Thermal Analysis of Polymeric Materials Part 12 pps

6 Single Component Materials ___________________________________________________________________ 646 Fig. 6.64 exotherm in the insert in Fig. 6.62. The following annealing of 70 min starts at ‘3’ with only a slight increase in diffraction intensity. The subsequent heating to 433 K is shown in the lower part of Fig. 6.63. It reveals, again, a period of constant intensity ‘1’ followed by an increase in crystallinity ‘2’ which leads to the beginning melting at ‘3’, which on annealing at ‘4’ indicates some recrystallization and a further, slight increase in crystallinity. The next heating leads to the main melting with disappear- ance of the crystalline diffraction pattern, in accord with the DSC trace. Further analysis of the two samples of Fig. 6.62 is done by TMDSC, as seen in the left graphs of Fig. 6.64. A comparison of the two reversing heat capacities shows that the cold crystallization and the transition mesophase-to- -monoclinic crystals do not show, i.e., they are nonreversing. There remains, however, a substantial reversing contribution which is larger for the quenched iPP than for the lamellar, melt-cooled iPP. The upper limit of the devitrification of the RAF seems to occur at 320 330 K, but cannot be separated fully from the reversing melting. The curves on the right in Fig. 6.64represent the latent-heat contributions of the apparent, reversing c p , obtained by subtracting the expected thermodynamic c p for the given crystallinity from the curves on the left. Results from quasi-isothermal experiments on heating are plotted also in Fig. 6.64. These are taken after the slow crystal perfection ceased and represent reversible latent heats. The slowly-cooled, lamellar sample begins to show reversible melting at 320 K, while the fast-cooled sample with globular morphology begins melting at the end of the glass transition. More qualitative data were generated throughout the melting peak and show a maximum in reversible melting on heating after fast cooling of 3.05 J K 1 mol 1 and after slow cooling of 2.80 J K 1 mol 1 [52]. The same thicknessin the chaindirection and thedifference in lateralextension places the reversible melting for iPP on the growth faces of the crystals. 6.2 Size, Extension, and Time Effects During Fusion ___________________________________________________________________ 647 Fig. 6.65 Figure 6.65 further supports this interpretation of reversible melting at the lateral growth face with an analysis of iPP samples annealed for long times at different temperatures. The standard DSC traces were takenafterthequasi-isothermal TMDSC could detect no further decrease in the apparent reversing heat capacity and were used for a measurement of the remaining crystallinity. Then, corresponding X-ray diffraction differences were taken in a temperature range of ±1.5 K after the metastable, global equilibrium was achieved in this temperature region. On the right side, Fig. 6.65 displays the normalized difference patterns of the X-ray diffraction experiments. Within experimental error, the increase in latent heat, marked in Fig. 6.64 as open squares, and the increased differential diffraction areas in Fig. 6.65 give the same changes in crystallinity [53], making it certain that the reversible melting refers to the same process seen in the irreversible melting and crystallization. Syndiotactic polypropylene, sPP, isknown almost as long as iPP, but only became readily availablewith high stereospecificity afterthe discoveryof soluble, single-site, metallocene catalysts (see Sect. 3.2.1). Several semiquantitative studies are available and have been reviewed in Ref. [1]. Figure 6.66 illustrates TMDSC traces of an sPP sample of low crystallinity (and stereospecifity). The reversing c p of the semi- crystalline state on initial cooling and on subsequent heating are identical. Above the low-temperature glass transition, the sample has a solid fraction of 59%. With a heat of fusion of 17%, this corresponds to an RAF of 42%. A broad glass transition of the RAF is indicated between 315 and 365 K. Above this glass transition, reversing melting is observed with both of the irreversible melting peaks showing a reversing component. The frequency dependence of the reversing heat capacity of sPP after isothermal crystallization at 363 K is shown in Fig. 6.67 [54]. At low frequency the sample shows reversing melting which at higher frequency reverts to a partially 6 Single Component Materials ___________________________________________________________________ 648 Fig. 6.66 Fig. 6.67 devitrified RAF. Both iPP and sPP are thus examples with a broad devitrification of the RAF before the major melting peak. The frequency dependence of the reversing specific heat capacity of sPP over the full temperature range is displayed in Fig. 6.68 and should be compared to similar results on other polymers (see, for example, the results for nylon 6in Figs. 4.116 and4.119). The frequency dependencewas extracted from the higher harmonics of the heating rate and heat-flow rate. 6.2 Size, Extension, and Time Effects During Fusion ___________________________________________________________________ 649 Fig. 6.68 Decoupling of segments of polymer chains was proposed as a mechanism for the examples of the reversing and reversible melting which are summarized with their properties in Table 6.1 [1,33]. To restate the main facts, it is shown in Figs. 3.75 and 3.91 that flexible molecules melt reversibly only up to a critical length, which for paraffins is 10 nm. Longer extended-chain crystals melt irreversibly (Figs. 3.89 and 6.28). A first observation of decoupling of melting concerns the limiting reversibility of crystals grown froma distribution of molecules of low molar mass. In thiscase two physical changes that should occur simultaneously are partially decoupled. The melting/mixing and crystallization/demixing are incomplete, probably due to slow diffusion in the available time. This can cause the limited reversibility seen in Figs. 6.25 and 6.35. Decoupling is also known for polymers with long side-chains, where side-chains may be largely independent of the main chain (see Sect. 7.2.3). Locally reversing and reversible melting within a globally metastable structure of semicrystalline, flexible macromolecules also exhibit decoupling. Such decoupling allows molecular nucleation to start with crystallization of segments of the polymer chains (see Fig. 3.74). The portions of the molecules not crystallized must be considered as amorphous and belonging to a different phase. Figure 6.69 illustrates such creation of points of decoupling and a possible mechanism of reversible melting at the growth face. Figure 6.69 is linked to the molecular nucleation in Fig. 3.74. Because of the transfer of stress across the interface and along the continuing molecule, this picture also offers an explanation of the broadening of the glass transition and the creation of a separate RAF. In case T g of the RAF is close to T m , melting is governed by the RAF and not the crystal, as documented in Fig. 6.56. For polyethylene, the melting point of decoupled segments can be estimated from the known T m of folded and extended-chain crystals, given in Figs. 2.90 and 3.04, as is shown in Fig. 6.70. The bottom one of the two equations permits the computation 6 Single Component Materials ___________________________________________________________________ 650 Table 6.1. The reversing behavior of the various polymers a Polymer cold cryst. hot cryst. confor- mational C p RAF ( b ) reversible melting reversing melting PET 88 0<T m 88 PTT 8 – 8 <T m 88 PBT 88 ? T m 88 PEN 8 – 8 <T m 88 PEEK – 88<T m 88 PC 8 –0T m 0(?) 8 PPO – – 0 >T m 00 PTFE c – 8 ? T g 88 nylon 12 – 88T g 88 nylon 6 – 88T g 88 PHB 8 –0<T m 88 PCL – 8 0 T g 88 POM – – 0 T m –– POE – 88T g 88 POTM – 88T g 88 PPDX 8 –0T g 88 iPP 88 0<T m 88 sPP – 8 0<T m 88 PE – 88T g 88 PE-alt-CO 88 0 T g 88 a 8 indicates “yes”; 0, “no”; –, not discussed or measured; ?, “questionable.” The table has been adapted from [1]. b T g of the RAF below (<), up to (), at (), or above (>) the melting temperature, T m ; T g indicates a broadening of the glass transition only. c Only the low temperature transitions of PTFE have been analyzed (see Fig. 5.31). See Sect. 5.5 also for TMDSC of other mesophase transitions. 6.2 Size, Extension, and Time Effects During Fusion ___________________________________________________________________ 651 Fig. 6.70 Fig. 6.69 of the extended-chain length of the decoupled segment. Inserting its melting temperature into the upper equationinstead of 414.2 K (themelting temperature of the extended-chain crystal ofinfinite length) and correcting themelting point lowering(in brackets) for the fraction of the surface covered with chain folds, yields the shown curves. (For one fold, the correction factor for the term in brackets is 0.50, for two folds, 0.67, for three, 0.75, for four, 0.80, and for five, 0.83.) 6 Single Component Materials ___________________________________________________________________ 652 Fig. 6.71 Poor crystals of polymers which are susceptible to reversing melting are either small, nanophase-separated crystals, as just discussed, or may be conformationally disordered. The conformationally disordered crystals are mesophases, described in Sect. 5.5. The reversing and reversible melting of the mesophases depends on the mobility within thecrystal, as seenfromFigs. 5.131, 5.144, and5.154. Liquid crystals are likely to have fully reversible transitions within the time scale of calorimetry, while the condis phases are sufficiently rigid to show less reversing behavior and are closer intheir behavior to the small polymer crystals with their slow crystal perfection due to melting and recrystallization and reversibility on the growth faces. Quenching from the melt or solution can lead directly to glasses (see Sect. 6.3), to small crystallites, dendrites, and spherulites (see Figs. 5.56, 5.60, and 5.75), or to mesophase glasses (see Fig. 5.146). Small crystals are also found on cold crystalliza- tion of a glass by growing crystals as little above the glass-transition temperature as possible and by drawing of fibers or other deformations (see Sects. 5.2.6 and 5.3.5). In addition, by partial crystallization at higher temperature it is possible to set up a network of larger crystals, separated by amorphous defects, as described by Fig. 5.87. On cooling, these amorphous defects may generate a second population of poor crystals, observed calorimetrically in the form of “annealing peaks” (see Sect. 6.2.5). In all these cases there is a chance to set up a configuration for locally reversible melting and crystallization. One expects these poorly grown crystals to be close to a fringed-micellar structure, schematically shown in Fig. 5.42, with many of the molecular chains being decoupled into crystals and amorphous segments. A unique method to generate poor crystals is shown in Fig. 6.71. It involves the compression of crystals of poly(4-methyl-1-pentene), P4MP1, with a glass transition T g = 303 K, melting transition T m o = 523 K, and a heat of fusion of H f =10kJmol 1 (see also Fig. 5.124). At room temperature, the tetragonal crystals of P4MP1 have a 6.2 Size, Extension, and Time Effects During Fusion ___________________________________________________________________ 653 Fig. 6.73 Fig. 6.72 lower density than the surrounding amorphous phase. The rather open helical conformation 2*7/2 for P4MP1 is the reason for the poor packing in the crystals at room temperature, compared to the amorphous phase. Increasing the temperature expands the amorphous phase faster than the crystals, so that at T m the crystals are somewhat denser. The change of T m as a function of pressure is shown in Figs. 6.72 and 6.73. After the initial, expected increase with pressure, T m decreases. 6 Single Component Materials ___________________________________________________________________ 654 Fig. 6.74 This unique behavior of the melting temperature with pressure indicates that increasing the pressure reverses the density difference in a similar manner as decreasing the temperature. At the maximum of the T m versus p curve in Fig. 6.73, the difference in molar volume between melt and crystal is zero, leading to a pressure of coefficient of zero in the Clausius-Clapeyron equation discussed in Sect. 5.6.5 and written in Fig. 5.168. In Fig. 6.73 a phase diagram for the various states is drawn based on high-pressure calorimetry. Only the melt appears as a true equilibrium phase. All other phase areas are semicrystalline, i.e., do not follow the phase rule (see Sect. 2.5.7). Because Fig. 6.73 does not represent equilibrium, it can mirror the actual processes. All exotherms and endotherms, except for melting, are small (0.5 to 2 J/g), indicating thatat the transitions, the changes in entropyof the phases are small. Major crystallization and melting, thus, is not possible. The nonequilibriumstates were identified and seem to be initiated in close concert with the existence of the glass transition of the majority glassy phase identified in Fig. 6.73. Isothermal disordering at room temperature along the horizontal arrow causes a transition to the condis glass on the right of the extrapolated phase boundary. It seems to be a frustrated tetragonal to trigonal transition. Although superficially showing an amorphous X-ray pattern in Fig. 6.71, the condis glass retains most of its heat of fusion and acquires largely sessile backbone and side-chain conformational defects. This disordering can be reversed by reduction of pressure. Based on Fig. 6.73one can speculate that the equilibrium phase-diagramisas given in Fig. 6.74. The tetragonal phase is the stable crystal. The trigonal crystal form with a 2*3/1 helix conformation, common for vinyl polymers with smaller side chains (see Fig. 5.14), exists as the low-temperature,high-pressurephase. As expected from a high-pressure, low-temperature phase, it is denser than the trigonal crystals. Neither of the two equilibrium polymorphs has dynamic conformational disorder. 6.2 Size, Extension, and Time Effects During Fusion ___________________________________________________________________ 655 Fig. 6.75 6.2.3 Annealing and Recrystallization Effects The word annealing derives from the Anglo-Saxon ælan, to bake or burn, and the prefix an, on. It was used to describe the burning-on of a glaze or enamel on ceramics and metals, and furthermore the strengthening, hardening, or toughening by heating. In the polymer field, the term implies imparting of a certain (advantageous) property by heat treatment without large-scale melting. If large-scale melting of crystals is involved in the process and renewed crystallization occurs before the whole sample turns liquid, the process is called recrystallization rather than annealing. Naturally, both processes may occur at the same time in a semicrystalline sample. Glassy polymers are mainly annealed close to the glass transition for strain release and densification (seeSect. 6.3), while crystalline polymers perfect theirinternal structure and their morphology on annealing (see Sects. 5.1–3). Crystal structure and morphology are best assessed byX-ray diffraction and electron microscopy, while the change in macroscopic properties is, as usual, best tested by thermal analysis. Point defects and dislocations are shown in Sect. 5.3 to appear as nonequilibrium defects, usually introduced during crystal growth or deformation, and as equilibrium defects, generated thermally. The dislocation density can be determined by the moiré method (see Appendix 17 and Fig. 5.93). On annealing of crystals above their crystallization temperature, the nonequilibrium dislocation density increases, as is demonstrated in the left graph of Fig. 6.75. Since the number of such nonequilibrium defects cannot increase by annealing without a cause which more than compensates their metastability, one assumes that the cause is a change in the packing within the crystals or between the lamellae that cause the moiré pattern. The right graph in Fig. 6.75 illustrates, based on X-ray diffraction, that the mosaic dimensions, indeed, [...]... electron micrograph of Fig 6.78 depicts melt-grown polyethylene lamellae of thicknesses of 18 22.5 nm, grown at 403 K on the surface of the (110) growth face of an extended-chain crystal, as seen in Fig 5.76 and 5.77 [24] This substrate lets one see the growth face of the folded-chain lamellae In the bottom of the figure the lamellae were annealed below the melting temperature of the substrate, at... temperature The analysis of crystallinity after quenching to room temperature showed that relatively quick melting is followed by recrystallization Above 518 K practically all crystals melt first This temperature is at the beginning of the melting peak of a standard DSC trace, as given in Fig 3.92 A qualitative correspondence exists between the dilatometry of 1960 and the analysis of the instantaneous,... thicknesses of 27 to 45 nm The underlying extended chain crystals are little affected by the annealing 6.2 Size, Extension, and Time Effects During Fusion 657 _ Fig 6.77 Fig 6.78 This short summary of microscopic observations of annealing reveals that there are at least two processes that must be considered when interpreting the macroscopic annealing experiments of thermal analysis: ... 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... 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... 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, ultrahigh-molar-mass polyethylene,... 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... The maximum of the reversing excess heat capacity is double as much as for the polyethylene of lower molar mass in Fig 6.30, and occurs at higher temperature As one goes into the double melting peak of Fig 6.106, which was ascribed to melting of orthorhombic crystals, transition into the hexagonal phase, and melting of the hexagonal phase, the reversible contributions ( ) are only a fraction of the reversing... original crystal morphology (2) A change in crystal shape towards equilibrium, which may involve several stages of crystal thickening 658 6 Single Component Materials _ The macroscopic thermal analysis, furthermore, proves that it is also common for melting of the original crystals to be followed by recrystallization to a more stable morphology Finally, some crystals... of the instantaneous, complete melting of thin films pressed against a hot surface at 518 K [54] The completion of melting was identified by the nature of the viscous flow at the high temperature and by dilatometry and DSC after quick quenching between cold plates In the same research isothermal primary melt crystallization was shown to be followed by insertion of secondary lamellar stacks and continuous . disappear- ance of the crystalline diffraction pattern, in accord with the DSC trace. Further analysis of the two samples of Fig. 6.62 is done by TMDSC, as seen in the left graphs of Fig. 6.64 glass transition, the sample has a solid fraction of 59%. With a heat of fusion of 17%, this corresponds to an RAF of 42%. A broad glass transition of the RAF is indicated between 315 and 365 K (Figs. 3.89 and 6.28). A first observation of decoupling of melting concerns the limiting reversibility of crystals grown froma distribution of molecules of low molar mass. In thiscase two physical