Thermal Analysis of Polymeric Materials

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Graph C, which shows the weight fraction of the melted polymer, wm, illustrates that as the heating rate increases from ‘a’ to ‘f’ more melting of the not perfected crystals occurs at Tm1. The zero-entropy-production melting can be seen at about 433 K. Melting peak Tm3 is seen only at the slowest rate which allows enough time for two recrystallizations during the analysis.

Curve E in Fig. 6.94 shows the minimal change in melting peak temperatures for crystals grown from the melt. These crystals do not reorganize significantly on heating. They remain metastable during the whole analysis. Their zero-entropy- production Tm is 448 K, close to that found for the recrystallized dendrites. Finally, the hedrites in Graph F of Fig. 6.94 are also grown from solution. Their crystal morphology is represented by stacks of lamellar single crystals as seen in Fig. 5.50. As seen in rhw dendrites, the hedrites also have triple melting peaks in their actual DSC traces. Their higher perfection is indicated by the higher Tm for zero-entropy- production melting ( 438 K) than found for dendrites ( 433 K). The thermal analysis of poly(oxymethylene) was an early example of a complete coverage from superheating to lamellar melting with major reorganization bracketing equilibrium and different zero-entropy-production melting temperatures.

6.2.5 Melting of PEEK

Poly(ether ether ketone), PEEK, poly(oxy-1,4-phenyleneoxy-1,4-phenylenecarbonyl- 1,4-phenylene), (O C6H4 O C6H4 CO C6H4 )x, can be considered to be made up of three straight segments joined into a zig-zag structure, each with a phenylene group in its center. The one (O )-group and two (C=O )-groups which link the rigid segments allow rotations about their bonds and induce flexibility. In addition, the phenylene groups may flip about their twofold symmetry axis. As a high-temperature polymer, PEEK is often used as a matrix in composites. The overall behavior of PEEK [57] has similarities to poly(ethylene terephthalate), discussed in in Sect. 6.2.2. The DSC-curve in Fig. 6.95 reveals that on quenching, crystallization can be avoided. The cold crystallization occurs above the glass transition, and crystallization and melting peaks have close to the same size. A more precise analysis would have to take into account the change of the heat capacities and the heat of fusion with temperature, as outlined in Fig. 4.80. The graph at the bottom of Fig. 6.95 represents a summary of the melting-peaks observed on crystallization at different temperatures, as detailed in Fig. 6.96. The line Tm = Tc corresponds to zero-entropy production for crystallization and melting, and was extrapolated to the experimental line of data to give an estimate of the equilibrium, Tmo, of about 668 K (compare with other polymers in Figs. 6.83–87, but see also the mentioned cautions in the discussion of Tmo on pp. 660 to 61). A model of primary and secondary crystallization of PEEK was developed to explain the various melting peaks [58].

The top DSC curves in Fig. 6.96 illustrate the melting behavior after cold crystallization at different temperatures. Enough time was allowed at the crystallization temperature to complete the crystallization and also to undergo annealing at the same temperature. The lower melting crystals give rise to the usual annealing peaks. They are representative of smaller crystals which may grow in the spaces left by the earlier grown, high-melting crystals. The melting peak at high temperature of the

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The bottom DSC-curves in Sect. 3.6.7 (Fig. 3.98) show the DSC curves of the same sample of PEEK as analyzed in Fig. 6.96, but obtained by immediate heating after partial crystallization at a temperature close to case B. The curves indicate that the high-melting crystals grow first, the low-melting ones later, and that there is an intermediate perfection of crystals that does not show up in the DSC traces after crystallization for long times. On cooling at increasing rates, it was found that the rigid-amorphous fraction increases, as is discussed in Sect. 6.3.4.

A more detailed analysis involves DSC, TMDSC, quasi-isothermal TMDSC, and TMDMA [42,59] (for other TMDMA experiments see also Figs. 6.51 and 6.58). Figure 6.97 is a comparison of the apparent total cp from standard DSC and the reversing cp from TMDSC at the same heating rate, and quasi-isothermal TMDSC at short and long times after heating to To. All measured apparent, specific heat

Fig. 6.97

capacities between 500 and 620 K are higher than the heat capacity of the liquid and, thus, must contain latent heat contributions. Compared to PET of Fig. 3.92, the reversing endotherm of melting is twice as high, and because of the identical heating rates for DSC and TMDSC, the melting peak positions are identical.

Slow, primary crystallization of PEEK is analyzed in Fig. 6.98 using long-time, quasi-isothermal TMDSC at 606.5 K, and in Fig. 6.99, using TMDMA at 605 K. The analysis of the TMDSC data is further illustrated in Fig. 6.100 and reveals an increase in heat capacity with time and crystallinity, instead of the decrease expected from a lower heat capacity of a semicrystalline sample. The crystallinity is derived from the integral of the total heat-flow rate, <.(t)>, with time in Fig. 6.98. From Fig. 6.99 it can be seen that the reversing melting reaches a maximum before the crystallization is complete. These data suggest that both crystallinity and crystal morphology need to be considered for the analysis of the reversing melting and crystallization.

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Fig. 6.98

Fig. 6.99

A considerable secondary crystallization, typical for PEEK, is seen in Fig. 6.99. Only about 50% of the total crystallinity can be assigned to primary crystallization. Secondary crystallization starts at about 28 h ( 105 s). At that time, the reversing amplitude of the storage modulus is almost constant. At later times, it decreases slowly, indicating very different processes for primary and secondary crystallization.

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Fig. 6.100

Note that this decrease in reversing melting is similar to the behavior of poly( -capro- lactone) in Fig. 6.58, and contrary to the behavior of PEN seen in Fig. 6.51.

The decrease in the quasi-isothermal, apparent, reversing cp in Fig. 6.97 was followed at 603 K and revealed that the melting endotherm reaches its steady state quicker than the crystallization exotherm. Combining kinetic information on the integral crystallinity of Fig. 6.99 derived from the storage modulus and the change of the reversing melting determined from the TMDSC, the two slow processes were proven to be independent of each other. This supports the assumption that local equilibria in an overall metastable phase structure cause the reversible melting.

6.2.6 Melting of Fibers

Specially difficult systems are represented by fibers of macromolecules. All the topics treated in the prior sections must be considered under the additional aspect of the presence of crystal deformations caused by the drawing process (see Sect. 5.2.6 and 5.3.6) and possible strain retained in the amorphous areas (see Sect. 6.3.3), as well as the existence of strain-induced mesophases (see Figs. 5.69–72 and 5.113–115).

An example of the changes in the thermal analysis results on drawing is given in Fig. 6.101. An undrawn sample as a reference is compared with a drawn sample, measured at constant length and free to shrink. The big difference between a drawn sample free to shrink during analysis, and one restrained can naturally only be caused by differences in the melting or/and annealing behavior since the actual polymer is the same. In fact, extensive experimentation with fibers of known crystal sizes which were either cross-linked to avoid reorganization or etched to remove the amorphous fraction, suggests that the outcome of both experiments with the drawn fibers in Fig. 6.101 is based on the increase of the melting temperature due to residual

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Fig. 6.101

orientation in the amorphous fraction (see Sect. 5.6). The sharper melting peak of the unrestrained fiber is due to a relaxation of the oriented amorphous part while melting. The restrained fibers, in contrast, lose their orientation more gradually and keep the superheating to higher temperature (see Fig. 2.120).

The special behavior of fibers is particularly well investigated on gel-spun, ultra- high-molar-mass polyethylene, UHMMPE. The as-polymerized UHMMPE shows a much higher amount of reversing melting than the lower-molar-mass folded-chain crystals of PE, described in Sect. 6.2.1, or the extended chain crystals of PE and its oligomers of Fig. 6.28. Figure 6.102 displays DSC and TMDSC traces of nascent UHMMPE and illustrates the changes on recrystallization from the melt and on second heating [60]. The apparent reversing heat capacity exceeds the standard DSC result on the low-temperature side of the melting peak. On crystallization from the melt, the reversing latent heat is less than on first heating, but increases on reheating, as was also seen for the lower-molar-mass PE in Fig. 6.30, just that the excess Cp on cooling and second heating is higher than for the lower-molar-mass PE, despite the likely lower crystallinity for the UHMMPE. The frequency dependence of the nascent UHMMPE was also analyzed. It showed a decrease in excess heat capacity reaching the heat capacity measured by standard DSC at 397 K, but at the peak temperature, at 411 K, the maximum of the reversing heat capacity retained is 67 J K 1 mol 1 (compare to Fig. 6.31). Model calculation of the frequency dependence of the reversing excess Cp allowed in this temperature range to speculate about a number of different irreversible and reversible processes. Quantitative information may be available as soon as the morphology is known for the specific sample.

On gel-crystallization from solution, the UHMMPE obtains a lamellar crystal morphology with 13-nm-thick crystals. Stretching such gel-crystallized samples below the melting temperature leads to big changes in the sample morphology.

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data in Fig. 6.103. Figure 6.104 illustrates that with decreasing sample weight, the melting range decreases. For the smallest mass, finally, there is practically no constraint, and direct melting of the orthorhombic crystals is approached, in agreement with Fig. 6.103.

Figure 6.105 illustrates the heat capacity between the glass transition and the beginning of melting, as measured by DSC, using a special step-scan modulation mode of overlapping 20 K heating and cooling steps for two types of samples in a comparison with ATHAS Data Bank information [63]. The fiber A is gel-spun without post-stretch annealing, while the three fibers labeled B are typical commercial

Fig. 6.105

fibers. Subtracting the heat capacity of the 100% crystalline PE (which includes the vibrational and the trans-gauche contributions within the crystal) from the measured Cp leads to the data in Fig. 6.105(b). As expected for drawn polymers, the glass transition, which starts at about 190 K, is broadened and extends to room-temperature. The increase in heat capacity, Cp, is a measure of the noncrystalline material, but does not agree with crystallinity from heat of fusion. To account for the discrepancy, the amorphous and the intermediate, oriented phase must contribute to Cp. To also bring the heat of fusion into agreement with the structural analyses, both the crystalline and the intermediate phase must have a latent heat. Mass, heat capacity, and energy balance suggest for fibers B and A a heat of disordering of the intermediate phase of 30 and 50% of the orthorhombic PE, respectively [63] . From Fig. 6.105, one can conclude that melting begins at about 375 K, but is initially reversible.