Thermal cross-link behavior of an amorphous poly (para-arylene sulfide sulfone amide)

Cross-link behavior of an amorphous poly (para-arylene sulfide sulfone amide) synthesized via low temperature solution polycondensation was observed for the first time, when the polymer was subject to a series of thermal curing at 260 °C in air condition. The formation of cross-link network was demonstrated by the DSC and TGA results that T_g of the polymer enhanced from 259.17 °C to 268.89 °C, and the 1% weight loss temperature increased remarkably from 243.75 °C to 345.87 °C. EPR analysis further suggested that two kinds of free radicals, CO and C, induced by thermal curing were responsible for this cross-link behavior. According to FT-IR spectrum, the origin of these free radicals was confirmed as amide CO group in the polymer backbone. The cross-linking type was attributed to conventional radical cross-link reaction and the cross-link mechanism was discussed in detail subsequently.     

Power law of molecular weight dependence of lateral growth rate of isotactic polypropylene

Molecular weight (M) dependence of the lateral growth rate (V) of α form crystal of isotactic polypropylene (iPP) was studied. Reliable equilibrium melting temperature determined in our previous study was used for the analysis of supercooling dependence of V. A power law of M of V, , was obtained, where H is a small constant (H = 0.7). The small H, which is similar to that of the hexagonal phase of polyethylene (H = 0.7) in comparison with the value of H = 1.7 for the orthorhombic phase of polyethylene, confirmed our prediction of smaller H for “rod like” chain polymers because of easier chain sliding within the interface between the crystalline phase and the melt. Thus, the universality of the important role of topological nature in polymer crystallization was confirmed. Lateral surface free energy (σ) of the α form of iPP was obtained as σ ≅ 1.59 × 10⁻⁶ J/cm².     

Polymorphism and β-form structure of poly(octamethylene 2,6-naphthalate)

It was revealed that poly(octamethylene 2,6-naphthalate) (PON) existed in two different crystal structures, α- and β-form, depending on crystallization process: The α-form crystal was dominantly developed from the cold-crystallization, whereas the β-form was from the melt-crystallization. The apparent melting temperatures of α- and β-form crystals were characterized to be 175 and 183 °C, respectively. On the basis of X-ray diffraction and molecular modeling studies, the crystal structure of β-form, developed dominantly from the melt-crystallization, was identified to be triclinic with dimensions of a = 0.601 nm, b = 1.069 nm, c = 2.068 nm, α = 155.68°, β = 123.25°, γ = 52.85°, and with the space group of . The calculated crystal density was 1.243 g/cm³, supporting that one repeating unit of PON exists in a unit cell. The octamethylene units in the PON backbone take largely all-trans conformation in the β-form unit cell.     

Comparison of poly(ethylene oxide) crystal orientations and crystallization behaviors in nano-confined cylinders constructed by a poly(ethylene oxide)-b-polystyrene diblock copolymer and a blend of poly(ethylene oxide)-b-polystyrene and polystyrene

A poly(ethylene oxide)-b-polystyrene (PEO-b-PS) diblock copolymer with a number average molecular weight of PEO blocks, =8.8 kg/mol, and a number average molecular weight of PS blocks, =24.5 kg/mol, (volume fraction of the PEO blocks, f_PEO, was 0.26) exhibited a hexagonal cylinder (HC) phase structure. Small angle X-ray scattering results showed that the PEO cylinder diameter was 13.3 nm, and the hexagonal lattice was a=25.1 nm. The cylinder diameter of this HC phase structure was virtually the same as that in the blend system constructed by a PEO-b-PS diblock copolymer (=8.7 kg/mol and =9.2 kg/mol) and a PS homo-polymer (=4.6 kg/mol) in which the f_PEO was 0.32. The cylinder diameter in this blend sample was 13.7 nm and the hexagonal lattice was a=23.1 nm. Comparing crystal orientation and crystallization behaviors of this PEO-b-PS copolymer with the blend, it was found that the crystal orientation change with respect to crystallization temperature was almost identical. This is attributed to the fact that in both cases the PEO block tethering densities and confinement sizes are very similar. This indicates that when the of PS homo-polymer is lower than the PS blocks, the PS homo-polymer is located inside of the PS matrix rather than at the interface between the PEO and PS in the HC phase structure. On the other hand, a substantial difference of crystallization behaviors was observed between these two samples. The PEO-b-PS copolymer exhibited much more retarded crystallization kinetics than that of the blend. Based on the small angle X-ray scattering results, it was found that in the blend sample, the HC phase structure was not as regularly ordered as that in the PEO-b-PS copolymer, and thus, the HC phase structure contained more defects in the blend. This led to a suggestion that the primary nucleation process in the confined crystallization is a defect-controlled process. The mean crystallite sizes were estimated by the Scherer equation, and the PEO crystal sizes are on the scale of the confined size.     

Copolymer crystallization: Approaching equilibrium

Melt crystallization of random copolymers leads to solids with crystalline fraction w_c and final melting temperature that are substantially below the predictions of Flory's equilibrium crystallization theory. Model ethylene/butene random copolymers, when crystallized as multilayer films by rapid solvent evaporation, exhibit increased w_c (50% relative) and (4 K) compared to melt crystallized values. For a copolymer with 0.92 mol fraction ethylene, the density-derived crystallinity w_c=0.6 is the same as that from Flory's theory, although the maximum observable crystal thickness from remains about 25% of the theory value. These effects are seen because crystallization from solution occurs without many of the constraints to segment dynamics that limit crystalline fraction during melt crystallization. Crystal thickness is dominated by secondary nucleation barriers in both melt and solution. Chain or sequence folding is much more regular in the solution crystallized material, and amorphous layer thickness is reduced from about 8 nm to 3 nm.     

Relative stabilities of Ce(IV) and Ce(III) limiting carbonate complexes at 5-50 °C in Na aqueous solutions, an electrochemical study

The normal potential of the Ce(IV)/Ce(III) redox couple was determined by square wave voltammetry (SWV) at different temperatures in solutions with a constant ratio [CO₃²⁻]/[HCO₃⁻] ≈10 for high ionic strengths (3.29 mol dm⁻³ at 4.39 mol dm⁻³): varies from 259.5 to 198.0 mV/S.H.E. in the 15-50 °C range. Linear variations were found for versus (RT/F)ln(mCO₃²⁻), leading to the stoichiometry, Ce(CO₃)₆⁸⁻ for the Ce(IV) limiting complex. But the slopes of these linear variations were actually found in the range 1.8-1.9, not exactly 2. This was interpreted as dissociation of the Ce(IV) limiting complex following the reaction: Ce(CO₃)₅⁶⁻ + CO₃²⁻ → Ce(CO₃)₆⁸⁻ and as dissociation of the Ce(III) limiting complex following the reaction: Ce(CO₃)₃³⁻ + CO₃²⁻ → Ce(CO₃)₄⁵⁻; for which maximum possible values of log₁₀ K_IV,6 and log₁₀ K_III,4 were estimated via fitting in the 15-50 °C temperature range (log₁₀ K_IV,6 = 0.42 (0.97) and log₁₀ K_III,4 = 0.88 (7.00) at 15 °C (50 °C). The normal potential was found to decrease linearly with T, these variations correspond to , with T⁰ = 298.15 K and . The apparent diffusion coefficient of Ce(IV) was determined by direct current polarography (DCP), cyclic voltammetry (CV) and square wave voltammetry. It was found to depend on the ionic strength and to be proportional to T.