Semicrystalline polymers experimental

The SFA, originally developed by Tabor and Winterton [56], and later modified by Israelachvili and coworkers [57,58], is ideally suited for measuring molecular level adhesion and deformations. The SFA, shown schematically in Fig. 8i,ii, has been used extensively to measure forces between a variety of surfaces. The SFA combines a Hookian mechanism for measuring force with an interferometer to measure the distance between surfaces. The experimental surfaces are in the form of thin transparent films, and are mounted on cylindrical glass lenses in the SFA using an appropriate adhesive. SFA has been traditionally employed to measure forces between modified mica surfaces. (For a summary of these measurements, see refs. [59,60].) In recent years, several researchers have developed techniques to measure forces between glassy and semicrystalline polymer films, [61-63] silica [64], and silver surfaees [65,66]. The details on the SFA experimental procedure, and the summary of the SFA measurements may be obtained elsewhere (see refs. [57,58], for example.). 

For a layer-stack material like polyethylene or other semicrystalline polymers the IDF presents clear hints on the shape of the layer thickness distributions, the range of order, and the complexity of the stacking topology. Based on these findings inappropriate models for the arrangement of the layers can be excluded. Finally the remaining suitable models can be formulated and tested by trying to fit the experimental data. 

Through the use of multiple experimental techniques, we have shown how both the NXL and XL phases of PILE interact and respond to applied tensile deformation. Strains transmitted to PILE crystals lead to two distinct slip modes and, at higher strains, to the breakup and alignment of lamellar fragments. In our experiments, crystallites in PTFE orient fuUy with respect to the draw direction at strains between 70 to 200%. With increasing strain, some chains originally in the XL phase are transformed to NXL material. Noncrystalline chains continue to orient until macroscopic failure is reached. This could be a fairly general microstructural response for semicrystalline polymers. 

Experimental studies of the temperature changes and heat effects resulting from reversible deformations of glassy and semicrystalline polymers have been carried out since the 1950s. The early results were summarized and analysed by Muller 1). Since that time, a number of experimental results on thermomechanical behaviour of glassy and semicrystalline polymers have been obtained. This part of our review is dealing with critical consideration of these results. 

In Fig. 27 experimental 1/e values are represented as a function of the temperature for samples with different compositions. As shown, the experimental values qualitatively follow Eq. 7 although 1/e achieves a non-zero value at the critical temperature. The intrinsic composite structure of semicrystalline polymers has been invoked to understand this effect [8, 6]. The order of magnitude of the constant A has been reported to be around 103 °C [11] which is consistent with the relatively high polarizability of these materials. At this point it is important to emphasize that the knowledge of morphological aspects of these copolymers might help, in future, to develop a theoretical framework capable of accounting for the experimental observations. 

A number of theories of the contribution of interdomain polymeric material to the stress-strain, modulus, and swelling behavior of block copolymers and semicrystalline polymers are examined. The conceptual foundation and the mathematical details of each theory are summarized. A critique is then made of each theory in terms of the validity of the theoretical model, the mathematical development of the theory, and the ability of the theory to explain experimental findings. 

Diffusion requires cooperative motions of both the polymer and the diffusant [27] and is therefore only low below T , and severely restricted below Tg. The diffusion of various stabilizers was elucidated in amorphous and semicrystalline polymers and in multiphase systems. In semicrystalline polymers, diffusion takes place almost exclusively in the amorphous phase and the value of D is sensitive to the total crystallinity and the morphology. It is difficult to predict D within a homologuous series of stabilizers. For a given molecular weight, long and flexible molecules diffuse more rapidly than more rigid and compact structures. For a given stabilizer, the value of D is usually lower in PP and HDPE than in LDPE. It was demonstrated that typical AO molecules have a very restricted mobility in polymers. They are, however, insufficient experimental data to correlate the AO mobility and the AO efficiency. 

Despite those difficulties, some reactive systems have been used to study the reinforcement of interfaces between semicrystalline polymers. Within the context of interfacial fracture, the possibility of crystallization introduces two new important aspects that need to be taken into account when interpreting experimental data ... 

They have been combined with FV models and they have been applied to semicrystalline polymer-solvent systems. The results are satisfactory but they are not predictive the and / parameters should be estimated from experimental data. However, the swelling of cross-linked polymers can be estimated with such equations. In one of the very few works reported on the prediction of SLLE for polymer solutions, the Entropic-P/ and UNIFAC models have been compared for semicrystalline polymer-solvent systems. These two models are shown to yield similar results for SLLE. 

In this section, the molar volume at room temperature will be denoted by V rather than V(298K) for the sake of brevity. Experimental values of V for amorphous polymers, as well as for the amorphous phases of a large number semicrystalline polymers, were gathered for 152 polymers. Extensive lists provided by van Krevelen [1] were used as the main sources of data, and supplemented or updated with data provided by several other sources [17-25],... 

The amorphous phase of a semicrystalline polymer will follow Cp T), while the crystalline phase will follow Cps(T), between the glass transition temperature Tg and melting temperature Tm. Measurement of Cp(T) as a function of percent crystallinity can therefore enable the extrapolation of both Cps(T) and Cp (T), as limits at 100% and 0% crystallinity, respectively, in this temperature range. Experimental values of both Cps and Cp1 can thus be determined between Tg and Tm if measurements are performed on samples of different percent crystallinity. 

Table 9.2. Experimental molar volumes V(exp) at 298K in cc/mole, predicted molar volumes V(pred) at 298K calculated by using equations 3.13 and 3.14, and experimental and predicted values of the molar polarization PLL in cc/mole, for 61 polymers. The calculations also utilize the experimental and fitted values of the dielectric constant, which are listed in Table 9.1. The V(exp) values listed for semicrystalline polymers are extrapolations to the amorphous limit. LL(exP) was not calculated for six polymers because V(exp) was not known.

Experimental refractive indices were used whenever available in these calculations. For semicrystalline polymers such as polyethylene, typical refractive indices of the semicrystalline specimens were used instead of taking the amorphous limit as was done whenever possible in Section 8.C. Whenever experimental refractive indices were not available, the best possible estimate was made for the refractive index, in most cases by using Equation 8.6. 

The Chow equations [5] and the Halpin-Tsai equations [8,9] are also useful in modeling the effects of the crystalline fraction and of the lamellar shape (see Bicerano [23] for an example) on the moduli of semicrystalline polymers. Grubb [24] has provided a broad overview of the elastic properties of semicrystalline polymers, including both their experimental determination and their modeling. Janzen s work in modeling the Young s modulus [25-27] and yielding [27] of polyethylene is also quite instructive. 

The dependence of E on the molecular diameter of the penetrant, calculated for several species of penetrants in a number of amorphous and semicrystalline polymers to agree acceptably well with experimental values ( 4). ... 

All these commercially produced BR systems are amorphous rubbers under atmospheric conditions. The Tg-value of these systems, depending on their structure, is described by the Gordon-Taylor relation, see Chapter 1. BR becomes a semicrystalline polymer under atmospheric conditions if the 1,4 trans-BR content is higher than about 70 %wt. or if a syndio-or isotactic 1,2-BR phase is present. This is shown by the results of thermo-analytical measurements on experimental BR systems with a high trans content and with a high syndiotactic 1,2-BR content which are reported in this chapter. Moreover, the Tg-values of two series of IR samples containing 1,2- and 3,4-IR are used to determine the Tg/structure relation for non-polar polymers with side-chains. 

Experimental Correlation Function y(r) Following a similar process, Vonk and Kortleve introduced the correlation function for semicrystalline polymers, but in a single dimension (Fig. 19.8) [20]. If the cosine Fourier transform of the correlation function is obtained, then... 

Because of the usual experimental ranges of crystallinity, the appearance of the baselines in the study of semicrystalline polymers is not very frequent [36, 37]. Therefore, the crystalline volume fraction should be obtained by other alternative techniques such as, for example, differential scanning calorimetry. In this way, if the crystallinity in volume is known, it is possible to obtain the lost baseline [37, 38]. However, if the crystallinity is either low or high, it is possible to directly obtain the baseline for which it is foreseen that the morphology is solved [36]. For example, Vonk and Pijpers and Vonk and Koga worked at high crystallinities and they observed the aforementioned baselines [39]. 

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