanizate. This effect (or lack of effect) has already been explained (7) as being caused by the fact that the size of the center block (MW — 60,000) is such as to include a sizeable number ( ~ 1 0 ) of chain entanglements, which really act as network restraints. The tensile strengths of the polymers in Figure 11 are shown as the last points in the curves. These, too, show an increase with increasing styrene content, and this could also possibly be explained as a "filler" effect. However, these samples were made by heat molding i n a press, and the high styrene polymer (40% ) d i d not exhibit its maximum strength unless it was "annealed" (slow cooling, at l°C/min instead of 20°C/min). This indicates a considerable influence of thermal treatment, presumably on the degree of phase separation. In view of this, further samples were always prepared by casting films from solution (11). T w o other aspects of the behavior of these block polymers are noteworthy. The first is demonstrated in Figure 11 by the curve for the 4 0 % styrene polymers. These exhibit an unusual "yield point" at very low strain, after which they show a typical curve for an elastomer. This yield point occurs only during the first draw and is not reproduced unless the sample is remolded or reheated. This behavior has been ascribed to the presence of a "continuous" polystyrene phase at such high styrene levels, and this has actually been observed by electron microscopy ( I , 6). The interdomain contacts are, of course, disrupted on the first extension and cannot reform unless the material is reheated. The second aspect that bears mention is the severe distortion of the polystyrene domains which occurs, especially at high strain. This is reflected in an unusually high "set," or unrecoverable deformation, that these materials show at the breaking point or even before. Figure 12 shows this deformation as " % set"—i.e., percent increase over original length, as a function of styrene content and strain. Thus the 4 0 % styrene polymers show as much as 50-60% set near the breaking point. This irreversible extension is completely recoverable upon heating the samples near the T of polystyrene ( — 1 0 0 ° C ) , and it is thus reasonable to ascribe it to an actual distortion, or "cold drawing" of the polystyrene domains. Such distortions have been demonstrated by electron photomicrographs (1,6) of strained polymers, which show the polystyrene domains capable of assuming shapes of high axial ratio. It is this ability of the glassy polystyrene domains to yield under stress that has been invoked (10, 12) to account for the high tensile strength exhibited by these materials. In other words, the polystyrene domains do not act in a similar manner to the above model fillers i n reinforcing vulcanizates—i.e., by raising the viscosity of the matrix, but instead they offer "energy sinks" for the dissipation of the strain energy by being capable of yielding. Hence, the strength of these block polymers y


502

MULTICOMPONENT POLYMER SYSTEMS 70 60 50 H

40


UJ

55 30 20 10 0

J=L

2

3

4

5

EXTENSION RATIO Figure 12.

6

7

8

(X)

Percent set vs. extension for SBS elastomers

should depend on the extent to which the domains can retain their integrity and still absorb energy. In this regard, the behavior of the styrene-isoprene—styrene (SIS) block polymers is illustrative. A n SIS series of polymers is shown in F i g ure 13, and these can be compared with the analogous SBS polymers shown in Figure 11. Although the stress levels of the SIS polymers depend mainly on the styrene content and not on block size (as expected), the tensile strengths appear to be independent of both, with one exception. Unlike the SBS polymers, the tensile strengths are very similar for all the SIS polymers except for the one case involving the shortest end block ( M W = 8400). The constancy of tensile strength of the SIS polymers, above a certain minimum end-block size, can be explained best on the basis of efficiency of phase separation. Since the latter depends on the three parameters—i.e., incompatibility of the two blocks, composition ratios, and block size—the SIS polymers must undergo a much better phase separation at equivalent composition and block size than the SBS polymers, and the latter require a higher styrene content and a longer polystyrene end block to accomplish good phase separation. This is in accord with the generally accepted fact that polyisoprene is more incompatible with polystyrene than is the case for polybutadiene. The importance of efficient phase separation on the "purity" of the domains is illustrated even better in Table III, which shows the tensile strength of the SIS polymers as a function of composition and end-block



31.

MORTON


503

length. The values at higher styrene content were taken from Figure 13, but the striking results are those in which the polystyrene end block drops from a MW of 8400 to 7000 and finally to 5000, where it can be assumed that a separate polystyrene phase no longer exists. Hence the "critical" value for the molecular weight of the polystyrene is in the vicinity of 6000. If the strength of these block polymers depends on the ability of the polystyrene domains to absorb strain energy by a yielding process, it should be possible to design a material of higher strength by creating domains capable of absorbing more energy—i.e., having a higher modulus. Such domains could possibly be formed, for example, from end blocks having a higher T . In accord with this approach, such a block polymer has been synthesized (5) in which polyisoprene is the center block and g

6

4

EXTENSION

Figure 13.

8 RATIO

-

10 X

Stress-strain properties of SIS elastomers

Table III.

Tensile Strength of SIS Block Polymers

Polymer

Styrene, wt %

MW, X 10~

40 40 30 20 20 19 11

21.1- 63.4-21.1 13.7- 41.1-13.7 13.7- 63.4-13.7 13.7-109.4-13.7 8.4- 63.4- 8.4 7.0- 60.0- 7.0 5.0- 80.0- 5.0

SIS-5 SIS-2 SIS-1 SIS-3 SIS-4 SIS-9 SIS-10

Tensile Strength, z

kg

cmr

2

310 306 321 270 160 22 0


504


poly-a-methylstyrene constitutes the two end blocks. The stress-strain properties of this polymer are compared with those of an equivalent SIS polymer in Figure 14; the stress levels of the a-methylstyrene polymer are consistently higher than those of the SIS polymer, up to and including the point of tensile rupture. Furthermore, since the poly-a-methylstyrene also has a much higher T ( 1 7 0 ° C ) than polystyrene ( 1 0 5 ° C ) , the corresponding block polymer shows a much better temperature-strength retention, as shown in Figure 15. These results with the a-methylstyrene block polymers are especially interesting in that they indicate that end blocks of higher T and rigidity are still capable of yielding under stress and thus absorbing energy rather than undergoing brittle fracture. It would, of course, be instructive to determine to what extent this yielding behavior prevails as the T and rigidity of the domains increase. The above data on these block polymers have thrown great emphasis on the critical role of the " r i g i d " domains in the strength mechanisms of these elastomers. It is interesting, therefore, to explore the effect of making "composite" domains by adding, for example, free polystyrene to these block polymers. The techniques w i l l be described in detail in a forthcoming publication, but suffice it to say that all blending of these additive polymers was done in solution prior to film casting. The properties of the blended polymers are shown in Table I V for both SIS and SBS block polymers. The added polystyrene was designed


g

g

g

500 M O L . W. (xl(T ) 8

2

Figure 14.

4

6 EXTENSION

% S

OR

(X-MS

8

RATIO

- X

10

Stress-strain properties of a-methylstyrene block polymers



31.

0

I

I

I

i

-40

0

'

i

Figure 15.

»

40

TEMP.

»

»

I

80

1

1 160

120

(°C.)

Tensile strength vs. temperature of a-methylstyrene block polymers Table IV.

Polymer Added

505


MORTON

Wt. %

Effect of Added Homopolymers Total Styrene Content, %

Stress at X = 4y kg cmr 2

Tensile Strength, kg cm' 2

None PS-15,000 PS-15,000 PS-15,000

— 5 10 20

S I S - 3 : 13,700-109,400-13,700 20.0 17.5 22.8 24.0 25.5 25.4 30.2 33.5

296 306 302 288

None PS-12,500

SIS-20: 13,700-109,400-13,700 — 20.0 — 14.9 30.0 —

127 342

None

—

S B S - 1 7 : 13,700-63,000-13,700 30.0 —

313

to have approximately the same molecular weight as the polystyrene i n the end blocks. A l l of the resulting cast films were optically clear, except for the blend having 2 0 % added polystyrene, which showed some haze. This can be taken to indicate that the added polymer was able to be included i n the polystyrene domains, which are, of course, too small to cause noticeable light scattering. It is not surprising, therefore, that the added polystyrene causes no change i n the strength of the SIS polymer since, at these styrene contents



506


and molecular weights, the tensile strength has been shown (Figure 13) to be independent of these parameters. The moduli, however, do show an increase with increasing styrene content, as expected. The significant aspect of these results is the fact that the added polystyrene does apparently enter the domains and does not form a separate polystyrene phase, as would be the case i n any attempted blend of the two homopolymers, i.e., polystyrene and polyisoprene. T h e same "compatibility" of added and block polystyrene is demonstrated i n the case of the SBS polymer, only here there is an effect on the tensile strength caused by the increased styrene content, regardless of whether the polystyrene is present as an end block or as added homopolymer. Again the evidence suggests that the added polystyrene enters the end-block domains, and further that i t actually improves phase separation, presumably by virtue of the higher styrene content of the mixture. The above experiments, besides demonstrating the phase separations which occur in these blends, can also be used to show the possible effect of synthesis problems on the properties of these block polymers. In other words, it is obvious that any homopolystyrene that is formed during the block polymerization, e.g., b y fortuitous termination arising from impurities, would have little effect on the tensile strength (provided it is of a suitable molecular weight (11) to be "compatible" with the polystyrene end blocks i n the domains). In this connection it is of interest to explore this approach one step further—i.e., to determine the effect of the presence of diblocks since these also could be the result of adventitious termination i n the synthesis process. Such a study has been carried out b y preparing solution blends of SB diblocks and SBS block polymers, having similar block sizes, and the results are shown i n Table V . Even minor amounts of the diblock polymer have a profound effect on the tensile strength. The fact that the modulus is not influenced to any extent again indicates that the latter depends primarily on the styrene (or "filler") content. However, the presence of the SB diblocks apparently introduces "network defects" which can act as flaws to initiate rupture. Table V . Added Diblock, % 0 1 2 5 67

Effect of Added Diblock Polymer Stress at X = 4, kg cmr 2

46.3 50.3 47.5 48.3 14.6

Tensile Strength kg cm~

0

% Set

2

319 308 264 244 49

50 50 50 50 200

SBS-5:21,100-63,400-21,100; SB diblock: 21,100-63,400.

a



31.

MORTON

507


THe network structure of these block polymers can also be altered by actually crosslinking the elastomeric polydiene chains, thus introducing a "chemical" network. In this approach it was necessary to use a crosslinking method which would not result i n any measurable chain scission. Dicumyl peroxide ( D i c u p ) was chosen for crosslinking an SIS polymer since this peroxide is known to accomplish exclusive crosslinking without any observable chain scission ( 3 ) . The D i c u p was dissolved i n a T H F solution of the polymer, and a cast film was prepared which was then vulcanized in a press at 155 °C for 35 min. A control sample, without Dicup was treated in the same way. Table V I .

Properties of SIS Polymers Crosslinked by Dicumyl Peroxide

Polymer SIS-2A" Vulcanized SIS-2A «MW 6

Swelling Ratio («-)

Stress at X = 4, kg cm'

Tensile Strength, kg cm.-"

X at Break

9.8 5.4

50 53

225 160

11 10

6

2

1

= (13.7-41.1-13.7) X 10

3

Isooctane, 48 hour at 25 C. a

The effect of such crosslinks is shown i n Table V I . Whereas the modulus was affected only slightly by the number of crosslinks introduced, the tensile strength was decreased markedly. Hence, although the actual number of crosslinks produced by the peroxide is not very great relative to the number of chain entanglements already present, these "fixed" crosslinks must introduce sites of stress concentration which cannot transmit the stress to the polystyrene domains as efficiently as the entanglements do. This conclusion is, of course, based on the premise that there is a negligible amount of chain scission occurring during the crosslinking reaction. Note should be taken of the fact that although the crosslinking d i d not raise the modulus much, it d i d reduce the swelling markedly. H o w ever, this anomalous lack of agreement between modulus and swelling may be ascribed to the unusual swelling behavior of these block polymers. This has already been noted (2) as "swelling creep" which occurs in presence of a specific solvent for the elastomer phase—e.g., isooctane—and renders swelling equilibrium difficult to obtain. Finally, it is instructive to compare the temperature effect on the tensile strength of the SBS and SIS block polymers. As noted previously (Figure 6) the tensile strength of an elastomer vulcanizate can be related to the difference between the test temperature and the T of the elastomer, in accordance with the viscoelastic theory of tensile strength. Since the T values for polyisoprene ( —65°C) and polybutadiene ( —95°C) differ y

g



508


-0.5

-

1 -80

i

I - 4 0

I

I 0

I

I 40

1

1

1

80

!!0

TEMPERATURE (°C.)

Tensile strength vs. temperature for SBS and SIS polymers. Strain rate, 2 inches/min.

Figure 16.

by 30°, one would expect the SIS polymers to have a consistently higher tensile strength at any given temperature and strain rate. However, this has already been shown not to be true at room temperature. Figure 16 demonstrates that the strength-temperature relations of the SIS and SBS polymers are virtually identical, again indicating that it is the response of the polystyrene domains i n relieving the stress concentration which governs the rupture of these materials. Since the modulus of these domains is several orders of magnitude higher than that of the elastic matrix surrounding them, the viscosity of the matrix contributes little, if anything, to the tensile strength. Summary and Conclusions The two systems discussed above demonstrate two mechanisms whereby the tensile strength of elastomers can be reinforced by the presence o f ' r i g i d " fillers. For the polymeric fillers dispersed within a vulcanizate, the filler operates b y raising the viscosity of the matrix, analogous to a decrease i n temperature, but without affecting the dynamic, high frequency response (there is ample experimental evidence of the independence of T on presence of filler). There is also some indication that the rigidity of the filler affects the extent of reinforcement. In the case of the heterophase systems resulting from the A B A block polymers, the strength is reinforced because of the ability of the plastic, g


31.

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509


glassy domains to yield and absorb energy, thus delaying the onset of crack propagation. This is undoubtedly caused by the efficient transmittal of the applied stress through the elastic chains which are bonded to the domains, as well as by the presence of some phase mixing at the interface. In this way, the strength reinforcement increases if the energy of distortion of the domains is greater—e.g., for domains of higher T . The absorption of energy b y the elastic matrix does not seem to be significant in these systems. These studies of the A B A block polymers also help to explain the results obtained by many investigators, as described i n a recent comprehensive review (4).


g

Acknowledgments The results described here are based on the excellent experimental work carried out during the past several years b y the following contributors: J. C . Healy, R. L . Denecour, J. Trout, J. E . McGrath, P . C . Juliano, F . C . Schwab, and C . R. Strauss. The interest and collaboration of L . J. Fetters is also gratefully acknowledged. Various aspects of this work were supported b y the Cabot Corp., B. F . Goodrich C o . , and the Materials Laboratory, U . S. A i r Force. Literature Cited (1) Beecher, J. F., Marker, L., Bradford, R. D., Aggarwal, S. L.,J.Polymer Sci., Pt. C 26, 117 (1969). (2) Bishop, E. T., Davison, S., J. Polymer Sci., Pt. C 26, 59 (1969). (3) Calderon, N., Scott, K. W., J. Polymer Sci., Pt. A 3, 551 (1965). (4) Estes, G. M., Cooper, S. L., Tobolsky, A. V., J. Macromol. Sci., Rev. Macromol. Chem. C4 (2), 313-366 (1970). (5) Fetters, L. J., Morton, M., Macromolecules 2, 453 (1969). (6) Hendus, H., Illers, K. H., Ropte, E., Kolloid-Z.Z., Polymere 216-217, 110 1967, MacLaren, London, 1968, p. 175. (7) Holden, G., Bishop, E. T., Legge, N. R., J. Polymer Sci., Pt. C 26, 37 (1969). (8) Morton, M., Healy, J. C., Denecour, R. L., Proc. Intern. Rubber Conf. 1967, 175 (1968). (9) Morton, M., Healy, J.C.,Appl. Polymer Symp. 7, 155 (1968). (10) Morton, M., McGrath, J. E., Juliano, P.C.,J.Polymer Sci., Pt. C 26, 99 (1969). (11) Morton, M., Fetters, L. J., Schwab, F. C., Strauss, C. R., Kammereck, R. F., "SRS-4," p. 1, Rubber and Technical Press, London, 1969. (12) Smith, T. L., Dickie, R. A., J. Polymer Sci., Pt. C 26, 163 (1969). RECEIVED August 3, 1970.


Mechanisms of Reinforcement of Elastomers by ... - ACS Publications

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