The molecular motions underlying the dielectric and dynamic-mechanical 8 relaxation in poly-(methyl methacrylate) (PMMA) have been elucidated in detail by means of two-dimensional (2D) and threedimensional (3D) 13C exchange NMR of the carboxyl moiety and 2D 2H exchange NMR of the methoxy group. The identity of the motions observed by NMR and the ^-relaxation dynamics is proved by the agreement of the measured correlation times. The selective-excitation "3D" NMR spectrum proves that, for every mobile side group, a relatively well-defined motion between two potential-energy minima occurs. The 2D spectral pattern shows that the OCO plane of the side group undergoes 180°(±20°) flips. Experiments with multiple exchange and selective saturation for analysis of the growth of exchange signals (MESSAGE) prove that the molecular motions responsible for the 8 relaxation are associated with a distribution of correlation times, which appears to be bimodal with both mobile and trapped side groups. Consistently, analysis of the integral 2D exchange intensity shows that around 330 K only about 50% of the side groups participate in the large-amplitude dynamical process on the time-scale of the ^-relaxation correlation time. The 2D 2H NMR spectra, while exhibiting narrowing due to methyl-group rotation around the 0-CH3 bond, exclude any significant motion of the methoxy group around the C-OCH3 bond. Both the 13C and the 2H 2D NMR spectra provide compelling evidence that the side-group flip is accompanied by a main-chain rearrangement which can be characterized as a random rotation around the local chain axis with a 20°r oot-mean-square amplitude. This is ascribed to the fact that the asymmetric side group, after the flip, does not fit into its original environment. These findings explain both the dielectric and the dynamic-mechanical 8 relaxations of PMMA.

The occurrence and rate of 180° chain-flip motions in the crystalline regions of two polyethylenes were studied by 13C NMR. In high-density polyethylene (HDPE) and in ultradrawn ultrahigh molecular weight polyethylene (UHMWPE) fibers, the changes in the 13C−13C dipolar couplings brought about by the reorientations of the 13C−13C internuclear vectors in the crystallites were observed. In the HDPE sample, which was labeled with 4% of 13C−13C pairs, the rotational motion was observed directly via two-dimensional exchange spectroscopy, stimulated-echo decays, and 1D line shape changes monitoring the 13C−13C dipolar coupling. The data show that the jumps occur between two sites, with a rotation angle of 180° and with a jump rate of ∼10/s at ambient temperature. The correlation function of the motion was found to be slightly nonexponential, with a stretched-exponential β parameter of 0.8 ± 0.1. The data yield an activation energy of 93 ± 10 kJ/mol for the 180° chain flips. In the fibers, the narrowing of natural-abundance 13C−13C dipolar satellites is a clear NMR signature of the chain motion, indicating a jump rate of 150/s at 360 K, which is 20 times slower than in the unoriented HDPE. The correlation time dependence of the 1H T 1 ρ relaxation time, which probes the modulation of H−H dipolar couplings in the crystallites, was determined directly. Relations between the chain flip motion, the dynamic-mechanical α-relaxation, creep, and drawability are discussed.

Stiff macromolecules with flexible side chains are investigated by proton spin diffusion experiments with 13C detection and by a recently developed wideline separation 2D lH-13C NMR experiment (WISE-NMR spectroscopy). The conformational order and the molecular mobility of the alkyl side chains (CieHsa) are characterized for samples with polyester, polyamide, and polyimide main chains. The side chains, which are phase-separated from the main chain in a layer-type structure, can form crystalline as well as amorphous phases. The sizes of these domains depend on the nature of the main chains and their organization.In the polyimide and the polyester with regular main-chain packing, crystalline as well as amorphous regions are observed extending over more than one layer spacing. The heterogeneity observed in the polyamide is only of the order of the layer spacing. The polyester can also be obtained in a modification with uniformly ordered but anisotropically mobile side chains and conformationally disordered main chains. These results indicate coupling between the main-chain and side-chain packing in the investigated stiff macromolecules with flexible side chains.

The coupling of the a and 8 processes in poly(ethyl methacrylate) has been investigated in detail by multidimensional 13C solid-state NMR of the carboxyl moiety. In the glassy state the underlying molecular motion is anisotropic and involves a it flip of the side group coupled to a rocking motion around the local chain axis with a ±20°a mplitude. Above the glass transition (Tg) the molecular motion remains highly anisotropic. The geometry of the molecular motion is similar to that in the glass; however, the rocking amplitude increases upon raising the temperature above Tt. This is indicative of a pronounced influence of the a main-chain motion on the 8 side-group motion which manifests itself by a marked increase of the rocking amplitude to a value of ±50°at 365 K (Tg+ 27 K). It eventually leads to a locally anisotropic uniaxial chain motion at 395 K (Tg + 57 K). This behavior differs significantly from that of other amorphous polymers above T( where the molecular motions of both the main chainandside groups are isotropic. The averaged correlation times extracted from NMR experiments are in good agreement with data from dielectric relaxation.

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Cellulose as well as two cellulose/poly(vinyl alcohol) blends with compositions 60/40 and 80/20 w/w exposed to water are investigated by 1H-, 2H-, and 13C solid-state NMR spectroscopy. For pure cellulose, the lower temperature, secondary dielectric relaxation process can be attributed to the onset of motion of adsorbed water molecules as revealed by 2H-NMR spectroscopy. This water is not crystalline below 270 K. Three distinct kinds of water bound to the polymer matrix are detected, as far as dynamic behavior is concerned. First there is nonfreezable, strongly bound water that is rigid but amorphous at low temperatures. The second component is highly mobile and exhibits isotropic motion even below 270 K. Interestingly, there is a third component of water molecules that undergo well-defined 180° flips around their bisector axis with a rate greater than 105 s-1 due to anisotropic constraints. In contrast to the first two kinds, this component cannot be removed from the polymer matrix by drying even at elevated temperatures and its motional process is observed over the whole temperature range, investigated from 190 to 370 K. All three kinds of matrix water coexist in a wide temperature range. In the blends, 2D 1H−13C heteronuclear wide line separation (WISE) NMR spectroscopy shows that at our low concentrations the water is predominantly associated with the cellulose backbone. No water can be detected in the immediate vicinity of the poly(vinyl alcohol). Applying spin diffusion, we detected nanoheterogeneities in the range of about 3 nm within these systems.

The main chain dynamics of amorphous poly(ethyl methacrylate) (PEMA) and poly(methyl methacrylate) (PMMA) below and above their respective glass transition temperatures T g are analyzed by two-dimensional solid-state exchange 2H NMR spectroscopy. In both polymers, a restricted mobility of the polymer backbone is already present in the glassy state, as is directly demonstrated and quantified using samples deuterated at the methyl and methylene moieties of the polymer main chain. The unusual main chain mobility below T g is coupled to the β-relaxation process, which involves 180° flips of the carboxyl side groups. At their respective glass transition temperatures, the coupling of the β-process to the main chain motions manifests itself differently in both polymers; the smaller ester side group reorients comparatively fast in PMMA, whereas in PEMA, the reorientation of the bulkier side group remains anisotropic and the correlation times are slower by about 1 order of magnitude. Therefore, in PMMA, the β-relaxation predominantly influences the time scale of the α-relaxation, leading to a particularly high mobility of the main chain itself. In contrast, in PEMA, a slow uniaxial diffusion of the main chain around its local axis sets in at T g, the β-process thus affecting mainly the geometry of backbone motions, as is further corroborated by comparing one-dimensional 13C NMR spectra with two-dimensional exchange 2H NMR spectra at higher temperatures. In summary, the coupling of the α- and β-processes leads to longer mean correlation times for the α-relaxation in PEMA than in PMMA.

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