The effects of DV on the morphology and rheology of TPVs and their nanocomposites.

Dynamic vulcanization (DV) is a process that involves the curing of a rubber while it is being mixed with a plastic to form a homogenous melt. This process was discovered by Gessler and Haslett (ref. 1) and was further applied to PP/ EPDM blends by Fisher (ref. 2). The work by Fisher, however, involved only the partial curing of the EPDM rubber using a free radical mechanism. Coran and Patel (refs. 3 and 4) developed the first commercial TPVs in which the rubber phase was fully crosslinked.

Since then, a significant amount of work has been done on the rheological properties of dynamically crosslinked PP/ EPDM blends. Ha et al (ref. 5) carried out studies on the viscosity of PP/EPDM blends that were dynamically cured by dicumylperoxide. They reported that EPDM-rich blends exhibit high viscosity in the low shear rate region. Goettler et al (ref. 6) and Han and White (ref. 7) have also carried out studies on the theological properties of PP/EPDM TPVs.

The dual phase morphology of the TPV is a very important parameter in determining its rheological properties. The curing agent content, amount of extender oil, speed and temperature of mixing and type of rubber all affect the morphology. Medintseva et al (ref. 8) studied the effect of oil and curing agent content on the properties of TPVs. In recent times, atomic force microscopy (AFM) has developed as a useful tool for cure state characterization and phase morphology of TPVs (ref. 9). It is our purpose in this article to study intentional changes to the dynamic vulcanization parameters and their resulting effect on the morphology and rheology of the TPVs.

Experimental

Material

V-5600 and V-2504 EPDM rubbers were obtained from ExxonMobil Chemical. The V-5600 has 68.5 wt. % ethylene, ENB (ethylidene norbornene) content of 5 wt. % and a Mooney viscosity of 72, whereas for V-2504 the ethylene is 57.5 wt. %, ENB content is 4.7 wt. % and Mooney viscosity is 25. Thus, the major difference between these two rubber components lies in their Mooney viscosity, which is reported as ML (1+4) at 125[degrees]C. They henceforth will be referred to as the HM (high Mooney) rubber and the LM (low Mooney) rubber, respectively. P-6523 polypropylene was obtained from Basell Polyolefins, and the extender oil was S-150 from Sunoco. The curing system consisted of SP-1045 octylphenol resin from Schenectady International activated with stannous chloride. The silicate nanoclay used was C-20A from Southern Clay Products. This clay is a natural montmorillonite modified with dimethyl dihydrogenated tallow quaternary ammonium salt at 90 meq/100 g clay, which renders the clay more hydrophobic to aid intercalation of the non-polar PP/EPDM system.

Preparation of TPVs and their nanocomposites

The TPVs and their nanocomposites were prepared on a laboratory scale using an internal mixer. Mixing was accomplished in two stages. First, the EPDM rubber was compounded with extender oil and cure activator at 80[degrees]C and 80 rpm using internal mixer type rotors. In the case of TPV nanocomposites, the silicate nanoclay was also added to the rubber phase in this first stage of mixing, along with extender oil and cure activator under the same mixing conditions. Once the rubber was compounded, it was then mixed in a second stage with PP at 180[degrees]C and 80 rpm until a homogenous melt was obtained. At this point, the curing agent was added and mixing was continued for another seven minutes. A control sample to be used as a point of comparison is described in table 1, along with intentional variations made to the composition and processing parameters to cause variations in morphology and rheology of the TPVs.

In the nanocomposites, the clay loading was kept at 6 wt. % with respect to the combination of HM type EPDM rubber and process oil. The HM rubber was used to make the nanocomposites. Changes were made to the cure system concentration to counteract clay effects on the cure characteristics of the EPDM rubber (ref. 10). Thus, 10 phr of curing agent and 2 phr of stannous chloride (with respect to EPDM rubber content) were used, and the EPDM in this case was 130 phr oil extended.

Rheological characterization

The rheological characterization of the TPVs and their nanocomposites was carried out at 210[degrees]C using an advanced polymer analyzer. No special sample preparation was necessary. A strain sweep was performed at a constant frequency of 6.28 rad/s over a range from 1.0% to 400% strain and a frequency sweep over a range from 0.2 to 209 rad/s at a constant strain of 7%. The rotors used were biconical with an angle of 7[degrees] and a diameter of 63.5 ram.

Phase morphology characterization

The phase morphology was characterized by atomic force microscopy (AFM) in non-contact mode. The sample surface was prepared for examination by cryogenic microtoming at -80[degrees]C using a cryo ultramicrotome.

Results

Rheology

Viscosity measurements plotted in figure 1 ,ate constant in the low strain region, indicating linear viscoelastic behavior. Subsequent frequency sweep tests were thus carried out in this region at a constant strain of 7%. At high values of strain, there is a drop in complex viscosity for all TPV batches, indicating a breakdown of the molecular and mesoscale network structure comprising a TPV. The viscosities of samples TPV-C and TPV-N, representing a higher cure state in the rubber phase, are greater than that of the control TPV. On the other hand, TPV-T and TPV-R exhibit comparatively lower melt viscosities with respect to the control TPV.

[FIGURE 1 OMITTED]

Figure 2 shows the dependence of the complex dynamic viscosity on frequency at 210[degrees]C for TPVs and a nanocomposite. All batches of TPVs exhibit Power Law type behavior, with no indication of approaching a plateau at low shear rate. The convergence of the curves at high shear rate indicates a lower Power Law index and hence greater network formation in those TPVs showing higher viscosity. TPV-N exhibits a higher viscosity as compared to the control TPV due to a combination of the silicate nanoclay reinforcement of the rubber phase, but it is lower than that of TPV-C, despite having the same higher curing agent concentrations, because of the deleterious effect of the silicate nanoclay on both the cure rate and cure state of the EPDM rubber.

[FIGURE 2 OMITTED]

At low frequencies, the viscosity of TFV-O is similar to that of the control TPV. At higher frequencies however, TPV-O and TPV-S show greater shear thinning behavior than the control TPV, even though they are not of higher viscosity. As a result, there is a crossover of the viscosity curves as shear rate increases. The lowest viscosities are exhibited by TPV-R and TPV-T, respectively. In the case of TPV-R, the rubber used is the LM rubber instead of the HM rubber, which is used in making the control TPV. The LM rubber has a Mooney viscosity of 25, compared to 72 in the case of the HM rubber.

The values of the storage modulus shown in figure 3 at 210[degrees]C follow a trend similar to that observed for viscosity in figure 2. This behavior is not surprising, since G' is the larger component of the dynamic viscosity. TPV-C has a higher value of modulus as compared to the control TPV, and TPV-T has the lowest value. The storage modulus of TPV-O is slightly higher than that of the control TPV in the low frequency region. With an increase in the frequency, there is a crossover in the modulus values of the melts, and in the high frequency region the control TPV has a higher storage modulus.

[FIGURE 3 OMITTED]

Figure 4 shows the dependence of tan [delta] on frequency for selected TPV samples. The lower tan [delta] values for TPV-C and TPV-N indicate less deformable and more elastic TPV melts than the control TPV, due to the increased cure system concentration in the ease of TPV-C and nanoclay reinforcement of the rubber phase in the case of TPV-N. When the LM rubber is used instead of the HM rubber, as in the case of TPV-R, there is an increase in the value of tan [delta], indicating a less elastic material.

[FIGURE 4 OMITTED]

Morphology

In the AFM photomicrographs of various TPV samples shown in figures 5-9, the rubber phase is white in color and the poly-propylene phase is dark. This higher cure system concentration of sample TPV-C generates a smaller dispersed rubber phase with respect to the control TPV, which can be observed by comparing figures 5 and 6.

[FIGURES 5-9 OMITTED]

A lower speed of mixing, i.e., 50 rpm, apparently does not generate sufficient shear forces to enable break up of the rubber particles. The resulting TPV-S has a larger size of rubber particles as compared to the control TPV, observed by comparing AFM photomicrographs in figures 5 and 7. These large rubber particles can coalesce to form agglomerates of rubber, forming a semi-continuous phase morphology.

TPV-R, made using the LM robber, has a melt viscosity that is lower than that of the control TPV and is observed in the AFM images of figure 8 to achieve a nearly co-continuous phase morphology. At the least, by comparing figures 5 and 8, it is observed that the size of the rubber particles (white) is larger in TPV-R than in the control TPV.

Figure 9 is an AFM photomicrograph of TPV-T made at 210[degrees]C. The rubber particles are larger in size as compared to those in the control TPV (figure 5). These large rubber particles again tend to coalesce to form the agglomerated type of morphology. TPV-T also exhibits the lowest melt viscosity of all TPVs made, as shown in figure 2.

Discussion

TPVs and their nanocomposites exhibit Power Law type behavior without indication of any Newtonian viscosity plateau. Changes in the rheological behavior of the TPV melt are observed when variations are made to the dynamic vulcanization parameters, namely change in cure system concentration, speed and temperature of mixing, type of rubber used and amount of oil added to the rubber. In addition to these, the rheological properties of the TPV melt can also be modified through nanoclay reinforcement of the lubber phase.

These changes alter the viscosity, as well as the elastic storage modulus, and are reflected in the tan delta values that represent a ratio of the viscous/elastic flow behavior of the TPV melts. In general, making the rubber phase more ductile results in a softer TPV with a larger and more elongated rubber phase that is manifested in lower viscosity and modulus, but a higher tan delta, i.e., a more dissipative melt. The opposite occurs when the rubber phase is made more deformation-resistant through reinforcement at either the molecular scale (cross-linking) or mesoscale (nanoplatelets). A larger rubber phase size, with the same concomitant effects on melt theology, occurs when the rate of crosslinking is lower relative to the rate of mixing, or if the final crosslink density is reduced. These changes can be effected either chemically (crosslinking reactant concentration) or physically (temperature or mixing intensity). Specific comments on the various actions taken to implement these changes are noted as follows:

An increase in cure system concentration generates a smaller dispersed rubber phase size, consequently exhibiting change in the rheological properties of the TPV melt. Reinforcement of the rubber phase with nanoclay also generates higher melt viscosity and storage modulus, along with lower tan [delta] in comparison to a control TPV.

At a lower speed of mixing, the rate at which the rubber particles are dispersed in the continuous PP matrix would be reduced relative to the rate at which the crosslinking of the rubber proceeds. In addition, lower shear forces are generated during mixing at lower speed. TPV-S, mixed at 50 rpm, exhibits a sinuous type of agglomerated rubber phase morphology that could be a reason for its comparatively low melt viscosity and storage modulus.

A lower molecular weight rubber should produce a looser molecular network at the same level of phenolic curative due to the greater concentration of molecular ends. According to Abdou-Sabet and Patel (ref. 11), weaker TPV compositions are in fact obtained from rubbers of low molecular weight. In this study, the LM rubber used to make TPV-R has a Mooney viscosity of only 25 prior to 100 phr oil extension. (This oil extension could further reduce the shearing forces generated during mixing.) A combination of these effects would lead to a larger rubber phase size, as observed in figure 8. Melt viscosity and storage modulus G' of TPV-R are also lower, and tan [delta] values are higher than the corresponding values for a control TPV.

[FIGURE 8 OMITTED]

Increasing the mixing temperature to 210[degrees]C decreases the viscosity of the homogenized rubber-plastic melt prior to the crosslinking stage. The resulting decreased viscosity again reduces the shearing forces that are necessary to enable the break-up of the rubber particles. At the same time, the vulcanization "kinetics are enhanced by the higher temperature, reducing the time period over which the rubber particle size can be refined. TPV-T made under these conditions thus exhibits a large phase size and the lowest values of melt viscosity and storage modulus.

The rubber phase in TPV-O has a higher oil extension of 130 phr. At low frequency, the melt viscosity and storage modulus values of TPV-O actually exceed those of the control TPV, although with increasing oil extension a drop in the above properties would be expected. The anomalous increase in these properties suggests that at low frequencies the mesoscale network overshadows molecular lubrication. With an increase in frequency, a crossover in values occurs wherein the melt viscosity and storage modulus values drop below those of the control TPV, suggesting breakdown of the network that allows lubrication to become the predominant effect. There is again the possibility of migration of the oil from the rubber phase to the PP phase, which would have a plasticizing effect on the PP melt. Unfortunately, micrographs of the rubber phase structure are not available.

A general relationship between the phase structure and rheology of TPVs emerges from these results. Namely, both the viscous response (dynamic viscosity) and the elastic response (storage modulus) to rate of strain increase with decreasing rubber phase size as determined qualitatively from the micrographs presented. The order of the dispersed rubber phase size is as follows: TPV-C > TPV-control > TPV-S > TPV-R > TPV-T. The interpretation of this effect is that smaller dispersed phase size generates a larger interfacial area that augments the viscosity, while the larger number of smaller rubber particles results in a higher particulate concentration with smaller interparticle separations. We speculate on the existence of a percolation threshold, above which an interparticle network facilitates elastic stress transfer. The lower tan delta for melts comprising a smaller rubber phase size further supports this model. The gradual convergence of the curves for both viscosity and storage modulus at increasing strain rate suggests the shifting equilibrium in rebuilding networks at both the molecular and mesoscale, respectively, as the rate of disruption increases. Similar effects accrue from increasing the strain level.

It is expected that these rheological changes will affect the processability of the TPVs in typical fabrication operations, as well as modify the mechanical properties of the formed articles. These issues are the subjects of further studies.

Conclusions

Controlled variations made to the composition and mixing of TPVs formed by dynamic vulcanization of EPDM rubber in polypropylene suggest an inverse relationship between the crosslinked rubber phase size and both the viscous and elastic rheological parameters. In addition, the melt becomes less "lossy" (lower tan delta) as the phase size is reduced. Atomic force microscopy indicates the coalescence of the less refined rubber particles into agglomerates, which can extend continuously through a substantial portion of the TPV melt. Various routes exist for generating these changes to the rubber phase, including the temperature and speed of mixing, the type of rubber used and the cure system concentration. Silicate nanoclay reinforcement of the rubber phase also increases the elasticity of the TPV, raising the complex melt viscosity and storage modulus of the nanocomposite, while reducing its tan 8 through a filler effect in the rubber phase. However, the presence of the smectic clay diminishes the cure of the EPDM rubber, requiring an increase in the cure system concentration to maintain crosslink density.

Table 1--variations to dynamic vulcanization parameters

Name                   EPDM,      PP,       Resin,     Stannous
                         %         %         phr     chloride, phr

Control                 70         30          4           1
TPV-T                   70         30          4           1
TPV-C                   70         30         10           2
TPV-S                   70         30          4           1
TPV-R (LM rubber)       70         30          4           1
TPV-O                   70         30          4           1
TPV-N (6 wt. %          70         30         10           2
 silicate nanoclay)

Name                    Oil,     Temp.,     Speed,
                        phr    [degrees]C    rpm

Control                 100       180         80
TPV-T                   100       210         80
TPV-C                   100       180         80
TPV-S                   100       180         50
TPV-R (LM rubber)       100       180         80
TPV-0                   130       180         80
TPV-N (6 wt. %          130       180         80
 silicate nanoclay)

References

(1.) A.M. Gessler and W.H. Haslett, U.S. patent no. 3,037,954 (1962).

(2.) W.K. Fisher, U.S. patent no. 3,806,558 (1974).

(3.) A.Y. Coran and R. Patel, U.S. patent 4,104,210 (1978).

(4.) A.Y Coran and R. Patel, Rubber Chem. Technol. 53, 141, (1980).

(5.) Ch. S. Ha, D.J. Ihm and S. Ctr. Kim, J. Appl. Polym. Sci. 32, 6,281, (1986).

(6.) L.A. Goettler, J.R. Richwine and F.J. Wille, Rubber Chem. Technol. 55, 1,448, (1982).

(7.) P.K. Han and J.L. White, Rubber Chem. Technol. 68, 728, (1995).

(8.) Tatiana Medintseva, Natalia Erina and Eduard Prut, Macromol, Symp. 176, 49, (2001).

(9.) Oansuk Chung and Hari P. Nadella, Annual Technical Conference Society of Plastics Engineers (2001), 59th (Vol. 3), 2,926-2,930.

(10.) Hemant Thakkar and Lloyd A. Goettler, Annual Technical Conference, Society of Plastic Engineers (2003), 61st, Nash-ville, TN, May 2003.

(11.) S. Abdou-Sabet and R. Patel, Rubber Chem. Technol. 64, 769, (1991)