Polymer backbone motion

Further, the electron-transport properties of the polymer 7 were enhanced by extending the separation between the redox center and backbone from a single Os—amino linkage to one that extends over 17 bonds. The goal was to provide mobility of the redox center independently of backbone motion, which is necessarily restricted by cross-linking. The mobility of the redox center can be characterized by an apparent diffusion coefficient, Z app- According to the relation proposed by Blauch and Saveant ... 

It was soon realized that a distribution of exponential correlation times is required to characterize backbone motion for a successful Interpretation of both carbon-13 Ti and NOE values in many polymers (, lO). A correlation function corresponding to a distribution of exponential correlation times can be generated in two ways. First, a convenient mathematical form can serve as the basis for generating and adjusting a distribution of correlation times. Functions used earlier for the analysis of dielectric relaxation such as the Cole-Cole (U.) and Fuoss-Kirkwood (l2) descriptions can be applied to the interpretation of carbon-13 relaxation. Probably the most proficient of the mathematical form models is the log-X distribution introduced by Schaefer (lO). These models are able to account for carbon-13 Ti and NOE data although some authors have questioned the physical insight provided by the fitting parameters (], 13) ... 

The other class of motion only now being introduced into interpretive models is oscillatory motion. Anisotropic oscillatory motions of substituent groups have been considered by Chachaty (12) but not in conjunction with a lattice description of backbone motion. No attempt to develop a model based on oscillatory backbone rearrangements is known to these authors, and this avenue may be very important for the interpretation of concentrated solutions, rubbery or amorphous solids, and especially glassy polymers... 

Polymer Backbone Motion. Alternate descriptions of molecular motion utilize an effectively non-exponential autocorrelation function to describe polymer dynamics. One formalism is the use of a log-/2 distribution of correlation times in place of a single correlation time(14). Such a description may simulate the various time scales for overall and internal motions in polymers. 

The acenaphthylene label, being incapable of motion independent of the polymer chain, is expected to monitor segmental relaxation of the macromolecule. The relaxation data adhere well to an Arrhenius relationship over the relatively restricted frequency-temperature range sampled. This behaviour is similar to that observed for the PMA relaxation as determined by phosphorescence depolarization(9,10) but contrasts with that of PMMA in which backbone motion of lesser activation energy was sensible at temperatures inferior to T (11),... 

The activation energies for the backbone motion of P(4HB) and the 3HB and 4HB units in the P(3HB-co-4HB)s, derived from the DEM model analysis, are found to be similar and in the range 42-47 kJ/mol [79]. This range is typical of amorphous polymers at temperatures above Tg, but they are greater than typical ones for polymers in solution, possibly due to the increased apparent viscosity exerted by the amorphous matrix on the moving backbone segment [79]. The activation energy observed for the backbone motion of P(3HB) in chloroform solution is 17 kJ/mol [72]. 

In this paper, we report the first extensive sub-ambient VT-MAS 13C T j and Tj data on macromolecules. The emphasis of the study was placed on isotactic poly(propylene)(PP) and atactic poly(methylmethacrylate)(PMMA) as they represent semi-crystalline and glassy polymers, respectively. Specifics of the investigation were directed to the issue of elucidating sidechain and backbone motions from the high frequency relaxation experiments. 

Localized motions, involving either in-chain movements or side groups laterally attached to the main chain, are the origin this process. This type of local dynamic stays active even when the polymer is in the glassy state [40], that is, when the large length scale backbone motions are frozen. 

Chain flexibility — In the process of aggregation to form a crystalline solid, polymer molecules are opposed by thermal agitation, which induces segmental rotational and vibrational motion. Polymers with flexible chains are more susceptible to this agitation than those with stiff backbones. Consequently, chain flexibility reduces the tendency for crystallization. 

Now, we focus on polymeric materials. Most of polymers have monomeric dipoles classified as type-A, type-B, and type-C. The type-A and type-B dipoles are directly attached to the chain backbone in the directions parallel and perpendicular to the backbone, respectively cf. Figure 3.2. The dielectrically observed fluctuation of these dipoles is activated only by the motion of the chain backbone. In contrast, the type-C dipole (not shown in Figure 3.2) is attached to the side groups and fluctuates through the side group motion even in the absence of the backbone motion. 

In the present study, we report the relaxation of three acrylate polymers in concentrated solutions with chloroform. The polymers were labeled in the methine position so that the backbone motions of the molecules were probed. In order to probe the relaxation, T i and T2 were used because they cover a wide range of spectra densities for the polymer. We have previously reported the behavior of poly(/so-propyl acrylate)-di (PIPA-di) in chloroform(5) where we showed that the relaxation of the polymer could be probed at very high concentrations with one technique. We previously found that at higher concentrations, the relaxation data could not be adequately fit with existing models. However, the subsequent development and correction of models allowed us to extend the concentrations where meaningful interpretations could be made.(6) The present study extends this work so that a comparison can be made among a series of polymers in the same family. 

The incorporation of a low level of plasticizer fills the available free volume, the lifetime of the o-Ps is decreased and the temperature to which the solid has to be raised before polymer backbone motion can occur is increased. The phenomenon of the addition of a diluent raising Tg is called antiplaticization and is commonly observed for low levels of plasticizer. Further increase in the plasticizer levels leads to a decrease in the temperature at which the lifetime plots change slope indicative of plasticization of the polymer by the diluent. 

The first exponential term represents the transport of molecules to the growing nucleus with U the activation energy for this process and (T - To) the temperature difference between the crystallization temperature T and the temperature at which backbone motions substantially cease. For most polymers To is about 50°C below the glass transition temperature. The second exponential term represents the work required to form a critical nucleus where TS, is the melting point of an infinitely thick lamellar crystal, AT is the supercooling and AH the heat of fusion. The term C contains various interfacial energy terms and depends upon the precise mechanism of the nucleation process. For homogeneous nucleation... 

From the threshold measurements one concludes that, for long spacer lengths where the side-chain mesogenic moiety motion is essentially decoupled from the backbone motion, the low molar mass compounds and the polymers have similar elastic constants. If, however, the flexible spacer length in the polymer is short, then the elastic constants (as related through k in eqn (4)) increase markedly. This result is also implicit in the work of Ringsdorf and Zentel on side-chain polyacrylic esters. ... 

If the shear rate is higher than the time for the first normal mode, the chain does not have time to respond to the applied perturbation, and only the higher modes are able to be activated. In other words, at times which are shorter than Tj the first normal mode is frozen out and hence cannot contribute to the observed viscosity. Further increase in the rate of shear will progressively remove further modes until the viscosity falls to a value which corresponds to that of the solvent. This simple description, with minor modifications, describes the behaviour of most polymer molecules in dilute solution. Because in solution the backbone motions are effectively liberated, so that the chains are fuUy flexible, the description of the viscosity of dilute polymer solutions is essentially independent of the chemical nature of the molecules. The modes are purely defined by the end to end length of the polymer chains and hence by the molar mass of the polymer. 

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