Thermal Analysis of Polymeric Materials Part 12 pps

More qualitative data were generatedthroughout the melting peak and show a maximum in reversible melting on heatingafter fast cooling of 3.05 J K1mol1and after slow cooling of 2.80 J K1m

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, slightincrease 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 theleft graphs of Fig 6.64 A comparison of the two reversing heat capacities shows thatthe cold crystallization and the transition mesophase-to--monoclinic crystals do notshow, i.e., they are nonreversing There remains, however, a substantial reversingcontribution which is larger for the quenched iPP than for the lamellar, melt-coolediPP The upper limit of the devitrification of the RAF seems to occur at 320330 K,

but cannot be separated fully from the reversing melting The curves on the right inFig 6.64 represent the latent-heat contributions of the apparent, reversing cp, obtained

by subtracting the expected thermodynamic cp for the given crystallinity from thecurves on the left Results from quasi-isothermal experiments on heating are plottedalso in Fig 6.64 These are taken after the slow crystal perfection ceased andrepresent reversible latent heats The slowly-cooled, lamellar sample begins to showreversible melting at 320 K, while the fast-cooled sample with globular morphologybegins melting at the end of the glass transition More qualitative data were generatedthroughout the melting peak and show a maximum in reversible melting on heatingafter fast cooling of 3.05 J K1mol1and after slow cooling of 2.80 J K1mol1[52].The same thickness in the chain direction and the difference in lateral extension placesthe 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 lateralgrowth face with an analysis of iPP samples annealed for long times at differenttemperatures The standard DSC traces were taken after the quasi-isothermal TMDSCcould detect no further decrease in the apparent reversing heat capacity and were usedfor a measurement of the remaining crystallinity Then, corresponding X-raydiffraction differences were taken in a temperature range of ±1.5 K after themetastable, global equilibrium was achieved in this temperature region On the rightside, Fig 6.65 displays the normalized difference patterns of the X-ray diffractionexperiments Within experimental error, the increase in latent heat, marked inFig 6.64 as open squares, and the increased differential diffraction areas in Fig 6.65give the same changes in crystallinity [53], making it certain that the reversiblemelting refers to the same process seen in the irreversible melting and crystallization

Syndiotactic polypropylene, sPP, is known almost as long as iPP, but only becamereadily available with high stereospecificity after the discovery of soluble, single-site,metallocene catalysts (see Sect 3.2.1) Several semiquantitative studies are availableand have been reviewed in Ref [1] Figure 6.66 illustrates TMDSC traces of an sPPsample of low crystallinity (and stereospecifity) The reversing cp of the semi-crystalline state on initial cooling and on subsequent heating are identical Above thelow-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 theRAF is indicated between 315 and 365 K Above this glass transition, reversingmelting is observed with both of the irreversible melting peaks showing a reversingcomponent The frequency dependence of the reversing heat capacity of sPP afterisothermal crystallization at 363 K is shown in Fig 6.67 [54] At low frequency thesample shows reversing melting which at higher frequency reverts to a partially

Fig 6.66

Fig 6.67

devitrified RAF Both iPP and sPP are thus examples with a broad devitrification ofthe RAF before the major melting peak The frequency dependence of the reversingspecific heat capacity of sPP over the full temperature range is displayed in Fig 6.68and should be compared to similar results on other polymers (see, for example, theresults for nylon 6 in Figs 4.116 and 4.119) The frequency dependence was extractedfrom 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 theirproperties in Table 6.1 [1,33] To restate the main facts, it is shown in Figs 3.75 and3.91 that flexible molecules melt reversibly only up to a critical length, which forparaffins is 10 nm Longer extended-chain crystals melt irreversibly (Figs 3.89 and6.28) A first observation of decoupling of melting concerns the limiting reversibility

of crystals grown from a distribution of molecules of low molar mass In this case twophysical changes that should occur simultaneously are partially decoupled Themelting/mixing and crystallization/demixing are incomplete, probably due to slowdiffusion in the available time This can cause the limited reversibility seen inFigs 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 ofsemicrystalline, flexible macromolecules also exhibit decoupling Such decouplingallows molecular nucleation to start with crystallization of segments of the polymerchains (see Fig 3.74) The portions of the molecules not crystallized must beconsidered as amorphous and belonging to a different phase Figure 6.69 illustratessuch 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 continuingmolecule, this picture also offers an explanation of the broadening of the glasstransition and the creation of a separate RAF In case Tgof the RAF is close to Tm,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 fromthe known Tmof 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

Table 6.1 The reversing behavior of the various polymersa

Polymer cold

cryst

hotcryst

mational Cp

confor-RAF(b)

reversiblemelting

reversingmelting

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 meltingtemperature into the upper equation instead of 414.2 K (the melting temperature of theextended-chain crystal of infinite length) and correcting the melting point lowering (inbrackets) for the fraction of the surface covered with chain folds, yields the showncurves (For one fold, the correction factor for the term in brackets is 0.50, for twofolds, 0.67, for three, 0.75, for four, 0.80, and for five, 0.83.)

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 conformationallydisordered The conformationally disordered crystals are mesophases, described inSect 5.5 The reversing and reversible melting of the mesophases depends on themobility within the crystal, as seen from Figs 5.131, 5.144, and 5.154 Liquid crystalsare 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 arecloser in their behavior to the small polymer crystals with their slow crystal perfectiondue 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), tosmall crystallites, dendrites, and spherulites (see Figs 5.56, 5.60, and 5.75), or tomesophase 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 aspossible 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 anetwork of larger crystals, separated by amorphous defects, as described by Fig 5.87

On cooling, these amorphous defects may generate a second population of poorcrystals, 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 reversiblemelting and crystallization One expects these poorly grown crystals to be close to afringed-micellar structure, schematically shown in Fig 5.42, with many of themolecular chains being decoupled into crystals and amorphous segments

A unique method to generate poor crystals is shown in Fig 6.71 It involves thecompression of crystals of poly(4-methyl-1-pentene), P4MP1, with a glass transition

Tg= 303 K, melting transition Tmo

= 523 K, and a heat of fusion of Hf= 10 kJ mol1(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 helicalconformation 2*7/2 for P4MP1 is the reason for the poor packing in the crystals atroom temperature, compared to the amorphous phase Increasing the temperatureexpands the amorphous phase faster than the crystals, so that at Tmthe crystals aresomewhat denser The change of Tmas a function of pressure is shown in Figs 6.72and 6.73 After the initial, expected increase with pressure, Tmdecreases

Fig 6.74

This unique behavior of the melting temperature with pressure indicates thatincreasing the pressure reverses the density difference in a similar manner asdecreasing the temperature At the maximum of the Tmversus 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 andwritten in Fig 5.168 In Fig 6.73 a phase diagram for the various states is drawnbased on high-pressure calorimetry Only the melt appears as a true equilibriumphase All other phase areas are semicrystalline, i.e., do not follow the phase rule (seeSect 2.5.7) Because Fig 6.73 does not represent equilibrium, it can mirror the actualprocesses All exotherms and endotherms, except for melting, are small (0.5 to 2 J/g),indicating that at the transitions, the changes in entropy of the phases are small Majorcrystallization and melting, thus, is not possible

The nonequilibrium states were identified and seem to be initiated in close concertwith the existence of the glass transition of the majority glassy phase identified inFig 6.73 Isothermal disordering at room temperature along the horizontal arrowcauses 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 superficiallyshowing an amorphous X-ray pattern in Fig 6.71, the condis glass retains most of itsheat of fusion and acquires largely sessile backbone and side-chain conformationaldefects This disordering can be reversed by reduction of pressure Based onFig 6.73 one can speculate that the equilibrium phase-diagram is as given in Fig 6.74.The tetragonal phase is the stable crystal The trigonal crystal form with a 2*3/1 helixconformation, common for vinyl polymers with smaller side chains (see Fig 5.14),exists as the low-temperature, high-pressure phase As expected from a high-pressure,low-temperature phase, it is denser than the trigonal crystals Neither of the twoequilibrium 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 isinvolved in the process and renewed crystallization occurs before the whole sampleturns liquid, the process is called recrystallization rather than annealing Naturally,both processes may occur at the same time in a semicrystalline sample Glassypolymers are mainly annealed close to the glass transition for strain release anddensification (see Sect 6.3), while crystalline polymers perfect their internal structureand their morphology on annealing (see Sects 5.1–3) Crystal structure andmorphology are best assessed by X-ray diffraction and electron microscopy, while thechange in macroscopic properties is, as usual, best tested by thermal analysis.Point defects and dislocations are shown in Sect 5.3 to appear as nonequilibriumdefects, usually introduced during crystal growth or deformation, and as equilibriumdefects, generated thermally The dislocation density can be determined by the moirémethod (see Appendix 17 and Fig 5.93) On annealing of crystals above theircrystallization temperature, the nonequilibrium dislocation density increases, as isdemonstrated 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 compensatestheir metastability, one assumes that the cause is a change in the packing within thecrystals or between the lamellae that cause the moiré pattern The right graph inFig 6.75 illustrates, based on X-ray diffraction, that the mosaic dimensions, indeed,

Fig 6.76

increase on annealing in the crystallographic a and b directions (see Figs 5.18 and5.19) A fast, but small adjustment from 12.5 to 20 nm in fold length occurs withinthe first minute along [001] (compare to the initial dimension, and see also Fig 6.34).The following long-time annealing affects all dimensions of the crystals

On a larger scale, crystals of polymers were shown in Sect 5.2 to commonly have

a fibrillar or lamellar morphology Neither of these is an equilibrium shape Figure6.76 illustrates how these nonequilibrium morphologies would have to change to reach

an equilibrium shape More detailed studies are available for the thickening of thelamellar crystals It involves an increase of fold-length by a mechanism as displayed

by the right-hand sketch in Fig 6.76 The intermediate states have only a small

increase in free enthalpy which acts as an activation energy for the process to the morestable state Actual observation of the thickening is seen in Fig 6.77 Separatedlamellae develop holes with rims of about double the original thickness From thedimension of the rims, it can be estimated that for each growth plane that crosses thehole, only one molecule thickens in the early stages of annealing The graph inFig 6.77 indicates that annealing speeds up above 395 K The crystals become moreirregular and have more than twice the fold length A hole in one lamella is naturally

an ideal position for the adjacent lamella to thicken into

The electron micrograph of Fig 6.78 depicts melt-grown polyethylene lamellae ofthicknesses of 1822.5 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 onesee the growth face of the folded-chain lamellae In the bottom of the figure thelamellae were annealed below the melting temperature of the substrate, at 411.6 K, for

18 h The lamellae are now irregular and show thicknesses of 27 to 45 nm Theunderlying extended chain crystals are little affected by the annealing

6.2 Size, Extension, and Time Effects During Fusion _657

Fig 6.78

Fig 6.77

This short summary of microscopic observations of annealing reveals that there are

at least two processes that must be considered when interpreting the macroscopicannealing experiments of thermal analysis: (1) A crystal perfection within the originalcrystal morphology (2) A change in crystal shape towards equilibrium, which mayinvolve several stages of crystal thickening

Fig 6.79

The macroscopic thermal analysis, furthermore, proves that it is also common formelting of the original crystals to be followed by recrystallization to a more stablemorphology Finally, some crystals may have different crystal structures at hightemperatures and, thus, show polymorphic transitions on annealing Annealing ofpolymers to a more stable state as shown in Fig 6.79, simply expressed by a decrease

in free enthalpy, may thus be a rather complicated process (compare to Sect 2.4.2)

On discussing the melting of polyethylene in Sect 6.2.1, the annealing of samples

on heating was observed in curves A and B of Fig 6.22 to cause a decrease in meltingtemperature with heating rate In case the kinetics of rearrangement of the crystalscould be outrun, a constant zero-entropy-production melting temperature wasobtained At the same time, the slower heating rate experiments give someinformation about the annealing process Besides using fast heating to minimizereorganization, it is possible to stop the thickening of crystal lamellae by immobilizingthe amorphous phase Two examples are shown in Fig 6.80, one for nylon 6, theother for polyethylene For nylon 6 theNH-groups of the amorphous nylon (seeFig 1.18) are methylmethoxylated toNCH2OCH3, so that they cannot be movedinto the crystal Figure 6.80 implies that after 4 h, the single, heating-rate independ-ent, sharp melting peak is related to the zero-entropy-production melting On longerreaction, this peak broadens because of reaction with the crystals In addition, themelting point seen after 4 h has also increased by about 3 K due to the permanentchange of the molecules, or possibly due to a change of the initially present-crystals

to the higher melting -crystals (see also Sect 6.2.6) For immobilizing theamorphous polyethylene, cross-linking by irradiation has been used, as also shown inFig 6.80 The irradiation caused a similar effect as fast heating shown in curves Aand B in Fig 6.22

6.2 Size, Extension, and Time Effects During Fusion _659

Fig 6.81

Fig 6.80

In the following figures, further experiments are linked to the annealing andrecrystallization of a number of additional polymers In Fig 6.81 the increase oflamellar thickness is illustrated on the example of solution-grown polyethylene Thecrystals were collected in dried mats The graph on the right side allows a comparisonwith data on melt-crystallized and annealed polyethylene (see also Sect 5.2).Sufficiently mobile and flexible molecules, such as polyethylene, can thicken by chain

Fig 6.82

extension, either on crystallization, or on subsequent annealing at higher temperatures.For polyethylene, long-time annealing at temperatures above about 400 K can reachthicknesses of over 100 nm Full extension of the polyethylene crystals, however,needs the intervention of the hexagonal condis phase (see Fig 5.156)

Of the polyamides, nylon 6 is perhaps best studied Figure 6.82 shows thesimilarity of the melting of solution-grown  crystals of nylon 6 lamellae withpolyethylene in Fig 6.22 The initial crystal thickness was only 56 nm, accountingfor the larger drop in melting temperature The melt-crystallized nylon 6 is initiallyabout 10 nm in fold length and leads to the correspondingly higher zero-entropy-production melting temperature On annealing, the perfection gained during slow

heating leads on faster heating to a small amount of superheating The curves inFig 6.82 show only the temperature of the last (major) melting peak As many as fourlower-temperature melting peaks can be observed, as are illustrated in the summary

in Fig 6.83 The form has a somewhat higher melting temperature () than the form () The horizontal appearance of lines 1 and 2 are an indication of reorganiza-tion of the initially less-perfect crystals, as in Fig 6.82 for the crystals Lines 3 and

4 are closer to zero-entropy-production melting Line 5, finally, refers to the

“annealing peak.” It corresponds to small crystals that form between prior growncrystals They melt somewhat above the crystallization or annealing temperature Asseen in Fig 8.62, most melting temperatures shift when changing the heating rate Onreorganization, the peak sizes will also change Sometimes, an attempt is made toextrapolate the assumed zero-entropy-production melting points (lines 3 and 4) to thepoint of intersection with the line that marks the position of Tm= Tcto estimate theequilibrium melting temperature, Tmo

This is only safe on using the production melting temperatures, on avoiding superheating of crystals grown at higher

zero-entropy-6.2 Size, Extension, and Time Effects During Fusion _661

6.2 Size, Extension, and Time Effects During Fusion _663

Fig 6.87

Fig 6.88

To illustrate recrystallization as a frequent occurrence in addition to reorganizationand lamellar thickening, the behavior of melt-grown polyethylene crystals andpoly(ethylene terephthalate) are chosen as examples in Figs 6.88 and 6.89 The freeenthalpy diagram for such process is given in Fig 6.79 The crystals melt partially orcompletely when reaching the annealing temperature, indicated for polyethylene inFig 6.88 by the increase in specific volume and decrease in density This is followed

Of interest is the larger rates ofrecrystallization when compared to melt crystallization at the same temperature in theright graph of Fig 6.88.

The amount of recrystallization for poly(ethylene terephthalate), PET, in Fig 6.89increases with annealing temperature, as in polyethylene At 520 K, the thin film ofPET could be melted almost completely by fast heating Thin films were chosen forquick temperature equilibration at constant temperature The analysis of crystallinityafter 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, 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 viscousflow at the high temperature and by dilatometry and DSC after quick quenchingbetween cold plates In the same research isothermal primary melt crystallization wasshown to be followed by insertion of secondary lamellar stacks and continuous crystalperfection by time-resolved small-angle X-ray scattering and DSC at different heatingrates The major high-melting peak at 525 K was shown not to be due to the originallygrown crystals, but was changed due to recrystallization and annealing

6.2 Size, Extension, and Time Effects During Fusion _665

Fig 6.90

Fig 6.91

More quantitative information than by calorimetry in Fig 6.89 could be gainedrecently by the use of ultra-fast, thin-film calorimeters based on integrated circuits (seeAppendix 10) [55] Figure 6.90 illustrates the faster recrystallization than crystalliza-tion from the random melt, as was also seen for polyethylene in Fig 6.88

A special effect in changing the melting characteristic is transesterification,illustrated in Fig 6.91 Transesterification is found in polyesters, as described in

Fig 6.92

Sects 3.1 and 3.4 and involves the possibility to change the backbone chain structure

to accommodate better crystallization Figure 6.91 contains a collection of annealingstudies by DSC, analyzed as a function of the heating rate The dotted lines refer tostandard crystals grown from solution, from the melt (isothermally grown at 523 K,sample S1), and obtained after annealing of sample S1at 533 K The three curvesillustrate decreasing to increasing melting peaks with heating rates, analogous to thepolyethylene and nylon 6 data in Figs 6.22 and 6.82, respectively In addition, thesample S1was etched with superheated steam at 453 K to remove the amorphoussegments to increasing amounts The lamellar morphology of the remainingcrystalline PET is seen in Fig 5.65, and the thermal properties are illustrated inFig 6.91 Depending on the degree of etching and the remaining amount ofamorphous, low-molar-mass debris, the zero-entropy-production melting temperaturedecreased Annealing these etched samples R, Q, Re, and Qe, identified by differentsymbols, lead to repolymerization to similar high-molar-mass crystals by transes-terification without intermediate melting, as seen by the upper lines A similarindustrial process is known as solid-state polymerization and used for tire-yarnimprovement Transesterification, thus, is important in annealing of PET crystals

6.2.4 Melting of Poly(oxymethylene)

Poly(oxymethylene), POM, is a polymer with well-known melting behavior [56].Fibrous, extended chain crystals can be produced by crystallization during polymeriza-tion, as described in Figs 3.104 and 3.105 These extended-chain crystals cansuperheat on melting due to their slow melting rate, as is also seen for extended-chaincrystals of polyethylene in Fig 6.21 The DSC melting endotherms in Fig 6.92 startfor all heating rates at about the same temperatures, but the peak- and melt-end

6.2 Size, Extension, and Time Effects During Fusion _667

Fig 6.93

Fig 6.94

temperatures increase, indicating a kinetic effect The plot of temperature versus rate of heating is shown in graph B and reveals less superheatingfor imperfect crystals The estimated equilibrium melting temperature is 457 K.The next thermal analyses are for solution-grown dendrites, hedrites, and melt-grown lamellae, represented by Figs 6.93 and 6.94 The dendrites melt at much lowertemperatures than the extended-chain crystals and show multiple melting peaks

melting-peak-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 crystalsoccurs 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 tworecrystallizations during the analysis

Curve E in Fig 6.94 shows the minimal change in melting peak temperatures forcrystals grown from the melt These crystals do not reorganize significantly onheating They remain metastable during the whole analysis Their zero-entropy-production Tmis 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 crystalmorphology 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 actualDSC traces Their higher perfection is indicated by the higher Tmfor 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 ing to lamellar melting with major reorganization bracketing equilibrium and differentzero-entropy-production melting temperatures

superheat-6.2.5 Melting of PEEK

Poly(ether ether ketone), PEEK, 1,4-phenylene), (OC6H4OC6H4COC6H4)x, can be considered to be made up

poly(oxy-1,4-phenyleneoxy-1,4-phenylenecarbonyl-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 rigidsegments allow rotations about their bonds and induce flexibility In addition, thephenylene groups may flip about their twofold symmetry axis As a high-temperaturepolymer, PEEK is often used as a matrix in composites The overall behavior ofPEEK [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 andmelting peaks have close to the same size A more precise analysis would have to takeinto 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= Tccorresponds to zero-entropy production for tion and melting, and was extrapolated to the experimental line of data to give anestimate of the equilibrium, Tmo, of about 668 K (compare with other polymers inFigs 6.83–87, but see also the mentioned cautions in the discussion of Tmo

6.2 Size, Extension, and Time Effects During Fusion _669

Fig 6.97

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

A more detailed analysis involves DSC, TMDSC, quasi-isothermal TMDSC, andTMDMA [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 thereversing cpfrom TMDSC at the same heating rate, and quasi-isothermal TMDSC atshort and long times after heating to To All measured apparent, specific heat

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, thereversing endotherm of melting is twice as high, and because of the identical heatingrates 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 Theanalysis 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 alower heat capacity of a semicrystalline sample The crystallinity is derived from theintegral of the total heat-flow rate, <0(t)>, with time in Fig 6.98 From Fig 6.99 itcan 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

6.2 Size, Extension, and Time Effects During Fusion _671

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 reversingamplitude of the storage modulus is almost constant At later times, it decreasesslowly, indicating very different processes for primary and secondary crystallization

6.2.6 Melting of Fibers

Specially difficult systems are represented by fibers of macromolecules All the topicstreated in the prior sections must be considered under the additional aspect of thepresence of crystal deformations caused by the drawing process (see Sect 5.2.6 and5.3.6) and possible strain retained in the amorphous areas (see Sect 6.3.3), as well asthe 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 inFig 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 drawnsample 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 thesame In fact, extensive experimentation with fibers of known crystal sizes whichwere either cross-linked to avoid reorganization or etched to remove the amorphousfraction, suggests that the outcome of both experiments with the drawn fibers inFig 6.101 is based on the increase of the melting temperature due to residual

6.2 Size, Extension, and Time Effects During Fusion _673

Fig 6.101

orientation in the amorphous fraction (see Sect 5.6) The sharper melting peak of theunrestrained 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 thesuperheating to higher temperature (see Fig 2.120)

The special behavior of fibers is particularly well investigated on gel-spun, high-molar-mass polyethylene, UHMMPE The as-polymerized UHMMPE shows amuch higher amount of reversing melting than the lower-molar-mass folded-chaincrystals of PE, described in Sect 6.2.1, or the extended chain crystals of PE and itsoligomers of Fig 6.28 Figure 6.102 displays DSC and TMDSC traces of nascentUHMMPE and illustrates the changes on recrystallization from the melt and on secondheating [60] The apparent reversing heat capacity exceeds the standard DSC result

ultra-on the low-temperature side of the melting peak On crystallizatiultra-on from the melt, thereversing latent heat is less than on first heating, but increases on reheating, as wasalso seen for the lower-molar-mass PE in Fig 6.30, just that the excess Cpon coolingand second heating is higher than for the lower-molar-mass PE, despite the likelylower crystallinity for the UHMMPE The frequency dependence of the nascentUHMMPE was also analyzed It showed a decrease in excess heat capacity reachingthe 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 K1 mol1(compare to Fig 6.31) Model calculation of the frequency dependence of thereversing excess Cpallowed in this temperature range to speculate about a number ofdifferent irreversible and reversible processes Quantitative information may beavailable as soon as the morphology is known for the specific sample

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

Fig 6.102

Fig 6.103

Details on calorimetry coupled with information on morphology, full-pattern X-rayanalysis (see Fig 6.4), solid-state NMR (see Figs 5.157 and 5.158), mechanicalproperties, and quantitative TMDSC were derived for a number of commercial gel-spun UHMMPE fibers The structure of an UHMMPE gel-spun fiber is rathercomplicated and changes on heating with constraint, as seen in Fig 6.103 It consists

of a metastable system of four interconnected, recognizably different phases At low

6.2 Size, Extension, and Time Effects During Fusion _675

Fig 6.104

temperature, the orthorhombic PE is the major and only stable phase (typically60–80%) Also present are some monoclinic crystals of PE (5–15%) The amorphousphase is as little as 1–10%, while an intermediate, oriented phase makes up thedifference The three unstable phases are arrested in a metastable state by theirmolecular links to the orthorhombic crystals The X-ray signal of the monocliniccrystals starts to disappear at 363 K On heating of the fibers so that they are free toshrink, the whole structure collapses when the orthorhombic crystals melt at about

415 K The intermediate, oriented phase was shown by solid-state13

C NMR to consist

mainly of trans conformations which are at room temperature intermediate in mobility

between the liquid and the crystal as illustrated in Figs 5.157 and 5.158 A solid-state13

C NMR analysis as a function of temperature gave for typical, gel-crystallizedUHMMPE fibers that at room temperature 7% monoclinic phase which disappeared

at about 380 K Its 78% orthorhombic crystals and about 15% intermediate phaseincrease somewhat to make up the loss in monoclinic phase at 380 K The ortho-rhombic phase showed melting from 393 K (75%) to 413 K (72%) while theintermediate phase still increased in amount (25% at 393 K and 28% at 413 K) [61].The amorphous phase remained negligible throughout the experiment at1%.The calorimetry of the UHMM fibers is beset by an additional difficulty due to theconsiderably lateral expansion on shrinking of the fibers on melting as is illustrated

in Fig 6.104 In case the fiber shrinks from an extension of 100× to its originallength, the diameter needs to expand by a factor of 10 and a lateral constraint isintroduced when the lid of the DSC pan was crimped onto the fibers which issufficient to hinder longitudinal shrinkage and may even be sufficient to deform theDSC pan and invalidate the calibration of the run [62] In the presence of suchconstraints, the orthorhombic crystals superheat sufficiently and undergo the phasetransition to the hexagonal mesophase, starting at about 415 K, as seen from the X-ray

The electron micrograph of Fig 6.78 depicts melt-grown polyethylene lamellae ofthicknesses of 1822.5 nm, grown at 403 K on the surface of the (110) growth face

of an extended-chain crystal,... summary of microscopic observations of annealing reveals that there are

at least two processes that must be considered when interpreting the macroscopicannealing experiments of thermal analysis: ... 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,