Thermal Analysis of Polymeric Materials
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magnified plot of the graph in the time period of fast, primary crystallization and illustrates the sawtooth-like modulation of the temperature. To interpret the data, it must be noted that the storage shear-modulus, G', is proportional to crystallinity. The increase in crystallinity stretches over a time interval of several days. At the end of the fast, primary crystallization, the amplitude of the storage modulus, which is a measure of the reversing melting, also continues to increase, in contrast to PET, where it reaches a constant value. Analogous TMDMA experiments for melt crystallization of poly( -caprolactone) (PCL) are discussed below in Fig. 6.58 and for poly(ether ether ketone), PEEK, in Sect. 6.2.5 (Fig. 6.99).
Polycarbonate, specifically poly(4,4'-isopropylidenediphenylene carbonate), PC, has the repeating unit:
3O
CH
(C6H4 C C6H4 O C O )x .
CH3
Its flexibility and segmental mobility are reduced so much that the crystallization of a quenched, amorphous sample of molar mass 28,000 Da needs an induction time of almost 50 h at 460 K, the temperature of fastest growth, and the half-time of crystallization is only reached after about 7 days [43]. The heat capacity of the solid PC has been analyzed, and the heat capacity of the liquid PC was measured and compared to the other aromatic polyesters.
The reversing specific heat capacity in the glass transition region is illustrated in Fig. 6.52 [21]. The analysis in terms of the ATHAS Data Bank heat capacities shows that there is no low-temperature contribution due to conformational motion below the glass transition. The glass transition of the semicrystalline sample is broadened to higher temperature relative to the amorphous sample, as found in all polymers. Of
Fig. 6.52
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particular interest is the sizeable RAF that is attained on crystallization. It is about equal to the crystallinity and makes up one third of the total amorphous fraction. Up to 460 K neither reversing melting nor a contribution from a glass transition of the RAF is seen for the semicrystalline of PC. The heat capacity of PC represents in this temperature range the behavior of an ideal, linear macromolecule without contributions of local large-amplitude motion.
The slow, irreversible cold crystallization is followed in Fig. 6.53 for more than 10 days with quasi-isothermal TMDSC to a fixed value of RAF. At the end of the crystallization there is no frequency dependence of the heat-capacity. The crystallization and the glass transition to the RAF occur simultaneously (see also Fig. 6.18).
Fig. 6.53
The melting of the semicrystalline PC is shown in Fig. 6.54 [44]. The total cp, depicted in the lower curve is similar to a standard DSC trace. It shows two melting peaks, one between 460 and 485 K for the secondary crystals, and the main melting peak at about 503 K. Single melting peaks are seen for PC of this molar mass only on crystallization above 486 K, a temperature above which no fringed-micellar, secondary crystals grow [43]. A similar change to one melting peak is seen for PEN in Fig. 6.49 at 500 K and many other polymers (see Sect. 6.2.3). The excess of cp(total) beyond cp(liquid) at the main melting peak in Fig. 6.54 is about 20 times that shown by the reversing cp in the upper curve. In the low-temperature melting range, the reversing cp shows an increase beyond the baseline for the semicrystalline polymer starting at about 470 K. Quasi-isothermal measurements separate the contributions. They seem to indicate increasing devitrification of the RAF with time and may be interpreted as for poly(phenylene oxide), described next. As long as some RAF surrounds a crystal, it cannot melt. The final exotherm in TMDSC may be due to loss of stationarity as often seen at a sharp end of irreversible melting (see Sect. 4.4.3).
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Fig. 6.54
Overall, there seems to be very little or no reversible melting for PC, a result which is reasonable, considering that the induction time for crystallization of an oligomer of a molar mass of 4,000 Da, i.e., with about 14 repeating units, is still about 200 min at 483 K [43] and may preclude sufficiently fast recrystallization of a once-melted chain segment. All excess reversing heat capacity, thus, seems to be accountable by irreversible reorganization which decays with time.
Poly(phenylene oxide), of these, most common is poly(oxy-2,6-dimethyl-1,4- phenylene), better called poly(oxy-1,4-m-xylylene), [O–C6H2(–CH3)2–]x, PPO™, General Electric trademark. When analyzing the heat capacity of a 30% semicrystalline PPO, a surprising result was obtained as represented by Fig. 6.55 [45]. Practically all of the noncrystalline material behaves as a rigid, amorphous fraction, RAF. In the center DSC trace of Fig. 6.55 there is no indication of a glass transition before melting. Annealing for increasing lengths of time at 496 and 502 K, which is above the exceptionally high Tg of amorphous PPO™, however, causes a slow development of the glass transition, as can be seen in the left and right traces, respectively. In addition, as expected, the melting peak moves to somewhat higher temperature due to crystal perfection, but unexpectedly, the crystallinity decreases sharply. On annealing at 505 K, about 15 K below (!) the melting peak temperature, the sample becomes almost completely amorphous within 65 min. It was concluded from these experiments that the glass transition of the RAF is above the melting temperature of the original crystals and hinders the melting at the zero-entropy- production temperature of, perhaps, 500 K. The high Tg of the RAF is most likely also the reason for the difficulty in crystallization of PPO from the melt. As soon as crystals start to grow, they are surrounded by networks of rigid-amorphous materials and further growth is not possible. Usually a plasticiser or solvent must be added to crystallize PPO from the melt.
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Fig. 6.55
Figure 6.56 shows quasi-isothermal TMDSC traces of an amorphous and a 30% crystalline PPO™. The glass transitions of the two phases are clearly separated. The Tg(amorphous) occurs at about 488 K, while the broadened glass transition of the RAF has a midpoint of about 502 K. Inspection of the semicrystalline trace shows no reversible melting, supporting the argument that polymer crystals surrounded by RAF show little or no melting, neither reversible nor irreversible.
Fig. 6.56
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With the known crystallinity from standard DSC, it is easy to calculate the expected Cp in Fig. 6.56 (+). The deficiency of the calculated Cp from the measured Cp is marked as the length a. The length b is then due to the RAF which has already gone through its glass transition and is now a mobile-amorphous fraction. The ratio a/ Cpo is the remaining rigid RAF at the given temperature. It is plotted in Fig. 6.57 together with the measured, reversing Cp and the crystallinity from a separate set of experiments. As the temperature increases, the crystallinity and the RAF initially decrease at different rates. Melting is completed at about 510 K. Up to about 495 K the crystallinity decreases very little, while the RAF loses almost 20% of its value, which is in accord with the assumption that the surrounding glass must become mobile first, before melting can occur. Between 495 and 510 K the decrease of both, the RAF and the crystallinity, is close to linear, with the RAF losing three times more solid
Fig. 6.57
fraction than the crystallinity. In this temperature range, the crystallinity is lost parallel to the loss of the RAF, proving that the glass transition of the RAF is coupled to the melting. This result can be compared to the parallel development of crystallinity and RAF in Figs. 6.18 and 6.53.
Poly( -caprolactone), PCL, is an aliphatic, linear polyester with the repeating unit O (CH2 )5CO . Figure 6.58 illustrates the change of the real part of the shear modulus of PCL along with its amplitude on temperature modulation when using TMDMA [42,46,47] (see also Fig. 6.51). The sample was cooled from the melt at 1.0 K min 1 to the crystallization temperature and reached a crystallinity of 50%. The insert shows that during the crystallization the modulation of the temperature produces a reversible crystallization and melting. As the primary crystallization is complete at about 2.0×105 s ( 56 h), the increase in G' slows, but indicates considerable secondary crystallization with some decrease in reversibility.
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Fig. 6.58
Figure 6.59 displays the corresponding quasi-isothermal TMDSC at slightly lower temperature. The expected reversing specific heat capacity was estimated from the latent heat and the ATHAS Data Bank Cp for the crystalline PCL and the melt. The measured value, although lower than the cp of the melt, never reaches this level. The small decrease in reversing cp goes parallel to the small decrease of the amplitude of the storage modulus in Fig.6.58 in the region of secondary crystallization.
Fig. 6.59
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Next, it needs to be determined whether this excess heat capacity is indicative of conformational defects or true reversible melting. To make sure that linearity was maintained in the melting range for measurement of the reversing heat capacity, its amplitude-dependence was tested by quasi-isothermal TMDSC and is shown in Fig. 4.115. The dependence of the reversing heat capacity on frequency is represented in Fig. 6.60 over a wide frequency range in terms of excess heat capacity [48]. The drawn line suggests that a sigmoidal decrease can represent the frequency dependence. For PCL it should take a frequency of more than 10 Hz to separate the reversible melting from the base heat capacity and determine how much conformational contribution is present in the measured baseline. Assuming a similar concentration of gauche-trans conformations as in polyethylene crystals at the same temperature (see Fig. 2.65), the defect contribution may be within the error limit of measurement. The equilibrium melting temperature, Tmo, of 342 K for PCL is below Tmo of polyethylene, so that one would expect fewer defects and even less than in nylon 6 and 12 with values of Tmo of 533 and 500 K, respectively.
Fig. 6.60
Figure 6.61 describes the TMDSC of reversing cp on cooling from the melt, followed by heating. Except for the small sharp endotherm at the beginning of crystallization and exotherm at the end of melting, which may be caused by errors in the deconvolution from the irreversible latent heats, the two reversing heat capacities differ little. The quasi-isothermal points give an assessment of the slow reorganization processes. Below about 300 K this reorganization seems to stop, but an excess heat capacity can be extrapolated down to about 270 K. The PCL is thus a semicrystalline polymer which has no rigid amorphous fraction. It also seems to have no significant contribution of conformational disorder to the heat capacity, so that the total apparent reversible cp should be due to reversible melting. Another, related flexible linear polymer, poly-p-dioxanone, is illustrated in Fig. 4.124. It seems similar in behavior.
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Fig. 6.61
Polypropylenes, (CH2 CHCH3 )x, PP, is commonly found as its isotactic stereoisomer, iPP. Figure 6.62 illustrates a standard DSC analysis on cooling, heating, and annealing which was subsequently studied quantitatively by TMDSC [49]. The melting and crystallization peaks in the apparent heat capacity plots on the left side of the figure show a typical supercooling of about 40 K. The first heating trace represents a sample that was cooled quickly. In this case, iPP forms a metastable,
Fig. 6.62
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conformationally-disordered glass (condis glass) with a smectic structure, disordered mainly by helix reversals and a crystallinity of about 39% (see Fig. 5.146). On heating, the condis glass transforms after a small endotherm at 320 K with a broad exotherm centered at 367 K to the stable, -monoclinic crystal structure (see Fig. 5.46). Cooling from the melt at 10 K min 1 leads to a 53% -monoclinic crystallinity of lamellar morphology with a similar fold-length. An RAF of as much as 30% was observed on a different sample, cooled at 0.5 K min 1. The RAF devitrified between 325 and 370 K in a broad, second step of the glass transition and led then, after quenching, to a normal glass transition on reheating, followed by cold crystallization of about 6% of the polymer and remelting with an annealing peak at about 340 K [50]. The second heating trace in Fig. 6.62 does not have enough resolution to identify the RAF, but it can be seen in Fig. 6.64, below.
The morphology and fold-length were determined on identical and similar samples as used in the calorimetry with wideand small-angle X-ray diffraction, as well as atomic force and electron microscopy, WAXD, SAXS, AFM, and electron microscopy [49,51]. The similarity in fold-length accounts for the comparable main melting peaks seen in the left heating traces in Fig. 6.62. Next, both samples were step-wise annealed for 40–400 min, as shown in the right DSC traces of Fig. 6.62. After each step, an annealing peak (endotherm) is observed on resuming the heating which is much smaller for the melt-cooled sample (see insert). Only few kelvins later, the apparent heat capacity shown in the left curves is resumed, which suggests that only the crystals that lose their stability in the vicinity of the annealing temperature are affected by the annealing. In Fig. 6.63, the heating and annealing of the quenched iPP is followed by time-resolved X-ray diffraction. The upper figure shows changes on heating after annealing at 373 K. First, the diffraction intensity marked by ‘1’ is constant, followed by an increase in crystallinity at ‘2’ which corresponds to the
Fig. 6.63
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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 disappearance 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,
Fig. 6.64
but cannot be separated fully from the reversing melting. The curves on the right in Fig. 6.64 represent the latent-heat contributions of the apparent, reversing cp, obtained by subtracting the expected thermodynamic cp 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 thickness in the chain direction and the difference in lateral extension places the reversible melting for iPP on the growth faces of the crystals.