Deformation orientation and

Fig. 2. The shape-memory process, where Tis temperature, (a) The cycle where the parent phase undergoes a self-accommodating martensite transformation on cooling to the 24 variants of martensite. No macroscopic shape change occurs. The variants coalesce under stress to a single martensite variant, resulting in deformation. Then, upon heating, they revert back to the original austenite crystallographic orientation, and reverse transformation, undergoing complete recovery to complete the cycle, (b) Shape deformation. Strain recovery is typically ca 7%.

Whether or not a polymer is rubbery or glass-like depends on the relative values of t and v. If t is much less than v, the orientation time, then in the time available little deformation occurs and the rubber behaves like a solid. This is the case in tests normally carried out with a material such as polystyrene at room temperature where the orientation time has a large value, much greater than the usual time scale of an experiment. On the other hand if t is much greater than there will be time for deformation and the material will be rubbery, as is normally the case with tests carried out on natural rubber at room temperature. It is, however, vital to note the dependence on the time scale of the experiment. Thus a material which shows rubbery behaviour in normal tensile tests could appear to be quite stiff if it were subjected to very high frequency vibrational stresses. 

In all cases of the processing conditions, TLCP domains were well dispersed and deformed to droplets in the core layer, but there was only a narrow distribution of their aspect ratio (about Hd 6) and less orientation. In both transition and skin layers, the domains were also well dispersed, but more oriented and fibrillated in the flow direction. From this reason, we give the distribution of aspect ratio Ud) and fiber number (N) versus fiber length class in Fig. 22, only for skin and transition layers, respectively. 

The transition into the oriented state is accompanied by the formation of a neck , a sharp and abrupt local constriction of the sample, in which the extent of orientation and the degree of extension are mudh higher than in the rest of the polymer. After the neck has been formed, further orientation of the sample occurs by spreading of the neck to the entire length of the polymer. When the sample is extended after passing into the oriented state, it undergoes further deformation and at some critical extension it breaks. 

Deformation Schemes and Theoretical Consideration of Orientation Distribution... 

There are two major aspects to this discussion of orientation in polymers. First, there is the question of defining orientation, and the information which can be obtained in principle by any given spectroscopic technique regarding orientation in a polymer. This leads directly to the problem of relating orientation to deformation mechanisms, because this may permit comparatively limited information to be put to optimum use. Secondly, there is the relationship between orientation and physical properties, especially mechanical properties, where such information has been valuable in stimulating and assessing practical developments such as high modulus polymers. 

One of the aims of the present research at Leeds University, of which the spectroscopic studies form a major part, has been to gain an understanding of mechanical behaviour. Both the uniaxially oriented and the biaxially oriented materials discussed in this review have also been the subject of studies of mechanical anisotropy and deformation. It is therefore of some interest to indicate the key guidelines which are emerging from these related studies. 

Figure 6. l-Alanine. Fit to noisy data. Calculation A. MaxEnt deformation density and error map in the COO- plane Map size, orientation and contouring levels as in Figure 2. (a) MaxEnt dynamic deformation density A uP. (b) Error map qME - Model. 

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