Second network modulus

Figure 1. Effective first network modulus, Gle, after complete removal of first network cross-links plotted against second network modulus, Gi. Calculated from the composite network theory of Flory (19J for G, — 0.75 MPa.

The result is important for the discussion in Part 3. Multiplication of the v-values by RT gives the corresponding moduli. The effective modulus of the first network after removal of first network crosslinks, Gie, has been calculated for a first network modulus, G-j, of 0.75 MPa. In Figure 1, G. e is plotted against the modulus of the second network before removal of the first network cross-links, G2. It can be seen that the memory effect increases with increasing modulus or degree of cross-linking of the second network. Gx and G2 max are related to the experiment to be discussed in Part 3. 

Figure 2. The principle of the two-network method for cross-linking in a state of simple extension. First network with modulus Gy is entirely due to chain entangling. Second network with modulus Gx is formed by cross-linking in the strained state. Both Gy and Gx can be calculated from the two-network theory.

Sequential IPN. The preceding analysis does not apply to the case of a sequential IPN. The formation of this system originates with the synthesis of the network (1). Then, network (1) is swollen with monomer (2) which is subsequently polymerized in situ to form a second network. Due to perturbed chain dimensions, the modulus of the first network is higher than the corresponding modulus in the unswollen state by a factor equal to v [ ] ... 

However, the validity of Eq. (21) for the systems montioned in chapter 3 cannot be proved. As shown in Fig. 17, hydrazide-cured resins having a maximum in 82 also exhibit a maximum E. Besides, the solubility parameter (82) computed from the values reported by Hoy 55) is 9.9 for the anhydride-cured resins and 10.6 for the DETA-cured resins, while E of the anhydride-cured resins is higher than that of the DETA-cured resin. These results have been derived from the ten-second shear modulus [3G(10)] and flexural modulus (Table 7) 40,46). in the resins obtained according to Eqs. (4) to (7) (e.g., anhydride-cured resins), the network formation is not perfect and there exist many dangling chains. 

The study of viscoelastic properties of sequential IPNs based on PU and styrene-DVB copolymer [192] has shown that increasing the content of the second network affects nonmonotonously the position and the shape of the curves of temperature dependence of the elastic characteristics. This dependence is shown in Figs. 36 and 37. The modulus diminishes only in a narrow... 

Fig. 36 Dependence of the shear modulus on the volume fraction of the second network at various temperatures (1) 243, (2) 253, (3) 293 K [192]...

The major limitation of rubber toughening of thermosets results from the fact that the increase in toughness can be achieved only at the expense of high-temperature performance or of mechanical properties, e.g., a decrease in modulus and yield stress. This can be unacceptable for structural and long-term applications (see Fig. 13.7). A second limitation is the lack of significant success in the toughening of high-Tg networks (see Fig. 13.8). 

The rate processes of diffusion and catalytic reaction in simple square stochastic pore networks have also been subject to analysis. The usual second-order diffusion and reaction equation within individual pore segments (as in Fig. 2) is combined with a balance for each node in the network, to yield a square matrix of individual node concentrations. Inversion of this 2A matrix gives (subject to the limitation of equimolar counterdiffusion) the concentration profiles throughout the entire network [14]. Figure 8 shows an illustrative result for a 20 X 20 network at an intermediate value of the Thiele modulus. The same approach has been applied to diffusion (without reaction) in a Wicke-Kallenbach configuration. As a result of large and small pores being randomly juxtaposed inside a network, there is a 2-D distribution of the frequency of pore fluxes with pore diameter. 

Viscoelastic measurements of ionomers have been used to indirectly characterize the microstructure and to establish property structure relationships. Forming an ionomer results in three important changes in the viscoelastic properties of a polymer. First, T usually increases with increasing ionization. This is a conseqi nce of the reduced mobility of the polymer backbone as a result of the formation of physical, ionic crosslinks. Second, an extended rubber plateau is observed in the modulus above T, again as a result of the ionic network. Third, a high temperaturi mechanical loss is observed above T, which is due to motion in the ion-rich phase. The dynamic mechan cal curves for SPS ionomers shown in Fig. 9 clearly demonstrate these three characteristics. 

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