Rheological and dielectric behavior was examined for concentrated solutions of a styreneisoprene-styrene (SIS) triblock copolymer in monomeric and polymeric I-selective solvents, n-tetradecane (C14) and a low-M homopolyisoprene (I-1; M ) 1.4K). The I blocks had symmetrically once-inverted dipoles along the block contour, and their midpoint motion was dielectrically detected. The SIS solutions exhibited rubbery, plastic, and viscous behavior at low, intermediate, and high temperatures (T). Dielectric and viscoelastic data strongly suggested that the S and I blocks were more or less homogeneously mixed in the viscous regime. In the rubbery and plastic regimes, the S blocks were segregated to form spherical domains, and the I blocks took either the loop or bridge conformation. In these regimes, the inverted dipoles of the I blocks enabled us to dielectrically estimate the loop fraction, φ1 = 60% in C14 and I-1. These loops, having osmotically constrained conformations, strongly affected the rheological properties of the SIS solutions. A strong osmotic constraint in C14 resulted in almost equal contributions of the loops and bridges to the equilibrium modulus. The loop contribution became less significant in I-1 that (partly) screened this constraint. Similarly, the yield stress σy in C14 was essentially determined by dangling (noninterdigitated) loops at relatively high T where the S/I mixing barrier was rather small, while the bridges and interdigitated loops had a large contribution when this barrier was enhanced, i.e., at lower T and/or in I-1 (a poorer solvent for the S blocks than C14).

Linear viscoelastic and dielectric relaxation behavior was examined for blends of styrene−isoprene (SI) diblock copolymers in a nonentangling homopolyisoprene (hI) matrix. The copolymers formed spherical micelles with S cores and I corona. The I blocks had type-A dipoles, and their global motion induced dielectric relaxation. The micelles exhibited fast and slow relaxation processes for both dielectric loss ε‘‘ and viscoelastic moduli G*. For the fast process, a nearly universal relationship was found for reduced moduli G r* = [M bI/c bI RT]G*SI and reduced frequencies ωτ*, where G*SI was the contribution of the SI micelles to G*, M bI and c bI were the molecular weight and concentration for the I block, and τ* was a relaxation time that increased exponentially with c bI for large c bI. Similar behavior was found also for ε‘‘SI (SI contribution to ε‘‘). These features were qualitatively the same as those for the relaxation of star chains, indicating that the fast process corresponded to starlike relaxation of individual I blocks tethered on the S cores. Effects of the S cores on τ* were discussed within the content of the tube model. For the slow process of the concentrated micelles entangled through their I blocks, a relaxation time τs was close to the Stokes−Einstein (SE) diffusion time τSE evaluated from the viscosity associated with the fast process. Thus, the slow process of those concentrated micelles was attributed to their SE diffusion governed by the relaxation of individual corona I blocks. On the other hand, for dilute micelles in the nonentangling matrix, τs was significantly shorter than τSE but close to the SE diffusion time evaluated from the matrix viscosity. This result suggested that the slow process of those micelles corresponded to the SE diffusion in the pure matrix.

Nonlinear stress relaxation after imposition of step strain γ (≤2) was examined for blends of styrene−isoprene (SI) diblock copolymers in a homopolyisoprene (hI) matrix. The blends contained spherical micelles with S cores and I corona. For most cases, the blends had no plasticity and exhibited complete relaxation. Fast and slow relaxation processes characterizing the linear viscoelastic behavior of the micelles (part 1) were observed also for nonlinear relaxation moduli G(t,γ). For sufficiently small γ, G(t,γ) agreed with the linear relaxation moduli evaluated from the G* data of part 1. However, G(t,γ) decreased for larger γ (mostly for γ > 0.1). This nonlinear damping was much more significant for the slow process than for the fast process. For quantitative analysis of the damping behavior, the linear viscoelastic relaxation time τ* of the fast process was utilized to successfully separate the G(t,γ) data into contributions from the fast and slow processes, G f(t,γ) and G s(t,γ), in the following way:  At t > 6τ* where the fast process had negligible contribution to G(t,γ), G s(t,γ) were taken to be identical to G(t,γ). By extrapolating those G s(t,γ) data to shorter time scales, G s(t,γ) were evaluated at t < 6τ*. G f(t,γ) were evaluated as G(t,γ) − G s(t,γ). For both G f(t,γ) and G s(t,γ), the terminal relaxation times were insensitive to γ and the time−strain separability held in respective terminal regions. This separability enabled us to define damping functions in those regions, h x(γ) = G x(t,γ)/G x(t) (x = f, s). For the fast process of the SI micelles, h f(γ) exhibited only modest γ dependence that was in good agreement with the dependence for homopolymer chains. This result indicated that the fast process corresponded to relaxation of individual corona I blocks, giving a strong support for the discussion of part 1. On the other hand, h s(γ) of concentrated micelles exhibited very strong γ dependence that was comparable, in both magnitudes and sensitivities to the I block concentration and molecular weight, with the dependence of the damping function h C 14 (γ) obtained for solutions of SI in an I-selective solvent, n-tetradecane (C14). Those solutions exhibited plasticity due to macrolattices of the micelles, and their nonlinearity was attributed to strain-induced changes in the micelle position. Thus, the similarity of h s(γ) and h C 14 (γ) suggested that the slow process of the concentrated micelles in the SI/hI blends was related to the changes in the micelle position and the subsequent micelle diffusion, again supporting the discussion of part 1.

The complex shear modulus was measured for four low molecular‐weight polystyrenes (Mw = 10,500, 5970, 2630, and 1050) near and above the glass transition temperature. For the lowest molecular weight sample, the method of reduced variables, the time–temperature superposition principle, was applicable, while it was not applicable for the higher M samples. For these higher M samples, it was assumed that the complex modulus is composed of two components (R and G components). The R component was estimated by subtracting the G component, which was assumed to be the same as the modulus of the lowest molecular weight sample. The time–temperature superposition principle was applicable to each of the R and G components, and the shift factors were different from each other. The contribution of the R component to the total complex modulus decreased with decreasing M. Anomalous temperature dependence of the steady‐state compliance for low M polymers as Plazek reported could be attributed to difference in temperature dependence of the two components. The estimated complex modulus for the R component was in accord with that calculated by spring‐bead model theory. © 1999 John Wiley & Sons, Inc. J Polym Sci B: Polym Phys 37: 389–397, 1999

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