The effect of cyclic stress on the physical properties of a poly(dimethylsiloxane) elastomer

The Effect of Cyclic Stress on the Physical Properties of a Poly(Dimethylsi1oxane)

Elastomer JOHN J. FITZGERALD, ARTHUR C. MAIYTELLOCK,

PAUL L. NIELSEN. andROBY" V. SCHILLACE

Eastman Kodak Company Rochester, New York 14650

The dynamic creep behavior of a filled poly(dimethylsi1oxane) elastomer was studied under cyclic stress. The stress level was chosen such that the increase in the internal temperature was small and that microcracks were not observed. This work has demonstrated that cyclic stress in combination with high temperature accelerates the degradation of the elastomer. The results suggest that because of the applied force, breaks in the load-bearing chains of the network occur. These breaks, while relieving the mechanical stress, create highly reactive ionic fragments. I t is believed that because of the subsequent reactions of the ionic fragments, changes in the specific gravity, storage modulus, effective crosslink density, and length of the sample (creep) are observed. The observed decrease in the storage modulus is thought to occur because of the reaction of the ionic fragments with moisture, which results in the formation of silanol chain ends that reduce the effective crosslink density. The results also show that contrary to the prediction of the Boltzmann's Superposition Principle, the rate of creep is greatly enhanced when the sample is subjected to a sinusoidally varying dynamic load as compared to a comparable static load. The polymer weight loss was found to be linear with time and strongly dependent on the level of applied dynamic and static force. In addition, the weight loss and rate of creep were also found to be strongly dependent upon temperature.

INTRODUCTION

hanges in stiffness or mechanical strength can C occur when a filled elastomer is subjected to cyclic stress. The failure of the elastomer can be caused by heat, ozone, or mechanical tearing. Heat buildup can be particularly severe for elastomers that exhibit low resilience since they are not able to efficiently dissi- pate the heat generated. Therefore, under cyclic stress, the internal temperature of the elastomer will rise, and with sustained exposure to high temperature the elastomer can thermally degrade by oxida- tion, reversion, or fatigue ( 1, 2). Mechanical failure of the elastomer can begin at small flaws or nonhomc- geneties. These flaws act as stress raisers, and cracks can begin at a sites where the local stress exceeds a critical level (3-5).

The use of filled poly(dimethylsi1oxane) (PDMS) elastomers for high temperature applications is necessary for many technologically important applications. The addition of filler substantially increases the tear resistance and tensile strength of these materials. The literature suggests that these materials are thermally stable and can withstand hot air aging

for 7500 h at 205°C before failing (6). In some applications, filled PDMS elastomers are subjected not only to high temperature, but also to a cyclic load. How- ever, there have been no published studies concerning the changes in physical properties that occur when a filled PDMS elastomer is subjected to cyclic loading at elevated temperature.

We have studied the response of a filled PDMS elastomer to cyclic stress. In this study, the increase in the internal temperature of the elastomer was small and the applied force was not so severe that microcracks were observed. The results indicate that contrary to the predictions of Boltzmann's Superposi- tion Principle (7), the rate of creep for a sample is faster when the sample is subjected to a sinusoidally varying load as compared to when the load is applied nearly continuously. This unique behavior has been reported previously for both filled natural rubber (8) and filled butyl rubber (9). To our knowledge, this is the first time this behavior has been observed for filled PDMS elastomers. In the previous studies the failure of the elastomer was attributed to a stress- softening (8). In this study, it was found that cyclic

1350 POLYMER ENGINEERING AND SCIENCE, SEPTEMBER 1992, Vol. 32, No. 18

Effect of Cyclic Stress

stress and elevated temperature accelerated the degradation of the filled PDMS elastomer, causing molecular changes that produced changes in the storage modulus, effective crosslink density, specific gravity, and length of the sample (creep). The changes in the storage modulus were influenced by the reaction with moisture and the loss of volatile cyclic products. The polymer weight loss was shown to be linear with time, and, as with the creep behavior, it was strongly affected by the stress level and temperature. In future reports, it will be shown how the concentration of filler, frequency, crosslink density, and atmo- sphere influence the response of the filled elastomer to cyclic stress at elevated temperature.

BACKGROUND

It is believed that the stress-induced changes in the physical properties of this filled elastomer are the result of chemical changes. Therefore, it is useful to review literature concerning the chemical degradation of PDMS. In the last forty years, there have been a number of studies concerning the thermal stability of linear PDMS (10-14). These studies have shown that the limiting factors controlling the thermal stability of PDMS are often either its oxidative stability or thermal rearrangement reactions (chain scission) (6). While the oxidative stability of PDMS can be improved with the addition of certain metal salts (10). the thermal rearrangement reactions are not so eas- ily controlled. In fact, these reactions can take place at elevated temperatures, either spontaneously or catalyzed by acids or bases. In thermogravimetric analysis experiments, it has been shown that volatile products typically do not evolve until approximately 340°C. The volatile products include cyclic trimers (73%), tetramers (13%), pentamers (4%), hexamers (6%), and higher molecular weight cyclics (1 1).

It has also been shown that the thermal stability of linear PDMS is adversely influenced by the presence of moisture. Lewis (1 5) has reported that water causes the random depolymerization of PDMS through the hydrolysis of the siloxane bonds. The reaction pro- ceeds quickly in the presence of both acidic and basic catalysts. Moreover, even without a catalyst present, the hydrolysis reaction will still proceed at elevated temperature, but at a slower rate.

The influence of PDMS fluid on the stability of crosslinked filled PDMS elastomers has been studied (16, 17). Henry (16) studied the high temperature degradation of aluminum and iron oxide filled silicone elastomers submerged in a high viscosity PDMS oil. A kinetic expression that involved the 3/2 power of the effective chain concentration was proposed to describe the degradation of the elastomer. Further, Henry found that the oil adsorption was accompanied by a decrease in the effective crosslink density of the samples, and it was concluded that the increased oil adsorption was associated with the chemical degradation of the elastomer.

Rice, et al. (17), studied the degradation of silica filled PDMS elastomers in the presence of Dow Corn-

ing 200 PDMS oil at temperatures b=tween 180 and 200°C. They concluded that the oil absorbs by two concurrent processes. The first process involved the rapid and reversible diffusion of ail into the elastomer. The second process was much slower and involved the irreversible chemical degradation of the elastomer. The authors also concluded that while the diffusion of oil did not appear to produce chemical changes in the elastomer, the elastomer was de- graded by the slower chemical process, which in turn promoted the adsorption of additional amounts of oil.

Stress relaxation can occur in these elastomers even at low temperatures. Stein and Prutzman (18) have recently studied the stress relaxation behavior of a PDMS room temperature vulcanized elastomer. Their work demonstrated that the tin crosslinking catalyst, together with water, produced siloxane bond rearrangements that resulted in chemical stress relaxation. An apparent activation energy of 10.3 kcal/mole was calculated for this process.

EXPERIMENTAL

Sample Preparation

A PDMS elastomer filled with 35 vol% aluminum oxide (Alcoa T-61) was studied. The filled elastomer was prepared by condensing 5 pph of tetraethoqsi- lane (TEOS) with a silanol terminated PDMS fluid, purchased from Petrarch, that had a viscosity of 0.750 Pa.s. The aluminum oxide filler was dispersed in the PDMS and TEOS with a three-roll mill. To the milled product, 0.50 pph of dibutyltin diacetate was added. The sample was well stirred and degassed before injection into a mold. The samples stood at room temperature for at least 48 h before demolding and then were post-cured for 16 h in a circulating air oven at 205°C.

Cyclic Stress Experiment

The samples were exercised with a dynamic mechanical analyzer (19) by repetitively exerting a com- pressive force. Typically, the static load was kept constant at 8 kg and superimposed on this was a sinusoidally varying dynamic load of 4 kg rms. Unless otherwise noted, the frequency was 30 Hz, and the temperature was held at 218°C. The initial samples were cylindrical (12 mm in length and 12 mm in diameter) and were run in an ambient air atmo- sphere. Depending on the applied load, the dynamic strain vaned from 4.9% to 1.5% during the course of the experiment. The storage modulus was not cor- rected for the compliance of the instrument.

Effective Crosslink Density

Because of the presence of filler and its influence on swelling, it was possible to calculate only an a p parent or effective crosslink density. To calculate the effective crosslink density, we used a modified ver- sion of the Flory-Rehner theory (16,20). The method outlined by Bueche (2 1) and Bobear (22) was used to extract and subsequently swell the samples after they

POLYMER ENGlNEERlNG AND SCIENCE, SEPTEMBER 1992, Vol. 32, No. 18 1351

John J. Fitzgerald, Arthur C. Martellock, Paul L. Nielsen, and Robynn V. Schillace

L

-1.0

- 0.8

- 0.6

- 0.4

were subjected to cyclic loading. The samples were extracted with toluene by placing them in a vial with an excess of toluene for at least 17 h at 50°C. The toluene was then discarded, fresh toluene was added, and the samples were allowed to equilibrate an additional 8 h at room temperature. The samples were then removed from the toluene, quickly blotted to remove the excess toluene, and the swollen weights of the samples were recorded. Studies showed that after 8 h no more toluene was imbibed by the sample. Each sample was then dried to constant weight in a vacuum oven at 110°C. The samples were weighed to the closest milligram and their specific gravities were determined (23).

RESULTS AND DISCUSSION Modulus Calculation

The response of a filled PDMS elastomer to dynamic creep was studied in order to understand how this material would behave when subjected to cyclic stress at elevated temperature. As shown in Fig. 1, there is an approximately 25% increase in the storage modulus and a 60% reduction in the height of the sample after 60 h. Compared with previous stress relaxation or dynamic stress relaxation experiments on PDMS elastomers, the observed changes are much more pronounced (18, 24). In a typical stress relaxation experiment, the sample is under high stress for only the first few moments; by studying dynamic creep, it is possible to obtain a higher average stress level during the course of the experiment. A disad- vantage is the lack of a simple theory [like Tobolsky's (25)] to mathematically describe the effects. It has been shown that in a stress-relaxation experiment, when the normalized decaying force log (F/F,), is plotted vs. linear time, a straight line is observed. Its

n lu n a W

I . I ! 0.2 0 20 40 60 ao

Time (hours) Fg. 1. The effect of cyclic stress on a PDMS elastomerJlled with 35 uol% aluminum oxide. A) Storage modulus and B) fractional length us. time.

slope has been shown to describe the rate of polymer chain breakage (24, 26).

FJF, = NJN, = exp- k t ( 1 )

where NJN, is the fraction of stress-supporting network chains remaining at time t. However, Ferry has shown that stress relaxation and creep are intercon- vertible (27). so in principle, at least, all of Tobolsky's insight applies.

Figure 1 shows the results of a typical 60 h isother- mal fatigue experiment wherein both changes in the length (creep) and modulus of the elastomer are o b served. However, the calculation of the scalar value of the complex modulus (E*) , and the derived storage modulus ( E') requires discussion.

The relationship developed by Payne and Scott (28) that is shown in Eq 2 is often used to calculate the complex modulus of an elastomer (E*) . This equation assumes that the elastomer 1) is bonded to the compression platens and 2) is perfectly elastic (no viscous flow):

Fd

A,(a, - a;2 -a, + G2) E* = 3

In Eq 2, Fd is the dynamic force (peak to peak), A, is the initial cross-sectional area, a, and a, are the compression ratios at the extremes of the cyclic deformation, and S , is an empirical shape factor. Based on the fact that the samples were cylindrical, and the expected modulus was between 4 and 6 MPa, Eq 3 was chosen to define the shape factor.

So = 1 + 0.063( d,/h,)2 (3) In Eq 3, d, is the original sample diameter and h, is the original sample height (28). The function of the shape factor is to correct for the apparent hardening of thin samples of an elastomer. In compression, when the diameter of a plug becomes much larger than the height, the ordinary equations for modulus give results that are too high (28). The shape factor, roughly the area under load to the free area, is an empirical correction for this effect.

In this case, it is obvious that the second assump tion in Payne and Scott's equation is not valid since the experiment is designed to study creep, which is equivalent to viscous flow in the terminal region. Initially the height and diameter of the sample are roughly equal: however, during the course of the 60 h experiment, the sample flows into a stable shape wherein the diameter is nearly four times its height. Clearly, a new method of calculating the storage and loss moduli is needed.

These new calculations were developed step-wise, without any untested assumptions. First, the simple definition of dynamic modulus was used, stress ( Fd/A,), dynamic force/initial area at temperature divided by strain ( hd/h,), the dynamic deformation divided by the thermally expanded height. The results of these calculations give the top curve in Fig. 2.



20

15 - m n 3 v) 3 = 10 '0 0 E

-

A

0 3 1 1 10

Time (hours) Q. 2. "his Figure illustrates the step-wise development of the calculation of the storage modulus. A) Simple equation: B) simple equation plus jlow; C) simple equation, JOW, and shape factor: D) Payne and Scott equation (27); E) simple equation, Bow, and changing shape factor: and F) simple equation,Jow, changing shape factor, and weight loss.

However, this simple equation for calculating the modulus also assumes there is no flow, so the first modification was to correct for flow by recalculating the area at each measurement ( A,, assuming a constant volume and that the sample remains cylindrical). This results in an increased time until the u p turn in the storage modulus occurs (Fig. 2 s ) .

A third modification is to correct for the changing shape and its influence on the apparent modulus. As shown in Eq 3, this is done by dividing by an empirical shape factor based on the original dimensions (28). As seen in Fig. 2C, division by the shape factor reduces the initial modulus of the filled elastomer and increases the time to the upturn in storage modulus. Since the shape factor is based on the initial sample dimensions, it is necessary to recalculate the shape factor at each time ( t ) as though it were a new sample. Now the measurements are referenced to time t.

S , = 1 + 0.063( d,/h,)' (4)

Note that this curve, shown in Fig. 2E, gives almost the same curve as calculated from the Payne and Scott equation (seen in Fig. 2D) above, despite totally different assumptions about flow and the use of a different shape factor.

The final adjustment is to incorporate the weight

loss data into the volume calculation. As seen in Fig. 3, both the static and dynamic weight loss are linear with time. Since no voids or microcracks could be observed in the samples, it was possible to show that the entire weight loss could be accounted for by the loss of polymer. This was simply done by assuming a specific gravity for the polymer of 0.97 and 3.9 for the filler, and calculating the change in specific gravity of the sample before and after exercising. Therefore, it is possible to calculate a gradually diminishing volume and to use this information to adjust the cross- sectional area of the sample as a function of time. In addition, Fig. 3 clearly shows the importance of these dynamic creep experiments in that the weight loss is about one order of magnitude higher than under static conditions during the 60 h experiment.

The final equation (Fig. 2F) for the scalar represen- tation of the dynamic modulus is:

E* = E J S , (5 )

E, = Kh,/A, (6)

K = Fd/hd (7)

A, = rr dB/4 (8)

where the apparent modulus

and the stiffness is given by

and the area at time t is,

where d , is the diameter at time t, calculated by taking into account the weight loss. Then the storage modulus and the loss modulus are given by

E' = cos(arctan 6 ) E* E = sin( arctan 6 ) E*

(9)

( 10)

where 6 is the phase angle. The observed changes in the storage modulus and

length of the filled elastomer are believed to be due to chemical changes. It is believed that the cyclic stress and high temperature caused breaks in the load- bearing chains in the polymer molecular network.

10 I

8 -

6 -

4 -

A

I 0-

I

0 20 40 60

Time (hours) Fig. 3. Percent weight loss us. time for the PDMS elastomer, A) heated at 218°C and B) subjected to cyclic stress at 218°C.

POLYMER ENGINEERING AND SCIENCE, SEPTEMBER 1992, Vol. 32, No. 18 1353

John J. Fitzgerald, Arthur C. Martellock, Paul L. Nielsen, and Robynn V. Schillace

Such breaks have been reported in static stress relaxation experiments on silicones by Tobolsb, et aL (291, and by Osthoff, et a1. (30). The breaks relieve the mechanical stress and at the same time create ionic chain fragments that are highly reactive. A similar process is believed to occur in these dynamic fatigue experiments.

The changes in the storage modulus and length of the sample are associated with the fate of the reactive ionic fragments. If the ions react with neighboring molecular chains and the neighboring chain is load- bearing, the new cut will appear in the creep measurement as a change in length. In principle, there would be no change in either the effective crosslink density or the storage modulus of the sample if this occurred.

The increase in the storage modulus at longer times is believed to be partially due to the loss of volatile cyclic products. This type of reaction is known to occur when a reactive ionic fragment “backbites” onto itself, liberating a low molecular weight cyclic oligomer (8). These cyclics, as described earlier, are typically trimers, tetramers, and pentamers that can diffuse to the surface and evaporate. This results in a loss of polymer, which decreases the total polymer fraction in the sample and increases its specific gravity and storage modulus. From measurements of the specific gravity it is possible to show that the volume fraction of filler increases from 35 to 42 vol% during the course of the 60 h cyclic stress experiment. In a future paper it will be shown that the Smallwood- Guth equation can be used to accurately model the observed changes in modulus of the elastomer (7).

In an experiment in which a series of samples were fatigued and their effective crosslink densities then calculated, it is seen that there is an initial decrease in the effective crosslink density (Fig. 4) of the elastomer. This may be due to the reaction of the polymer chains with moisture. The moisture, which may be liberated from the surface of the filler, is present in

400 I I

0 15 30 45 60

Time (hours) Fg. 4. Effective crosslink density us. time for the PDMS elastomer subjected to cyclic stress at 21 8°C.

sufficient amounts to lead to the hydrolytic degradation of the elastomer. The reaction usually yields polymers with silanol chain ends (1 5). This lowers the effective crosslink density, and therefore the elastic modulus, of the elastomer. As seen in F@. 4, the effective crosslink density reaches a minimum after about 20 to 22 h. This corresponds to the time when the sample begins to exhibit an increase in its storage modulus (see Fg. I). Therefore, it is believed that while the reaction with moisture acts to lower the effective crosslink density and storage modulus, con- currently there is a loss of polymer that causes an increase in the effective crosslink density and storage modulus of the elastomer.

Analysis of Creep Curve

In Fig. 5 the rate of creep of the elastomer under a static load is compared with the creep rate under a comparable dynamic load. For viscoelastically simple materials, it is expected that the creep under a static load should be greater than that of similar loading that is applied cyclically. However, the results indicate the rate of creep is significantly faster when a 4 kg rms dynamic load is superimposed on top of an 8 kg static load compared with the nearly static loading conditions. Because of the instrument’s design, it was necessary to apply both a small dynamic load (0.5 kg rms) and frequency (0.5 Hz) along with a 13.1 kg static load. The static load was chosen because it is approximately the peak load that the sample sees during cyclic testing. This observation, while being contrary to predictions of linear viscoelasticity theory, is consistent with the observations of Derham and Thomas (8) and McKenna and Zapas (91, who have studied other filled systems.

The results indicate that the rate of creep is greatly enhanced by the application of a sinusoidally varying

1.2-1

0 20 40 6 0 80

Time (hours)

Fig. 5. Comparison of the creep rate for a PDMS elastomer subjected to A) static load of 13.1 kg and 0.5 kg nns dynamic load and B) a dynamic load of 4 kg rms along with an 8 kg static load.



dynamic load. Under nearly static conditions the weight loss of the elastomer is only about 0.9%, compared with nearly 10.1% after 60 h under a dynamic load (Table 1). The enhanced rate of creep and higher weight loss of the elastomer may be the result of polymer motion around the filler particles. Since the fillers are so much more rigid than the polymer, they essentially remain undefonned during the cyclic loading. This leads to motion of the polymer chains and friction in regions of very high local stresses. There- fore, it is more likely that the silicone-oxygen bond would rupture. As described earlier, once the ionic fragments are formed they are highly reactive, and these reactive ions can react with other chains in the network to relieve the mechanical stress (creep) or can undergo a cyclization reaction and subsequently be volatilized.

In 0. 6, fractional length is plotted as a function of logarithmic time. The creep curves for the filled siloxane elastomers exhibit a maximum slope defined by the peak in the derivative curve ( dh/d log t ) (where h is the height of the sample). Thus, there appears to be an acceleration of the chain-breaking process, which later slows down. The reduced rate of creep, at long times, may be the result of how the experiment was run. It should be recalled that in these experiments both the dynamic and static force are held constant. Therefore, as the length of the sample de-

Table 1. The Effect of Stress Level on Weight Loss.

Static Force Dynamic Force Percent Weight (kg) (kg rms) Loss/GO h

- 2 1 - 4 2 -6 3 - 8 4 - 10 5 -13 0.5

2.0 6.0 9.0

10.1 15.6 0.9

10

5

0

1 .o

0.8

0.6

0.4

0.2

0.0 1 10 100

Time (hours) Fg. 6. Creep behavior of the PDMS elastomer under a dynamic load: A) fractional length and B) the derivative of the creep curve us. time.

creases and the cross-sectional area increases, both the dynamic strain and the stress on the elastomer decrease. Lower stresses, as described below in The Influence of Force, can result in lower rates of creep. From a chemorheology point of view, initially the high rate of creep is probably occurring as a result of both the stress induced cleavage and the moisture catalyzed hydrolysis of the silicone-oxygen bonds. As described earlier, the presence of moisture can lead to hydrolysis of the polymer chains, causing cuts between the crosslinks and formation of silanol chain ends. This results in elastically ineffective crosslinks and a lowering of the effective crosslink density, and therefore, a more rapid rate of creep.

The Influence of Temperature

The dynamic creep results for a series of samples subjected to cyclic stress at various temperatures is given in Fig. 7. The results show that temperature influences both the rate of creep and the weight loss of the elastomer. For example, the weight loss of the elastomer at 166°C is only 1.796, compared with nearly 10.1% at 218°C after 60 h. In addition, the higher the temperature, the further accelerated the bond breaking and the rate of creep. From an Arrhe- nius analysis of the data (time it takes the for the sample to creep 25%), shown in Fig. 8, an activation energy of 1 1.1 kcal/mol is calculated. Literature values for activation energies for uncatalyzed stress relaxation (29) and uncatalyzed hydrolysis (1 5) in PDMS are 24 kcal/mol. The calculated value of the activation energy is most similar to the value obtained by Stein and Prutzman (10.3 kcal/mol) (18) for the catalyzed stress relaxation of a PDMS elastomer. Poten- tially, the surface of the aluminum oxide and/or the residual dibutyltin diacetate catalyst could catalyze the degradation of the filled elastomer.

8 8p fit3 0.9 - 0.8-

0.7-

0.6-

0.5-

5 a

m 0

- B .-

; 0.4-


John J. Fitzgerald, Arthur C. MarteUock, Paul L. Nielsen, and Robynn V. Schillace

The Influence of Force

In Fig. 9 the experimental results show that the rate at which the PDMS network bonds break is a function of the force level. In this experiment both the dynamic and static force were varied between 2 kg static/l kg rms dynamic and 10 kg static/5 kg dynamic rms. As shown in the Figure, as the force increases there is a corresponding increase in the rate of creep. Moreover, as shown in Table 1, as the force increases, there is also an increase in the amount of polymer volatilized during a 24 h period. The creep curves can be brought closer together by plotting them vs. absorbed power (Fig. 10). By plot-

100

n

L in

0

Q)

a 5 E i=

10 I I I 0.0020 0.0021 0.0022 0.0023

Temperature -’ ( K) Fig. 8. Arrhenius analysis of the data in Fg. 7 after the sample had crept 25%.

1 .o

0.8

0.6

0.4

0.2 -I I 1 10 100

Time (hours) Fg. 9. Fractional length us. time at a series of applied stress levels. A) - 2 kg static/ 1 kg dynamic rms; B) - 4 kg static/ 2 kg dynamic rms; C) - 6 kg static/ 3 kg dynamic rms; D) - 8 kg static/4 kg dynamic rms: and E) - 10 kg static/5 kg dynamic rms.

1’21

0 10 20

Power Absorbed (MJ) Fig. 10. Fractional length us. power absorbed at various loads. A) - 2 kg static/ 1 kg dynamic rms: B) - 4 kg static/ 2 kg dynamic rms; C) - 6 kg static/ 3 kg dynamic rms: D) - 8 kg static/4 kg dynamic rms; and E) - 10 kg static/5 kg dynamic rms.

ting the data in this manner, it suggests that the mechanism by which the samples creep is similar at all the forces studied, and it suggests that it is the work absorbed by the sample that influences bond breaking.

CONCLUSIONS

This paper has demonstrated that cyclic stress, in combination with high temperature, accelerates the degradation of a filled poly(dimethylsi1oxane) elas tomer, causing molecular changes that result in changes in the effective crosslink density, creep, and the storage modulus of the elastomer. The polymer weight loss was found to be linear with time and strongly dependent on the level of dynamic and static force applied to the sample. It is believed that because of the applied cyclic stress, breaks in the load- bearing chains in the elastomer network occur. These breaks act to relieve the mechanical stress on the chains, and also create highly reactive ionic fragments. The ionic fragments can react with other chains in the network or with water or undergo a cyclization reaction that results in the loss of polymer.

Contrary to the predictions of linear viscoelastic theory, the rate of creep is enhanced when the sample is subject to a sinusoidally varying load as compared with a comparable static load. In future reports, we intend to investigate whether the failure behavior of these elastomers can be explained in terms of either a cumulative dzmage approach or a cyclic dependent fatigue approach wherein the number of cycles to failure is a constant independent of frequency (9).



Reaction with other chains in the network is b e lieved to cause the observed creep behavior, and the activation energy for this process is similar to the reported values for the catalyzed degradation of PDMS. The observed changes in the storage modulus are complicated. It is believed that the storage modulus increases primarily because of the loss of polymer. Initially a decrease in the storage modulus is observed and is believed to be caused by the reaction of the ionic fragments with adventitious or added water resulting in the formation of silanol chain ends. This results in a decrease in the effective crosslink density, which reaches a minimum after about 20 h. In addition, there is a concurrent loss of polymer through volatilization that results in a decrease in the spacing between filler particles, and thus increases the viscous retardation, and at longer times results in an increase in the storage modulus and the effective crosslink density of the filled elastomer.

ACKNOWLEDGMENTS

The authors would sincerely like to thank Drs. W. Staudenmayer, J. Pavlisko, and L. R. Whitlock for many helpful discussions. Thanks are also due to R. Charlebois for materials preparation.

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The effect of cyclic stress on the physical properties of a poly(dimethylsiloxane) elastomer

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