Amorphous ductile-brittle transition

Within this framework, one can imagine a two dimensional ductile-brittle transition criterion which would be composed of two perpendicular boundaries in the graph la-Mw (Fig. 7). The vertical boundary would correspond to a critical molar mass M c- In polymers of low crystallinity or fully amorphous, M c would be sharply linked to the entanglement molar mass M c 5Me. In polymers of medium to high crystallinity, M c would be the molar mass below which it is impossible to have values higher than lac- The horizontal boundary would correspond to U = lac-... 

The temperature dependence of tensile behavior is summarized in Table 9.2. In this table, open circles represent ductile behavior without stress whitening, and closed circles brittle behavior with stress whitening. Bold lines in the table denote Tg. The table shows that the ductile-brittle transition in the immiscible blends is independent of temperature, suggesting that the fracture will be associated with the boundary separation at the interface between iPP matrix and EHR domains. However, the EHR51 material plays a role of a plasticizer for iPP because the addition of EHR51 lowers the Tg region. This will be because of the incorporation of the EHR molecules into the amorphous region of iPP. The similar results were obtained in the iPP/EBR blends (50). 

Figure 130 shows the thermograms of the La55Al2sNi2o amorphous samples annealed for 1.8 ks at different temperatures ranging from 390 to 450 K including the ductile-brittle transition temperature (Inoue et al. 1991b). Here, and Cpji denote the specific heat curves of the as-quenched and annealed samples, respectively, and Cp represents the... 

In conclusion, the deformation behavior of poly(hexamethylene sebacate), HMS, can be altered from ductile to brittle by variation of crystallization conditions without significant variation of percent crystallinity. Banded and nonbanded spherulitic morphology samples crystallized at 52°C and 60°C fail at a strain of 0.01 in./in. whereas ice-water-quenched HMS does not fail at a strain of 1.40 in./in. The change in deformation behavior is attributed primarily to an increased population of tie molecules and/or tie fibrils with decreasing crystallization temperature which is related to variation of lamellar and spherulitic dimensions. This ductile-brittle transformation is not caused by volume or enthalpy relaxation as reported for glassy amorphous polymers. Nor is a series of molecular weights, temperatures, strain rates, etc. required to observe this transition. Also, the quenched HMS is transformed from the normal creamy white opaque appearance of HMS to a translucent appearance after deformation. 

Mechanical Properties. The room temperature modulus and tensile strength are similar to those of other amorphous thermoplastics, but the impact strength and ductility are unusually high. Whereas most amorphous polymers arc glass-like and brittle below their glass-transition temperatures, polycarbonate remains ductile to about — 10°C. The stress-strain curve for polycarbonate typical of ductile materials, places it in an ideal position for use as a metal replacement. Weight savings as a metal replacement are substantial, because polycarbonate is only 44% as dense as aluminum and one-sixth as dense as steel. 

The first secondary transition below Tg, the so called fj-relaxation, is practically important. This became evident after Struik s (1978) finding that polymers are brittle below Tp and establish creep and ductile fracture between Tp and Tg. The p-relaxation is characteristic for each individual polymer, since it is connected with the start of free movements of special short sections of the polymer chain. In view of more recent data of Tp Boyer s relation, Eq. (6.29), is very approximate and fails completely for amorphous polymers with high Tg s (e.g. aromatic polycarbonates and polysulphones). Some rules of thumb may be given for a closer approximation. 

We have seen already (Sect. 13.4.7) that every amorphous material (including that in semi-crystalline polymers) becomes brittle when cooled below the first secondary transition temperature (Tp) and becomes ductile when heated above the glass transition point (Tg). Between these two temperatures the behaviour - brittle or ductile - is mainly determined by the combination of temperature and rate of deformation. 

Finally, when the material is melted and crystallized, the material is tough and ductile compared with the single-crystal material. Therefore, this analogy is useful to understand the hypothesized transition from brittle to ductile behavior based on increased crystalline-amorphous interactions. 

Conversely, nodules can significantly increase the amount of energy dissipated during deformation of notched specimens (Fig, 5., 6.). They contribute to decrease brittle to ductile transition temperature. Non-reactive nodule can toughen amorphous PET to a certain extend, being nevertheless always less efficient than reactive one. On the contrary, only the reactive nodules exhibit a certain level of efficiency in semi-crystalline PET, provided that concentration is at least 21 %. 

At room temperature, PP is close to its Tg(0-25°C) and well above its normal brittle-ductile transition temperature ( -30°C). However the presence of surface cracks in the photo-oxidized film is apparently sufficient to promote brittle failure at room temperature. According to the Griffith crack theory, once a critical crack length has been exceeded, a critical crack velocity is required to propagate the crack. If this velocity is not exceeded, cold drawing of the amorphous zones ensues. 

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