Petr

Physical and Microstructural Effects of Heat Setting in Polyester Films J. GREENER,'A. H. TSOU,2 andT. N. BLANTON Eastman Kodak Company Rochester, New York 14650-2158 The effects of a controlled heat-setting treatment on the properties and microstructure of biaxially stretched poly(ethy1eneterephthalate) (PET) and poly(ethy1ene-2.6-naphthalene dicarboxylate) (PEN) films are described. Substantial changes in crystalline fraction, crystallite size, glass transition temperature and significant enhancement in dimensional stability are observed for both film types upon increase in heat-set temperature. A distinct melting peak in the vicinity of the heatset temperature, observed in differential scanning calorimetry thermograms for both materials, is shown to be a n effective marker of the heat-setting process underlying the dual nature of the morphology of the heat-treated films. We also observe that the PET films undergo significant molecular realignment on heat setting, while the orientation of the PEN films is only weakly modified by the heatset conditions. A morphological transition is detected for both films at high heat-set temperatures (THs*) near the onset of the primary melting range, marking a qualitative change in the physical response of the biaxially oriented films; Measurable drop in planar orientation and glass transition temperature, and a sharp rise in crystallite size are noted for films heat-treated above THs*.The morphological models of previous workers (Schultz et al. and Fischer and Fakirov), invoking the idea of fibrillar-to-lamellar transformation, are used to explain some of our observations. 1. INTRODUCTION

H

eat setting, or constrained high-temperature annealing, is a n important step in the process for making biaxially oriented polymer films. The essential microstructural features of the f i m are established during the sequential stretching steps as the material undergoes molecular alignment and strain-induced crystallization. However, because of the low temperature and rapid crystallization in these steps (the kinetics of strain-induced crystallization is ordersof-magnitude faster than that of spontaneous thermal crystallization (1, 2)). the crystalline domains formed during stretching contain many lattice and conformational defects. These defects can be partially eliminated and replaced with a more dense crystalline structure by suitable annealing at high temperatures. In turn, the buildup and "perfection" of the crystalline domains and the attendant changes in morphology and orientation can significantly alter the physical performance and uniformity of the stretched film. The importance of heat setting as a means for optimizing the structure and performance of oriented films -

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'Correspondmg author 2hesent address Exxon Chemical Company. Baytown. TX 77520

and fibers has been well recognized. The morphological changes underlying the heat-setting process -in poly(ethy1eneterephthalate) (PET) and other semicrystalline polymers have been studied extensively over the past three decades by means of X-ray diffraction, IR spectroscopy, calorimetry, rheology, and other characterization techniques (2-1 7). Most studies conclude that the increase in the Crystalline fraction and the size of the crystalline domains during heat setting sign@ a qualitative change in the superstructure and morphology of the oriented matrix, but the exact Mture of this microstructural change has been the subject of some debate in the open literature. In studies on biaxially stretched PET films and uniaxially drawn fibers, Schultz and coworkers (15, 16) contend, in line with earlier observations by Fischer, Fakirov and coworkers (4-6).Gupta et aL (10, 11)and others (9), that the heat setting process not only reduces the concentration of lattice defects but it also leads to the formation of lamellar crystals extending normal to the chain (or drawing) axis. More recently, in a study of the effects of heat setting on the physical properties of biaxially stretched PET films, Gohil (17)has noted that the heat setting process can be divided into two temperature regimes marked by different crystallization

POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

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J. Greener, A. H . Tsou, and T. N . Blanton rates and Werent temperature dependencies of some key physical properties. In this study we examine the effects of a controlled heat setting treatment on the morphology and physical performance of two polyester films-poly(ethy1ene terephthalate) (PET) and poly(ethylene-2,6-naphthalene dicarboxylate) (PEN)-and assess the similarity and differences between these materials stemming from their Merent molecular structures. The microstructural changes induced by heat setting and the implications of these changes to some macroscopic properties of the stretched films are briefly discussed in Section 4.

G"n1 + n, + n3, B R = n, - n,, and 3

PBR =

2. EXPERXHENTAL

Materiala The materials used in this study are commercial resins of PET and PEN with inherent viscosities of 0.62 and 0.69 dl/g, respectively. Non-heatset (NHS) biaxially stretched PET and PEN films with nominal thicknesses of 178 and 84 p,m, respectively, were produced on a pilot film machine. The PET film was produced by stretching a cast sheet sequentially 3.5 X 3.5 at temperatures of 80 and 103°C. respectively, for the machine and transverse direction stretches, while the NHS PEN film was produced by stretching sequentially 3.6 X 3.6 at temperatures of 139 and 141°C. Heat satting Controlled heat-setting experiments were conducted using a thermostated oil bath. The NHS films were attached to a rigid rectangular metal frame with its long dimension aligned with the machine direction (MD) of the stretched film.The film was clamped tightly to the frame to ensure that the film boundaries are fully constrained during the heat treatment, to simulate the lateral constraints imposed in the machine during the film process. The clamped film was then immersed vertically in a silicone oil bath, which was maintained at a constant temperature ( 2 0.5OC). After a specified soak (anneal) time, the sample was removed from the bath and allowed to drip excess oil for 10 sec and was then rinsed with a household detergent and warm tap water at -55°C. The clean f i l m was removed from the frame, sectioned and characterized by various techniques.

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-w Density was measured by the flotation method using a density-gradient column with a mixture of tetrachloroethylene and toluene. All samples were measured twice and averaged. The accuracy of the measurement is estimated at ? 0.00025 g/cc. Birefringence A polarizing Abbe refractometer (B&S Model 60HR) was used to obtain the refractive indices of the test samples in three orthogonal directions, nl. n, and n3, 2404

where 1 is the machine direction (MD), 2 is the transverse direction ('I'D), and 3 is the thickness (normal) direction (ND).All measurements were conducted with a monochromatic Sodium D line (589.6 nm) source. Diiodomethane was used as the contact liquid for the PET samples while Gem oil (n = 1.81)was used with the PEN samples. The refractive indices were used to evaluate the average index and the in-plane and planar components of birefringence through

n, + n, - 2n3 2

where BR expresses the relative in-plane alignment of polymer chains in the machine vs transverse directions rbalance"), while PBR represents the relative alignment of polymer molecules parallel to the film plane ('stratification').

ThermomechanidAnolJ.ia A Seiko TMA/SS 100 (SeikoInstruments) was used to obtain information on the dimensional stability of the test samples. Dimensional change in the machine direction as a function of temperature was measured by scanning the sample (10 X 3 mm) from 20 to 240°C at a rate of 2"C/min. No tension was applied during the scan.

DJirPmic-llbechrmicrrli. The isochronous dynamic moduli and loss tangent of the test samples were measured using a Polymer Labs DMTA Mark I1 system. Samples, ca. 13 X 12 mm, were cut along the machine direction and tested in the oscillatory tensile mode by scanning from 0 to 200°C at a rate of 2"C/min and a frequency of 1 Hz.

DiHerentid8crrPnimgCalorimetry A DuPont 990Thermal Analyzer was used to evaluate the thermal response of the heat-set fims. The films were tested as received, i.e., without drymg or preconditioning, so as not to mod@ their thermodynamic state after the heat setting treatment. The DSC thermograms were generated by scanning at a rate of 10°C/min from ambient to ca. 300°C. Only 'first heat' scans were analyzed. Wide Angle X-Ray Difthction

X-ray diffraction was utilized to probe the microstructure of the crystalline phase, particularly the crystal orientation and average cxystallite size. Survey scans in reflection and symmetrical transmission modes were collected to obtain preliminaq crystallinity and orientation information, relative peak intensity ratios, and crystallite size values. These data were collected with a Rigaku RU-300 pole figure goniometer used in the Bragg-Brentano geometry. A detailed de-

POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

Effects of H e a t Setting in Polyester Films scription of the experimental procedure and data analysis is given elsewhere (18). In-plane and planar orientation functions of the crystalline phase were extracted from azimuth scans and complete pole figures. In-plane orientation is expressed in terms of the "azimuthal" Herman's Orientation Function, HOFa, while the planar orientation is expressed by the "tilt" Herman's Orientation Function, HOFc, where

3 (COS2P) HOFa =

2

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1

and H O E =

Rgures 1 and 2 for PET and PEN films, respectively. In both cases, the density is increasing asymptotically to a value that remains constant a long time (- 60 min) after saturation, but this asymptotic value is dependent on the applied heat-set temperature. The kinetics of the densification process appears to be independent of THS, but some differences in saturation time between the materials are noted. These differences are attributed to variations in film thickness, i.e., the kinetics of the heat-setting process at the experimental time scales is controlled, most likely, by thermal diffisivity. Indeed, in a study by Lee and Schultz (16),crystallization half times for the heat setting process in PET fibers were found to be of the order of 10-100 msec-much less than the annealing times covered in this study, i.e., the crystallization process may be considered to be instantaneous. In both cases, saturation is attained at times <20 sec, with the actual equilibration time depending primarily on the thickness of the sample.

3 (cos2Q) - 1 2

(2)

P is the azimuth angle and Q = 90" - a where a is the tilt angle. Also, (cos2p)and (cos2@)are the second moments of orientation distribution with respect to p and Q, respectively. 3.RESULTS The evolution of density with anneal (soak) time (tHS) at several heat-set temperatures (TH.Jis shown in

L /---------

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Fig. 1 . Evolution of density with Effect of anneal timefor PET-. heat-set temperature. Arrow Mic a s NHsfiln

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Time (sec) POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Yo/.39, No. 12

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J. Greener, A. H . Tsou, and T. N. Blanton

crystallinity. A correlation between X, and THS for PET and PEN is given in m e 4. We note that, over the wide range of temperatures covered, the change in crystalline fraction is substantial for both materials but especially for PET (25-+ 52% vs. 29 -+ 41% for PEN) and, within the scatter of the data, X, increases approximately linearly with THs. (Buchner et aL (21) have proposed a higher value for pa for PEN, 1.340 g ~ m - which ~, will tend to lower the estimated crystalline fraction by >30%, so the results listed above may be considered a 'high" estimate for PEN). This result indicates that the heat-setting process brings about a significant transformation in the material, the extent of which depends strictly on the heat-set temperature once thermal equilibration has been attained. Also, the "efficiency" of the process in terms of its ability to crystallize amorphous matter is somewhat lower for PEN than for PET. This may be due to lower inherent crystallizability of the PEN matrix or

Figure 3 shows the effect of THS on density at saturation ( t H s = 30 sec). A simple, monotone relationship is found for both materials, indicating a steady increase in the crystalline fraction with heat-set temperature, with PEN exhibiting lower densities overall and a lower rate of change of density with THS. Assuming a two-phase model, the corresponding changes in the crystalline volume fraction (&) can be estimated from

&-- P - Pa Pc

- Pa

(3)

where pc and pa are the densities of the crystalline and amorphous domains. The values of pa a n d pc are 1.335 and 1.455 g ~ m for - ~ PET (19) and 1.325 and 1.407 g ~ m for - ~ PEN (201,respectively. Although the use of Eq. 3 has been called into question on the grounds that a two-phase model is not strictly valid for highly oriented semicrystalline polymers (5), it does provide a reasonable estimate of the degree of

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Fig. 3. Density us. heat-set temperature for tHs = 30 sec. Arrows indicatemsjxms.

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POLYMER ENGINEERING AND SCIENCE, DECEMBER 7999, Vol. 39, No. 12

Effects of Heat Setting in Polyester Films

simply a result of some uncertainty in the values of the amorphous and crystalline densities of this material (see above comment). Differential scanning calorimetry of the heat-treated films also reveals some interesting features of the heat-setting process. Figures 5 and 6 show DSC scans of samples heat set at various temperatures for a fixed soak time of 30 sec. The thermograms for both materials feature a low-temperature 'secondary' melting peak (Tm')at temperatures that nearly coincide with the actual heat-set temperature used to prepare the sample. The enthalpy associated with the secondary endotherm is considerably smaller than that of the primary peak, but even if heat setting takes place within the primary melting range of the material, the secondary peak retains its identity and appears to simply superimpose on top of the primary melting endotherm; The two melting events, which can be

readily deconvoluted, represent, most likely, two distinct morphologies, one (Tm') intimately associated with the heat setting process and the other (T,.,) practically independent of the heat setting treatment. A plot of the secondary melt temperature vs. heat-set temperature, Figure 7, shows a near identity between these temperatures, suggesting that some of the crystalline domains formed during the heat-setting process are relatively unstable and thus can be melted out at temperatures only slightly higher than tbeir formation temperature. Studies on the dual-crystallization behavior of PET and other polymers were reported by Fakirov et aL (6),Holdsworth and TumerJones (22). and others (17, 23-26),all of whom detected secondary melting peaks at temperatures close to or above the annealing temperature. However, Fakirov et al. observed a dual-melting behavior, similar to that observed in this study, only for undrawn amorphous

NHS

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Fg. 5. DSC thermogramsfor PET frlms. Effect of heat-set temperature (tHs= 30 sec).

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Fg. 6. DSC thennograms for PEN

Jim. Effect of heat-set temperature (tHs = 30 s=).

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J. Greener, A. H . Tsou, and T. N.Blanton PET, while the secondary melting temperatures for highly drawn (5x)fibers were generally much higher than the treatment temperatures and strongly overlapped with the primary peaks. A similar result for highly drawn fibers was noted by Buckley and Salem (12). Oswald et aL (26)who studied the annealing of PET fibers without constraint (not strictly heat setting!) show similar results to the ones reported here except that their secondary melting peaks appear to be weakly dependent on anneal time beyond thermal equilibration, contrary to our observations. The evolution of the secondary peak with time for PET film is illustrated in Figure 8. The peaks for short anneal times (i.e., before equilibration) appear at lower temperatures, i.e., the effective temperature "seen" by the material prior to full equilibration is lower due to thermal lag in the film.These results as well as similar results for PEN film,together with the transient density data (Figures 1 , 2).validate the conclusion that,

Hg. 7. Position of secondary melting peak (Tm') US. heat-~et temperature.Line represents identity.

under the experimental conditions of this study, the kinetics of the heat setting process is controlled by thermal diffusivity. Another characterization method used to probe the microstructure of the heat-treated films is thermomechanical analysis m. The TMA test traces the dimensional change (shnnkage or expansion) of the sample by scanning over a set temperature range at a given rate. TMA scans for the test materials at various heatset temperatures (tHs = 30 sec) are shown in FIgwes 9 and 10. The most striking feature in these traces is the appearance of severe shrinkage in the f ilm above some critical temperature, which seems to roughly coincide with the corresponding heat-set temperature. The results for PET and PEN suggest that the secondary crystalline structure formed during heat setting has an "interlocking"effect on the sample, which prevents shnnkage or disorientation of the amorphous phase at temperahms below the melting of the "secondary"crys-

p - 200 t-' 150

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POLYMER ENGINEERINGAND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

Effects of Heat Setting in Polyester Films

tallites. However, once these crystallites are melted out, at T > THs, the highly strained amorphous phase is free to disorient (shrink) and thereby lower its free energy. By contrast, it appears that the "primary" crystalline phase formed during the stretching steps (and partly during the DSC scan (22))does not have the "interlocking" structure needed to prevent slippage and retraction of the oriented amorphous domains above the glass transition temperature. It is interesting to note, however, that even at high THs. both films exhibit a mild inflection in the TMA trace (slight shnnkage) at a temperature that is nearly coincident with the glass transition temperature of the material (- 94 and 120°C for PET and PEN, respectively). This indicates that a small portion of the amorphous phase is not fully "locked" by the crystalline network formed during heat setting and it thus can partially retract once the amorphous domains gain sufficient mobility at Tg The evolution in the shnnkage response of PET film with an-

neal time for THs = 200°C is shown in Figure 1 1 . The onset of shrinkage in the TMA traces appears to be shifted to lower temperatures for short exposure times, in line with our observation that incomplete heating due to thermal lag corresponds to exposure to lower THs(cf.Rgwe 9). The attendant morphological changes during heat setting also have a pronounced effecton the dynamicmechanical response of the films. Figures 12 and 13 show isochronous plots of loss tangent vs. temperature for the test materials at several heat-set temperatures. Two outstanding features are noted for both materials: (i) the primary CY peak, representing the effective glass transition, is strongly suppressed by heat setting, and (ii)the position of the CY peak (TJ first increases and then decreases with THS. These effects are illustrated in Figures 14 and 15. In both cases, the peak height drops with temperature to an asymptotic value at 200°C. with the absolute change in peak

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C)

T,s( R g . 9. RMA scan for PET@

o NHS

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change in sample dimension (along MD) us. temperahre. Effect of heat-set temperature (tHs= 30 sec).

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J. Greener, A. H . Tsou, and T. N . B h t o n

The optical properties of the stretched films are also strongly affected by the heat-setting process. Generally, in line with the Lorentz-Lorenz equation, we observe a nearly linear dependence of the average refractive index on density. But, beyond this change, the molecular orientation (or optical anisotropy) in the polymer matrix, initially formed during the stretching steps of the film process, is also modified by the heat setting treatment. The effect of THS on the planar and in-plane birefringence components (cf.Eq 1) is shown in Figures 16 and 17. The value of the planar birefringence (PBR)is considerably higher for PEN owing to a higher planar alignment of the bulky and "flat" naphtyl moiety in this polymer. But in both cases, PBR shows a similar dependence on temperature; it initially increases then drops with THs. as in the case of T, (Rgure 15).The shallow peak in PBR occurs at a slightly higher temperature than the peak in T,, and the value

height being generally greater for the PEN film. Similarly, the peak temperature (closely related to Tgl appears to go through a shallow maximum (at 180°C for PET and 200°C for PEN) then drop substantially with THS (Figure 15). The suppression of the tan 6 peak can be attributed to the decrease in the amorphous fraction, consistent with the results in Rgure 4, but the decrease in the position of the peak at high heat-set temperatures suggests the possibility of some microstructural transformation beyond a certain temperature threshold (THs*). This transformation apparently removes some constraints from the interlocked amorphous domains, thus increasing their segmental mobility and, thereby, lowering the glass transition temperature of the film. A similar conclusion was arrived a t by Illers and Breuer (27) based on similar results on the effect of crystallinity on the a transition temperature of PET.

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Rg. 11. TuA scan for P E T m % change in sample dimension (along MD) us. .Eff& of heat-settime pr,, = 200°C).

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.36 .32 --

.28 Rg. 12. Isochronous dynamic-mechanical response of PETjilms. Loss tnngent us. ternperdure at 1 H z . Eff& of heat-settemperature: (4)N H S f i l m 160, (c) 180,(4 200,(e)220 and (fl 240°C. (tHs =

30 see).

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POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Voi. 39, No. 12

Effectsof Heat Setting in Polyester Films

.36 .32 .28 lQg. 13. Isochrono~~~ dynamicmechanical response of PEN films. Loss tangent us. temperature at 1 Hz. Effzct of heat-set temperature: (a) NHS jXm. [b) 140,[cl 180, (d) 220, and (el 240°C (t,, = 30 sec).

a c CCI

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.Rg. 14. Loss tangent peak height (at T,] u s . heat-set temperature It, = 30 sec). Arrows indicate

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J. Greener,A. H . Tsou. and T. N. Blanton

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mode 8/28 diffraction scans for PET and PEN films annealed at 200°C for 30 sec are shown in Ffsure 18. The PET pattern is characterized by a strong 1 0 0 diffraction peak at 28 = 25.5'. The crystal structure of PEN is different from that of PET (20,28).The peak at 28 = 26"in Figure 19 (for PEN) is assigned to the T10 reflection. Transmission mode 8/28 diffraction patterns are shown in Ffsure 18. The PET diffraction pattern features a n intense peak at 28 = 43'. which is assigned to the 705 plane, while the PEN dif€raction pattern is dominated by an intense 010 peak at 28 = 14.5".Since the 105 (for PET) and 010 (for PEN) planes are closely aligned with the chain axes of the respective polymers, the corresponding transmission peaks were used to generate complete pole figures for both film types. Figure 20 shows the effect of heat-set temperature (@ 30 sec anneal time) on the crystallite size as estimated from the breadth of a characteristic reflection

for PEN (- 220°C)is 20°C higher than for PET. This behavior, again, underlies some morphological changes taking place at some high temperature. T H d , near the onset of the primary melting range of the polymer. Essentially, the planar alignment is enhanced by heat setting at low temperatures but is clearly suppressed at temperatures approaching the primary melting range of the material. The in-plane birefringence (BR) for PET also shows a systematic change with THs.see Figure 17,with the material becoming more aligned towards the transverse direction (negative BR) at higher temperatures > 220°C).In line with the PBR data, the orientation of PEN is nearly independent of the heat-set temperature, indicating that the relatively high molecular ahgnment attained during the stretching steps, is not modified to any measurable extent by the heat-setting treatment of these films. Wide-angle X-ray ( W e ) diffraction was used to probe the structure of the crystalline phase. Reflection

nHs

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FXg. 16. Planar birefringence us. heat-set temperature (tHs = 30 sec). A m w s indicate NHSJilms.

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Effects of Heat Setting in Polyester Films (105 for PET and 010 for PEN), based on the Schemer equation (18).Despite significant differences in crystal habit, both materials exhibit similar levels of crystallite size and similar dependence of crystallite size on THS; over the range of temperatures covered, the crystallites grow steadily with THs with a relatively steep upturn at the highest temperatures for both materials. These results are in line with observations of other investigators (6, 9, 10, 17) and may suggest, again, the existence of some morphological transformation at temperatures near the foot of the primary melting range of the material. The effect of heat-set temperature on the planar and in-plane crystalline orientation functions, as expressed by HOFc and HOFa (cf. Eq 2). is shown in FIgues 21 and 22. The planar orientation (HOFc)of the PET films follows closely the results for the planar birefringence (Figure 16). with the orientation peaking at 220°C. However, the planar orientation of the PEN films is considerably lower and appears to be independent of THs. Since the birefringence represents the combined contributions of the amorphous and crystalline phases to the molecular orientation, the data in Figure 21 imply that the main contribution to the high planar birefringence of the PEN films is made by the highly oriented amorphous phase. The in-plane orientation function (Figure 22) follows a similar pattern as the in-plane birefringence (Figure 17). The orientation of the PEN

films is relatively unaffected by the heat-set conditions, while the PET crystals seem to realign towards the transverse direction with increase in THs. The PET crystallites are initially w e d in the machine direction but, with increasing heat-set temperature, they become increasingly aligned in the transverse direction. The overall orientation patterns as extracted from the X-ray and birefringence data are generally consistent and they indicate that the heat-setting process is capable of imparting significant molecular realignment in PET films but only subtle changes in the orientation of the PEN films are observed. (Note that the reference direction in the X-ray analysis is the transverse direction, while for the birefringence measurements the orientation is referenced to the machine direction).

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4. DISCUSSION

In line with previous investigations, this study illustrates the pronounced effect of heat setting treatment on the microstructure and macroscopic properties of two polyester films-PEN and PET; these effects are manifested by notable changes in crystallinity, crystallite size, density, orientation, dimensional stability and other related properties. Despite differences in molecular structure, crystal habit and crystallization kinetics, both polyesters appear to undergo similar morphological changes in response to the controlled

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J. Greener, A. H . Tsou, and T.N.Blanton in the crystal morphology of the oriented matrix: (i) sharpening of diffraction nodes upon annealing, which represents a reduction in lattice and conformational defects, (ii) formation of a 'secondary" crystalline morphology in the form of folded chain lamellae, which grow perpendicular to the chain axis or drawing direction. These lamellae are superimposed on top of the "primary" fibrillar structure generated during the stretching steps by a strain-induced crystallization process, thus forming a paracrystalline "mosaic" network. It is very likely that the lamellar crystals nucleate and grow epitaxially on the existing microfibrils and are organized in a "shish kebab"-like superstructure. These morphological changes are represented schematically in Figure 23. The lamellae grow in size with increase in THs because of the increased fraction of taut-tie amorphous chains with sufficient mobility to partake in the secondary crystallization process that involves chain fold-

heat setting treatment applied in this study. Aside from differences in temperature dependence which scale with the correspondingglass transition temperatures, the main difference between the PET and PEN films concerns the effect of heat setting on molecular orientation. Based on birefringence and X-ray data, the PET films exhibit significant molecular realignment during heat setting while the apparent orientation of the PEN films is only weakly affected by this treatment. These differences notwithstanding, the morphological transformation imparted by heat setting and the attendant changes in physical response appear to be similar for both materials. The morphological models of Fischer and Fakirov (5, 6)and Schultz and coworkers (15. 16) provide a useful framework for discussing the results of this study. Both models deduce from small-angle X-ray scattering (SAXS) data on oriented PET fibers that constrained annealinggives rise to two distinct changes

80 I

I o PET

70

PEN

ii

Q) .-N 60

0

Rg. 20. Crystdite size us. heatset temperature itHs = 30 sec). Arrows indicate NHS-.

50

40

,i

30 50

100

I

I

I

150

200

250

Heat-set temperature ("C)

o PET

PEN -0.1

Rg.21. planar crystalline orientation (HOFc) us. heat-set temperature (tHs = 30 sec). Arrows indi-

0

bI -o.2

CateNHSjiIms.

-0.3

-0.4 50

100

150

Heat-set temperature

2414

200 ("C)

250

POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

Effectsof Heat Setting in Polyester Films 0.8 o PET

0.6

0

I

PEN

/

0.4

a u-

Fig. 22. In-plane crystalline orientation HOFd us. heat-set temperahve (tHs = 30 sec). Arrows indicate NHSfilms.

0 I

0.2 0

/

-0.2

/-

A

-0.4 100

50

150

200

250

Heat-set temperature ("C) ing. Consequently, the density (or the crystalline fraction). the average refractive index as well as the crystallite size grow with annealing temperature. However, since the secondary crystals are attached to highly oriented amorphous ("taut-tie") chains, they are relatively unstable and have, therefore, a suppressed melting point Indeed, the corresponding melting point, based on DSC data, is nearly coincident with the formation temperature of the crystal (%we 3, indicating that once the taut-tie chains regain sufficient mobility. the secondary structures melt out due to the potential gain in free energy resulting from an increase in entropy of the surrounding amorphous domains, i.e.,

urn').

AH:, T' = __ rn AS,!,,

recrystallize into more stable folded chain crystals that ultimately melt at T, so that the secondary melting peak represents only a fraction of the crystalline matter formed during heat setting. Also, as noted above, some of the crystallinity gain is associated with the perfection of the existing fibrilar crystallites (rather than buildup of secondary lamellae) and is, therefore, not represented by the secondary melting peak. It should be noted that the dual-melting behavior of PET

NHS

THS

< 'AS

'HS

''is

(4)

where AH,,,' and ASrn' are the changes in enthalpy and entropy for the secondary melting process. The depression in the secondary melting point can also be attributed to the observed changes in lamellar thickness through (29):

where Tmoand AHf are the melting point and latent heat of fusion of a perfect crystal (crystal with infinite dimensions), 1 is the average thickness of the lamellar crystal and uc is the surface free energy of the crystal. By adjusting the parameters in this equation it is possible to show that the observed changes in crystallite size, 1, (Figure 20) can be directly correlated with the changes in the secondary melting temperature. The apparent enthalpy associated with the secondary melting endotherm, based on the DSC data (Figures 5 , 6).is relatively small compared to the total amount of crystalline matter formed during the heatsetting process (cf. Figure 4 ) . This implies that upon melting (at Tm')the secondary crystallites immediately

Fibrillar

Fibrillar-Lamellar

Lamellar

Q. 23. Schematic m~rphologicalmodel o f f l d h - b - h l h

transifion (After Schultz et aL (15,16))showing the evolution of the crystalline network with increase in heat-set temperah e . Scheme shown here for miaxially d r a w n m can be generalizedfor biaxially oriented-

POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

2415

J. Greener, A. H. Tsou, and T. N. Blanton

and other polymers has been the subject of some debate in the open literature (6, 17, 22-26) and the general picture painted above is not universally accepted. It is generally agreed, however, that the two melting endotherms represent two distinct crystalline morphologies that depend on the thermal history and onentation of the polymer matrix and annealing mode (constrained vs. free) of the film sample. One form, corresponding to the secondary melt temperature, T,’, is closely associated with the heat-setting process while the second form, corresponding to T, is practically independent of the heat-set conditions. Holdsworth and TurnerJones (22) have proposed, in fact, , , is the melting that the primary melting peak, T point of stable folded chain crystals that are formed throughout the DSC scan by a process of continuous melting and recrystallization, while the secondary peaks are the melting transitions of the thin crystalline lamellae formed during heat treatment. Oswald et aL (26)have concluded, from studies on annealing of unconstrained PET fibers, that the secondary melting peak is a result of lamellar crystals (“nuclei”)formed during heat treatment on the present fibrils, with the melt temperature being related to the size of the crystals. They observed, however, that the secondary melting temperature increased somewhat with anneal time, unlike the result of this study in which the asymptotic value of the secondary melt temperature is constant (annealingtime of ca. 60 min gave essentially identical T,’ as 30 sec). This difference can be attributed to the fact that our samples were annealed under lateral constraint. Oswald et d ’ s interpretation of the effects of heat treatment on the nature of the secondary melting process is, however, similar to the mechanism given above. Although the general morphological changes during heat setting observed in this and other studies are consistent with the phenomenological models of Fischer and Fakirov (5, 6)and Schultz and coworkers (15, 16), neither model pays particular attention to the apparent morphological transition at high heat-set This transition is marked by distemperatures (THs*). tinct maxima in planar birekingence and glass transition temperature (TJand abrupt change in crystallite size and in-plane orientation at temperatures close to the primary melting range of the material. Gohil (17) first noted this “transition,”which he attributes to a peak in crystallization kinetics, and assigned it as the boundary between two temperature regimes for heat setting. This transition can also be interpreted by the phenomenological model of Schultz et aL as the onset of the primary melting range of the material where the fibrillar core structure starts to melt out and recrystallize into a more stable lamellar form such that the general population of fibrilar crystals is depleted and replaced with lamellar crystals. In support of this idea, Statton et aL (7) and Wilson (8)observed that anneahg of drawn PET at high temperatures (>2OO0C) gives rise to a significant increase in the number of chain folds which is characteristic of lamellar crystals. 2416

Through their interlocking (crosslinking) effect, the secondary crystallites restrict the mobility of the oriented amorphous chains, thereby enhancing the dimensional stability and s W e s s (modulus)of the matrix. As soon as these crystallites are melted out (at T,‘), the oriented amorphous (taut-tie)chains are free ilm exhibits significant shrinkage to disorient and the f (Figures 9-11). The large distortion beyond T,’ suggests that once the secondary crystals are melted, there is nothing to prevent the highly strained amorphous domains from collapsing and moving the system into a lower free energy state. These secondary crystallites formed during heat setting are qualitatively different and apparently much less stable than those formed during stretching. The prevailing morphology of the strain-induced (primary)crystallites is that of a “shish kebab” or a bundle-like microfibril (30).see Figure 23, which contains a relatively small number of chain folds in its fibrillar core (“shish”)and small lamellae growing perpendicular to the draw direction (“kebab”).Unlike the secondary lamellae, these crystals apparently cannot prevent the amorphous chains, interspersed between the microfibrils, from disorienting (shrinking)at temperatures above the glass transition, although they contribute to the overall mosaic structure of the crystalline network and thus enhance the stiffness and strength of the matrix (at least along the draw direction).Once the melting range of the fibrillar crystals (or their stable recrystallized form) is approached, these crystals start to melt and transform into more stable folded chain lamellae. Very likely, this transformation gives rise to the decrease in planar birefringence and the glass transition temperature, T, , above THs*. The decrease in T, suggests also that the transformation at THs*‘loosens up’ the amorphous domains and imparts higher segmental mobility to the noncrystallized chains, as suggested by Illers and Breuer (27)in a study on the effect of crystallinity on the dynamic-mechanical properties of PET. Elenga et aL (25) also studied high-temperature annealing of oriented PET films, but they argue that the observed changes in mechanical properties after long annealing at high temperatures are due to a chemical transformation caused by a transesterification reaction. This, of course, cannot be ruled out and it may be happening in conjunction with the melting of the primary fibrillar crystals (or recrystallized folded chain crystals) over longer time scales. Whatever the origin of the morphological transformation, it must be recognized that there exists a critical heat-set temperature, THs*, beyond which the mechanical and physical performance of the film is qualitatively altered. 5. SUMMARY AND CONCLUSIONS

The effect of a controlled heat setting treatment on the microstructure and physical performance of PET and PEN films has been studied over a wide range of conditions. Heat setting was simulated by soaking laterally constrained, biaxially oriented films in a thermostated oil bath at a constant temperature for a

POLYMER ENGINEERING AND SCIENCE, DECEMBER 1999, Vol. 39, No. 12

Effects of Heat Setting in Polyester Films

set time. The heat-treated films were then characterized by densitometry, differential scanning calorimetry, refractometry, dynamic-mechanical analysis, thermomechanical analysis, and wide-angle X-ray diffraction. The salient observations of this study can be’ summarized as follows: 1. Heat setting can substantially increase the degree of crystallinity and crystallite size of biaxially stretched films. The crystalline fraction grows nearly linearly with heat-set temperature and is nearly doubled over the range of temperatures covered in this study. The changes in the crystalline fraction correspond to similar changes in density and average refractive index. 2. A secondary melting endotherm was detected in

DSC thermograms for both film types at a temperature that nearly coincides with the applied heat-set temperature. The secondary melting temperature does not change with anneal time once thermal equilibration is reached and it involves considerably lower enthalpy change compared to the primary melting transition of the polymer.

3. The heat-set temperature has a pronounced effect on the dimensional stability of the oriented film. Both PET and PEN films exhibit sigmficant shrinkage a t temperatures near and above the corresponding heat-set temperature. Thus, raising the heat-set temperature should improve the overall high-temperature dimensional stability of the stretched film. 4. X-ray and birefringence data indicate that the

changes in crystalline morphology are accompanied by some molecular realignment in the polymer matrix during heat setting. The changes in orientation of the PEN films are small compared to those in the PET films. In both cases, however, the planar orientation goes through a broad maximum at temperatures close to the onset of the primary melting range of the material. Also, the crystallite size grows steadily with heat-set temperature, with some indication of a sharp upturn near the onset of the primary melting range. 5. Heat setting leads to a strong suppression and a shift in the position of the loss tangent peak corre-

sponding to the (Y transition (related to Tg). Like the planar birefringence, the position of the o transition (position of tan S peak) first increases and then decreases with THs. The maximum in the a transition temperature occurs near the onset of the primary melting range of the material. 6. The maxima in planar birefringence and Tg and, in the case of PET, the steep increase in in-plane

orientation at temperatures close to the onset of the primary melting range of the respective materials, suggest the presence of a morphological transformation at some high temperature, THs*. This transformation is associated with changes in

the physical and mechanical properties of the film, in line with observations made by Gohil ( 17). 7. A general morpholo@cal model, based on previous suggestions by Fischer and Fakirov (5, 6) and Schultz et al. (15,16), is invoked to explain our data. The basic premise of the model is that the heat setting process induces the formation of chain folded lamellae, which nucleate epibxially on the fibrillar structures formed during the sequential stretching steps. The lamellae grow perpendicular to the stretching direction, with the lamellar size being directly dependent on the heatset temperature, thus producing an “interlocking” paracrystalline network that enhances the dimensional stability of the oriented matrix. The transition corresponding to THs*may be due to partial melting and recrystallizing of the fibrillar core crystals in the vicinity of the onset of the primary melting range of the polymer. ACKNOWLEDGMENTS

We thank Mr. G. Mosehauer, Ms. E. Priebe, Ms. J. S . Machell, Ms. B. A. Contestable, Ms. P. Sciotti, Mr. R. Schlotzhauer, Mr. C . Barnes and Ms. A. M a m e for assistance with the experimental work. REFERENCES 1. J . M. Schultz, in Solid State Behauwr of Linear Polyesters and Polyamide~,pp. 97-103, J. M. Schultz and S. Fakirov, eds., Prentice-Hall, New York (1991). 2. F. S. Smith and R. D. Steward, Polymer, 15,283 (1974). 3. D. S. Prevorsek, G . A. “lrpak, P. J. Harget and A. C. Reimchuessel, J. MacromoL Sci Phys., BB,733 (1974). 4. S. Fakirov. E. W. Fischer and G . F. Schmidt, MacromoL Chen,176,2459 (1975). 5. E. W.Fischer and S. Fakirov, J. Mater. Sci, 11, 1041 (1976). 6. S. Fakirov, E. W. Fischer. R. Hoffmann a n d G. F. Schmidt, Polymer, 18, 1121 (1977). 7. W. 0. Statton. J. L. Koenig and M. Hannon, J. Appl. Phys., 41,4290 (1975). 8. M. P. Wilson, Polymer, 15.277 (1976). 9. A. M. Hindeleh and D. J. Johnson, Polymer, la, 27 ( 1978). 10. V. B. Gupta and S. Kumar, J . Appl. Polyrn Sci., 26. 1865 (1981). 11. V. B. Gupta, C. m e s h and A. K. Gupta, J. AppL Polym Sci. 2s. 3115 (1984). 12. C . P. Buckley and D. R. Salem, Polymer, 28, 69 (1987). 13. C. P. Buckley and D. R. Salem, J. AppL Polyrn Sci, 41, 1707 (1990). 14. C. D. Bechev and J. T. Mishinev, J. Appl. P o l y m Sci. 48,29 (1992). 15. H. Chang, J. M. Schultz and R. M. Gohil, J . MacromoL Sci Phys.. BS2, 99 (1993). 16. K.-G. Lee and J. M. Schultz, Polymer, 34,4455 (1993); see also H. Chang, K. G. Lee and J. M. Schultz, J. MacromoL Sci Phys., BSS, 105 (1994). 17. R. M. Gohil,J. AppL Polyrn Sci,52,925 (1994). 18. L. E. Alexander, X-Ray Dt@?action M e t h o d s in Polymer Science, Wiley Interscience, New York (1969);see also, B. D. Cullity, Elements o f X - R a y DifF.action, 2nd Ed., Addison-Wesley, Reading, Mass. (1975). 19. R. de P. Daubeny. C. W. Bunn and C. J. Brown, Proc. Roy. Soc. London Ser. A 226, 531 (1954). 20. 2. Mencik, C h e m Prim, 17,78 (1967).

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J. Greener, A. H. Tsou, and T.N. Blanton 21. S. Buchner. D. Wiswe and H. G. Zachmann, Polymer, 30, 480 (1989). 22. P. J. Holdsworth and A. Turner-Jones, Polymer, la, 195 (1971). 23.J. P. Bell and J. H. Dumbleton. J. PoZym Sci. PoZym J. P. Bell and T. Murayama, Phys. Ed,7,1033 (1969); Ibid., p. 1059. 24.D. L. Nealy, T. G. Davis and C. J. Kibler, J. Polym Sci, Polym Phys.Ed.8,2141 (1970). 25. R. Elenga, R Seguela and F. Rietsch, PoZymer,Sa, 1975 ( 1991).

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26. H.J. Oswald, E. A. Turi, P. J. Harget and Y. P. Khanna, J. MacromL Sci Phys.,B13.231 (1977). 27. K.H. Illers and H. Breuer, J. Colloid Sci, 18,1 (1963). 28. M. Cakmak.Y. D. Wang and M. Simhambhatla, Po4m Eng. Sci. SO,721 (1990). 29. B. Wunderlich, Macromolecular Physics. VoL 3. Crystal Melting,pp. 30-33,Academic Press,New York (1980). 30.J. H. Magill. in ?)-eatiseon Materials Science and Techno@~, VoL 10, Properties of Solid Polymeric Materials (part A),pp. 89-122.J. M. Schultz, ed., Academic Press. New York (1977).

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