Chain orientation semicrystalline polymers

Figure 8.2 Schematic diagram indicating the structure of a highly oriented semicrystalline polymer. A, amorphous region B, intercrystalline bridges C, chain-folded crystal block TM, tie molecules.

These are just some typical responses that may be observed in an un-oriented semicrystalline fiber while a tensile stress is applied. In reality, the chemical and crystalline stmctures affect how the polymer chains behave under stress. As a result, for some un-oriented semicrystalline polymer fibers, the responses of polymer chains under tensile stress may be different from what is discussed above. 

The effects of processing will be illustrated by considering injection moulding of a semicrystalline polymer. The molten plastic is injected into the mould under high pressure and temperature. The edges of the mould retard flow and cool more rapidly, leading to a boundary layer of high shear, which in semicrystalline polymers leads to orientation of the polymer chains and of short fibre reinforcements parallel to the direction of flow. At the centre the structure is less oriented. Where two separate flow streams meet, there is a... 

With the exception of PC, amorphous, non-oriented polymers did not produce measurable amounts of broken segments when subjected to tension. As has been shown in previous paragraphs, large axial stresses capable of chain scission in amorphous polymers can only be transmitted into the chain by friction of slipping chains requiring strong intermolecular interactions. In addition, macroscopic fracture occurs before a widespread chain overloading and scission occurs, which is opposite to the behavior of semicrystalline polymers. 

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. 

The phenomenon of strain hardening in polymers is a consequence of orientation of molecular chains in the stretch direction. If the necked material is a semicrystalline polymer, like polyethylene or a crystallizable polyester or nylon, the crystallite structure will change during yielding. Initial spherulitic or row nucleated structures will be disrupted by sliding of crystallites and lamellae, to yield morphologies like that shown in Fig. 11-7. 

Practical problems associated with infrared dichroism measurements include the requirement of a band absorbance lower than 0.7 in the general case, in order to use the Beer-Lambert law in addition infrared bands should be sufficently well assigned and free of overlap with other bands. The specificity of infrared absorption bands to particular chemical functional groups makes infrared dichroism especially attractive for a detailed study of submolecular orientations of materials such as polymers. For instance, information on the orientation of both crystalline and amorphous phases in semicrystalline polymers may be obtained if absorption bands specific of each phase can be found. Polarized infrared spectroscopy can also yield detailed information on the orientational behavior of each component of a pol3mier blend or of the different chemical sequences of a copoljnner. Infrar dichroism studies do not require any chain labelling but owing to the mass dependence of the vibrational frequency, pronounced shifts result upon isotopic substitution. It is therefore possible to study binary mixtures of deuterated and normal polymers as well as isotopically-labelled block copolymers and thus obtain information simultaneously on the two t3q>es of units. 

In addition, most semicrystalline polymers, particularly those produced commercially, are partially oriented i.e. their chains have an overall alignment that may impart to the bulk polymer certain advantageous properties, e.g. increased mechanical strength or dielectric polarizability. Molecular orientation, whether arising from crystallization under stress or deformation of a solidified polymer, or in naturally occurring oriented crystalline polymers such as cellulose or keratin, is always associated with an orientational morphology. 

Semicrystalline polymers may crystallize from solution, as well as from the melt, in the form of chain folded lamellar crystals. The high spatial resolution of ATM enables one to assess lamellar thicknesses from images of these lamellar crystals in edge-on and flat-on orientation. As discussed in this section, images of flat-on oriented lamellae are particularly suitable for a quantitative determination of lamellar thicknesses. 

The deformation of long chain polymer molecules has always been of great industrial interest as more value can be placed on a material that has improved properties. Molecular extension, or alternatively molecular orientation, is of particular interest as it can enhance mechanical properties of an otherwise weak polymer. Oriented materials are inherently anisotropic. These anisotropic regions can be found directly in semicrystalline polymers where chains organize themselves into crystalline domains. 

The effects of orientation via mechanical deformation on Tg have been reviewed [65]. Tg increases in those amorphous regions of a semicrystalline polymer that are either attached to crystallites or so close to them that their chain segment mobilities are hindered because of the interference of the crystallites. On the other hand, orientation has little effect on Tg in amorphous regions far away from crystallites as well as in completely amorphous polymers. 

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