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Synthesis and non-isothermal crystallization behavior of poly(ethylene phthalate-co-terephthalate)s.


Poly(ethylene terephthalate) (PET) has excellent thermal and mechanical properties, high chemical resistance, and low gas permeability (1-3). Because of these advantageous properties, this polyester is widely used in synthetic fibers, packaging films, recording and photographic tapes, bottles for beverages and food, and engineering plastic components (1-3). However, PET has some undesirable properties, for example poor dyeability, pilling, low moisture regain, generation of static charges, poor adhesion to metals, and poor processability because of its high melting temperature. Copolymerization generally provides a facile means of modifying crystallinity, morphology, melting point ([T.sub.m]), glass transition temperature ([T.sub.g]), solubility, and permeability. However, when copolymerization is used to modify a particular property, it may also cause unwanted alterations of other properties. Thus, to gain the advantages offered by copolymerization, it may be necessary to incorporate a third component into the PET polymer in order to overcome the property changes brought about by the additional undesirable alterations.

A representative PET copolyester is poly(ethylene isophthalate-co-terephthalate) (PEIT), which contains crystallizing and noncrystallizing components. PEIT is currently used in the polymer industry in thermally shrinkable package films as well as in heat-sealable laminating films for steel cans, and for metal and ceramic sheets (4-8). The structure, crystallization behavior, and properties of the PEIT copolymer have been studied by several research groups (3, 4, 9-11). Another analogue of PET copolyester is poly(ethylene phthalate-co-terephthalate) (PEPT). PEPT contains the phthalate unit, which is an isomer of the terephthalate unit in the PET homopolymer and of the isophthalate unit in the PEIT copolymer. In contrast to the many studies of PET and PEIT polymers, the PEPT copolyester has rarely been studied (12-14).

In general, the properties of a crystallizable polymer are strongly dependent on the morphological structure (size, shape, perfection, volume fraction, and orientation of crystallites) formed by crystallization from the molten state. Thus for crystallizable polymers, elucidation of the crystallization behavior is crucial to the aim of controlling the polymers' morphological structure and properties. In particular, the behavior of crystallizable polymers during non-isothermal crystallization from the melt is of increasing technological importance, because these conditions are the closest to real industrial processing conditions. Therefore, information on the non-isothermal crystallization behavior is needed if we are to optimize industrial processes and to understand the properties of the processed products.

In this study, a series of PEPT copolyesters of varying composition were synthesized by the conventional bulk polycondensations of dimethyl terephthalate (DMT) and ethylene glycol (EG), using dimethyl phthalate (DMP) as a third component in order to combine the advantageous properties of the respective polymers in one polymer system and to minimize their disadvantageous properties. The compositions and molecular weights of the copolymers were determined by proton nuclear magnetic resonance ([.sup.1]H-NMR) spectroscopy and viscometry respectively. Their thermal properties were measured by differential scanning calorimetry (DSC). In addition, their non-isothermal crystallization behaviors were measured by DSC, with particular emphasis on analyzing the effect of the third component.


Materials and Polymerization

Polymerization-grade DMT was obtained from SKC Limited (Korea), and polymerization-grade EG was obtained from LG Chemical Company (Korea). All other chemicals were purchased from Aldrich Chemical Company (USA). All chemicals in this study were used as received.

A poly(ethylene phthalate-co-terephthalate) (PEPT) of 50/50 DMT/DMP (molar ratio) composition (DMP50) was synthesized by a two-step reaction sequence as follows (see Fig. 1). In the first step, DMT (0.5 molar equivalents) and DMP (0.5 molar equivalents) were added into EG (2.2 molar equivalents) in a two-neck flask equipped with a mechanically sealed stirrer and a motor, followed by addition of zinc acetate dihydrate (2.7 X [10.sup.-4] wt% of the total amount of DMT and DMP) as an ester interchange reaction catalyst. After the addition was complete, a condenser was added to the flask. Then the ester interchange reaction was conducted with stirring for 2.0-2.5 h at 180[degrees]C-210[degrees]C. During the reaction, byproduct methanol was removed from the reaction mixture with a yield of 92%. In the second step, antimony acetate (9.9 X [10.sup.-4] wt% of the total amount of DMT and DMP) and trimethyl phosphate (1.5 X [10.sup.-3] wt% of the total amount of DMT and DMP) were added into the reaction mixture as a polycondensation catalyst and thermal stabilizer respectively. The polymerization was carried out with stirring for 2.5-3.5 h at 260[degrees]C-280[degrees]C under a vacuum of approximately 1 X [10.sup.-2] Torr as described elsewhere (10, 15). In the same manner, a series of PEPT copolyesters with other compositions were prepared by varying the proportions of the DMT and DMP monomers with respect to that of the EG monomer.



The compositions of the synthesized copolyesters were determined in a mixture of trifluoroacetic acid-[d.sub.1] (C[F.sub.3]COOD) and chloroform-[d.sub.1] (CDC[l.sub.3]) (3:1 in volume) using a proton nuclear magnetic resonance ([.sup.1]H-NMR) spectrometer (Bruker Aspect 300 MHz); and the intrinsic viscosities [[eta]] were measured in trifluoroacetic acid at 30.0[degrees]C [+ or -] 0.1[degrees]C using an Ubbelohde suspended level capillary viscometer. The weight average molecular weights ([bar.M.sub.w]) were estimated from the measured intrinsic viscosities with the Mark-Houwink-Sakurada equation, using values for the constant [K.sub.[eta]] of 4.33 X [10.sup.-4] and for the exponent [alpha] of 0.68, as determined previously for the PET homopolymer (16).

The thermal characterizations were performed under a dry nitrogen atmosphere using a differential scanning calorimeter (DSC) (Seiko DSC 220CU). Both temperature and heat flow at a chosen rate of heating and cooling were calibrated using indium and tin standards. Film sample specimens of 3.0-5.0 mg were used. In the DSC measurements, each polymer sample was preheated at 290[degrees]C for 10 min in order to remove its thermal history and then cooled to 0[degrees]C at a rate of 10.0[degrees]C/min or quenched to 0[degrees]C. Subsequently, each cooled sample was again heated to 290[degrees]C at a rate of 5.0[degrees]C/min, whereas the quenched samples were again heated to 290[degrees]C at a rate of 10.0[degrees]C/min; in case of the quenched samples the heating rate of 10.0[degrees]C/min (i.e. a properly high heating rate) showed glass transition more clearly, compared to the heating rate of 5.0[degrees]C/min. [T.sub.m] was determined for each sample from the DSC thermogram as the temperature corresponding to the maximum of the melting transition during the heating run. The heats of fusion ([DELTA][H.sub.f]) were also determined from these thermograms. The [T.sub.g] of each sample was taken as the temperature at the middle point of the glass transition during the heating run.

For the crystallizable polymers, non-isothermal crystallizations were performed with varying cooling rates in the range 2.5[degrees]C/min-20.0[degrees]C/min. Each sample was heated to 290[degrees]C at a rate of 80[degrees]C/min and then held at that temperature for 10 min to remove thermal history effects, followed by cooling to 0[degrees]C at a chosen cooling rate. Crystallization peak temperatures ([T.sub.p]) were estimated from the maxima of the DSC thermograms. In addition, the heats of crystallization ([DELTA][H.sub.c]) were obtained from the DSC thermograms.

The equilibrium-melting temperatures ([T[degrees].sub.m]) were measured as follows. Each polymer sample was first heated to 290[degrees]C at 80[degrees]C/min and maintained for 10 min at that temperature, after which it was quenched to a chosen crystallization temperature ([T.sub.c]) and isothermally crystallized for 40 min at that temperature. The sample was then reheated at a rate of 10.0[degrees]C/min to the melt state. From the DSC thermogram measured during the reheating run, [T.sub.m] was obtained. For each sample, the measured melting temperatures were plotted as a function of crystallization temperature [T.sub.c] and then extrapolated to the line of [T.sub.m] = [T.sub.c], giving [T[degrees].sub.m].

Modified Avrami Analysis

The non-isothermal exotherms obtained in the DSC measurements were analyzed by the modified Avrami method (10, 17-20) as follows. For non-isothermal crystallization at a given cooling rate, the relative crystallinity ([X.sub.t]) is a function of the crystallization temperature (10, 17-20):


where T is the crystallization temperature and [T.sub.0] and [T.sub.[infinity]] represent the onset and end temperature of crystallization respectively. [X.sub.t] can be expressed as a function of crystallization time t according to the Avrami equation (10, 17-20):

1 - [X.sub.t] = exp (-k[t.sup.n]) (2)

where n is the Avrami exponent and k is the growth rate constant. The crystallization time t (which is the socalled equivalent time) is obtained from the cooling rate C and the crystallization temperatures T and [T.sub.0] as follows:

t = [[T.sub.0] - T]/C (3)

The rate of non-isothermal crystallization depends upon the cooling rate employed, so the growth rate constant k needs to be corrected as follows:

log k' = [log k]/C (4)

The crystallization half time [t.sub.1/2] is defined as the time at which the development of crystallization is 50% complete, and can be determined from the corrected crystallization rate constant k':

[t.sub.1/2] = ([ln 2]/k')[.sup.1/n] (5)



Poly(ethylene phthalate) (PEP), PET homopolymer, and their copolymers were synthesized by varying the feed ratio in the bulk polymerization of DMT and DMP with respect to the proportion of the EG monomer. The synthesized homopolymers and copolymers are listed in Table 1. The chemical compositions of these polymers were determined by [.sup.1]H-NMR spectroscopy. Figure 2 shows the [.sup.1]H-NMR spectrum of the DMP18 copolymer, which is typical of the polymer spectra. This copolymer was determined to have 17.7 mol% ethylene phthalate units and 82.3 mol% ethylene terephthalate units from the integrations of the peaks associated with these units, as shown in Fig. 2. All the other polymers were characterized in the same manner; the results are summarized in Table 1. In addition, intrinsic viscosity [[eta]] measurements were performed for all the synthesized polymers; the [[eta]] values range from 0.125 to 0.572 (see Table 1). From the measured [[eta]] values, the [bar.M.sub.w]s of the polymers were found to vary with composition over the range 4100-38,000 (see Table 1).

The DMP content of the polymers is always less than that fed into the bulk polymerization, as shown in Table 1. Moreover, for the copolymers containing > 32 mol% DMP, [bar.M.sub.w] decreases as the DMP content increases. These results are influenced by two main factors. The first factor is the steric hindrance of the DMP monomer and its effect on the polycondensation reaction. DMP is the ortho-isomer of the DMT monomer and hence has carboxylate groups at the 1- and 2-positions of its phenyl ring, creating highly kinked units in the polymer backbone when incorporated into the polymer chain. Such highly kinked units result in large steric hindrance of the growing polymer chain during polymerization, restricting the growth of the polymer chain. In contrast, the DMT monomer has carboxylate groups at the 1- and 4-positions of the phenyl ring, producing a linear unit in the polymer chain. Thus DMT favors the growth of the polymer chains during polymerization. The second factor is the approximate nature of the Mark-Houwink-Sakurada constants employed for the estimations of the [bar.M.sub.w]s from the measured intrinsic viscosities. In fact, the Mark-Houwink-Sakurada constants used in this study were measured only for the PET homopolymer, rather than for the PEP homopolymer and PEPT copolymers. However, the values of these constants for the PEP and PEPT copolymers may differ from those for the PET homopolymer. Thus the [bar.M.sub.w] values of the PEP homopolymer and the copolymers might be underestimated or overestimated.

Thermal Properties

DSC measurements were conducted for all the polymer samples prepared by cooling at 10.0[degrees]C/min from the melt. The results are summarized in Table 2. As shown in Table 2, PET and the copolymers that contain [less than or equal to] 21 mol% DMP units are crystallizable but the other copolymers are amorphous. For the crystallizable polymers, the melting temperature [T.sub.m], the heat of fusion [DELTA][H.sub.f], and the crystallinity [X.sub.c] all decrease as the DMP content increases. These results may be due to several factors that relate to the DMP content of the polymer chain. First, the incorporated DMP units disturb the chain regularity that is necessary for molecular ordering. The disruption of the chain regularity may negatively affect the crystal formation, thus reducing [DELTA][H.sub.f] and [X.sub.c]. Second, the incorporated DMP units shorten the length of the crystallizable ethylene terephthalate block within the polymer chain. Such shortened crystallizable blocks may crystallize as relatively small, thin lamellar crystals that have low [T.sub.m]. Finally, the incorporated DMP units may be excluded from lamellar crystal formation, increasing the fraction of amorphous phase in the polymer.


For each sample quenched from the melt, the glass transition temperature [T.sub.g] was measured by DSC. The measured glass transition temperatures are listed in Table 2. The PET homopolymer has a [T.sub.g] of 78.5[degrees]C, whereas the PEP homopolymer has a [T.sub.g] of 38.2[degrees]C. The copolymers exhibit glass transition temperatures that are intermediate between those of the homopolymers; the glass transition temperatures of the copolymers decrease as the DMP content increases. This is due to the weakening of the cohesive force between PET molecular chains by the incorporation of DMP units.

As shown in Table 2, each copolymer exhibits a single glass transition temperature. The variation of [T.sub.g] with composition was analyzed by the Fox (21) and the Gordon-Taylor equations (22):

1/[T.sub.g] = [[w.sub.1]/[T.sub.g1]] + [[w.sub.2]/[T.sub.g2]] (6)

[T.sub.g] = [[[w.sub.1][T.sub.g1] + k[w.sub.2][T.sub.g2]]/[[w.sub.1] + k[w.sub.2]]] (7)

where [T.sub.g], [T.sub.g1], and [T.sub.g2] are the glass transition temperatures of the PEPT copolymer, the PET homopolymer. and the PEP homopolymer respectively. [w.sub.1] and [w.sub.2] are the weight fractions of the DMT and DMP comonomers respectively. k is an adjustable parameter depending on the cubic expansion coefficient and the specific volume of each component.

The results are displayed in Fig. 3. The glass transition temperatures of the copolymers rich with DMT units are well fitted by both the Fox and the Gordon-Taylor equations; from the fitting with the Gordon-Taylor equation by nonlinear least-squares regression, a value of 0.54 was obtained for k. In contrast, the glass transition temperatures of the copolymers rich in DMP units show large deviations from those predicted by the Fox and the Gordon-Taylor equations. In general, the glass transition temperature of a polymer is strongly dependent on its molecular weight; in particular, [T.sub.g] varies very sensitively with molecular weight for polymers with a [bar.M.sub.w] of less than 20,000 (23). Thus the large deviations might be due to the relatively low molecular weights of the DMP-rich copolymers, compared to the high molecular weights of the DMT-rich polymers. Collectively, these [T.sub.g] analyses lead to the conclusion that the DMT-rich copolymers are random copolymers.


Melting Behavior and Equilibrium Melting Temperatures

The DSC measurements were extended to the polymer samples (PET homopolymer, DMP6 and DMP11 copolymer) that had been crystallized isothermally for 40 min at the chosen temperatures in the range 190[degrees]C-226[degrees]C. The resulting DSC thermograms are displayed in Fig. 4. Regardless of their crystallization temperatures ([T.sub.c]), the PET samples crystallized at temperatures below 220[degrees]C show multiple melting behavior, as reported previously (10, 24-27), each exhibiting three melting peaks. The melting peak (peak A) in the low-temperature region is very broad, and shifts to the high-temperature region as [T.sub.c] increases. Thus this melting peak is assigned to the melting of crystals formed secondarily in the crystallization. On the other hand, peak B in the intermediate temperature region is less broad. As [T.sub.c] increases, this peak becomes more intense and narrower, and gradually shifts to the high-temperature region. Peak B is thus assigned to the melting of primarily formed lamellar crystals. The third melting peak (peak C) is in the high-temperature region. The position of this peak varies very little with the crystallization history. However, this peak becomes weaker as [T.sub.c] increases. This behavior is in contrast to that of peak B (owing to the melting of primarily formed lamellar crystals). This fact suggests that the population of crystals corresponding to peak C depends on that of the primarily formed crystals. Thus the crystals that correspond to peak C form in small numbers when the population of primarily formed crystals is large. From these considerations, peak C is assigned to the melting of crystals recrystallized during the subsequent heating run after crystallization at a chosen [T.sub.c]. In contrast, the PET samples crystallized at temperatures [greater than or equal to] 220[degrees]C exhibit only two melting peaks (A and B), which correspond to the meltings of the secondarily and primarily formed crystals. This suggests that recrystallization during the subsequent heating run is much more favorable in polymer samples crystallized at a higher degree of supercooling. Similar multiple melting behaviors are observed in the DMP6 and DMP11 polymers, as shown in Fig. 4.


To determine the equilibrium melting temperature [T[degrees].sub.m] of each selected polymer, the melting temperatures of the primarily formed crystals were determined from the DSC thermograms shown in Fig. 4 and then plotted with respect to [T.sub.c] in accordance with the Hoffman-Weeks method (28):

[T.sub.m] = [T[degrees].sub.m] (1 - 1/[gamma]) + [T.sub.c]/[gamma] (8)

where [gamma] is a parameter that depends on the lamellar thickness. More precisely, [gamma] = l/l*, where l and l* are the thicknesses of the grown crystallite and of the critical crystalline nucleus respectively. As shown in Fig. 5. [T[degrees].sub.m] was determined to be 272.8[degrees]C for the PET homopolymer, 266.4[degrees]C for the DMP6 copolymer, and 262.2[degrees]C for the DMP11 copolymer. This value of [T[degrees].sub.m] for the PET polymer is in good agreement with those (275.4[degrees]C-278[degrees]C) reported previously (10, 29).



The [T[degrees].sub.m] of the copolymers decreases with increasing DMP content. The estimated equilibrium melting temperatures were further analyzed by the Flory equation (29):

[1/[T[degrees].sub.m,co]] - [1/[T[degrees].sub.m]] = - [R/[[DELTA][H.sub.f]]]ln x (9)

where [T[degrees].sub.m] is the equilibrium melting temperature of the PET homopolymer, x is the mole fraction of the crystallizable ethylene terephthalate units, and R is the universal gas constant. The results are shown in Fig. 6. As shown in the figure, the equilibrium melting temperatures follow the Flory equation. This result confirms that both DMP6 and DMP11 polymers are random copolymers, as was concluded above from the variation of the glass transition temperature with composition. Further, this suggests that the DMP units incorporated as minor components of the backbone are excluded from the lamellar crystal formation of the major DMT component.

Non-isothermal Crystallizations

Non-isothermal crystallizations of the crystallizable PET homopolymer and copolymers DMP6, DMP11, and DMP18 were performed from the melt by DSC at various cooling rates in the range 2.5[degrees]C-20.0[degrees]C/min. A representative selection of the resulting DSC thermograms is shown in Fig. 7. The PET homopolymer and copolymers exhibit crystallization exotherms typical of common crystallizable polymers. The homopolymers show narrower exothermic peaks than the copolymers. For each polymer the exothermic peak becomes broader and shifts to the low-temperature region as the cooling rate increases. When the crystallizations of the crystallizable polymers are carried out at the same cooling rate, the crystallization exothermic peak shifts to the low-temperature region as the DMP content increases. This effect might be due to a reduction in the regularity of the chains of the crystallizable polymers as their DMP content increases. From the DSC thermograms, we estimated the crystallization peak temperatures [T.sub.p], heats of crystallization [DELTA][H.sub.c], and crystallinities [X.sub.c]. These results are listed in Table 3. For all the polymers, a rapid cooling rate causes a decrease in [T.sub.p] and a reduction in both [DELTA][H.sub.c] and [X.sub.c]. [T.sub.p], [DELTA][H.sub.c], and [X.sub.c] also decrease as the DMP content increases.


To understand the crystallization kinetics under non-isothermal conditions, both the modified Avrami (10, 17-20) and/or the Ozawa analyses (30-33) of non-isothermal crystallization exotherms have been used. As a result, it was found that the values of the Avrami exponent obtained in the modified Avrami analysis are always higher than those of the Ozawa exponent even though both of the analyses have been derived from the same Avrami approach. Thus we analyzed the non-isothermal crystallization exotherms obtained in the present study with the modified Avrami method (10, 17-20). In general, the modified Avrami analysis is valid only during primary crystallization. Thus, the modified Avrami method was only applied to the early stages of each DSC exotherm in which primary crystallization occurs.

As shown in Fig. 8, the resulting plots of log[-ln (1 - [X.sub.t])] versus log t are linear in the early stages of the crystallization, indicating that the modified Avrami technique is suitable in this region for the analysis of the non-isothermal crystallizations of the PET homopolymer and its DMP copolymers. For each cooling run, the Avrami parameters were estimated from the slopes and intercepts of the plots of log[-ln(1 - [X.sub.t])] versus log t in Fig. 8. The results are summarized in Table 4. The resulting Avrami exponents n range from 2.7 to 3.8, depending upon the cooling rate and the composition. These values are comparable with those reported previously for the PET homopolymer (10). These results suggest that the non-isothermal crystallizations of the DMP copolymers have the same heterogeneous nucleation and three-dimensional spherulitic growth mechanism as the PET homopolymer. In conclusion, the DMP units incorporated into the PET polymer as a minor component do not significantly influence the growth mechanism.

In addition, the corrected crystallization rate constants k' and the crystallization half-times [t.sub.1/2] were determined. These results are listed in Table 4. As shown in Table 4, the rate constant k' for a given polymer increases with increasing cooling rate, and the crystallization half-time [t.sub.1/2] shortens as the cooling rate increases. On the other hand, for non-isothermal crystallization at a chosen cooling rate, both the rate constant k' and the crystallization half-time [t.sub.1/2] vary only slightly with changes in the DMP content, suggesting that the crystallization mechanisms of the DMP copolymers are identical to that of the PET homopolymer.

Activation Energies of the Non-Isothermal Crystallizations

The activation energies ([E.sub.a]) of crystallization were determined from the crystallization peak temperatures ([T.sub.p]) measured at the chosen cooling rates (C) using the Kissinger method (34):

[[d(ln(C/[T.sub.p.sup.2]))]/[d(1/[T.sub.p])]] = - [E.sub.a]/R (10)

where R is the universal gas constant. Figure 9 shows Kissinger plots of [T.sub.p] versus C for PET and its DMP copolymers; all the Kissinger plots exhibit relatively good linearity. The resulting crystallization activation energies [E.sub.a] are -224.4 kJ/mol for the PET homopolymer, -123.9 kJ/mol for the DMP6 copolymer, -117.2 kJ/mol for the DMP11 copolymer, and -104.6 kJ/mol for the DMP18 copolymer. Note that since the transition of the molten fluid into the crystalline state releases energy, the activation energies of these melt crystallizations are negative. Note that the [E.sub.a] of the PET polymer is almost two times lower than those of the DMP copolymers. The activation energy is correlated to the crystallization rate. Thus, these [E.sub.a] values suggest that the more regular PET polymer chains are more readily packed into an ordered arrangement than the less regular DMP copolymer chains. However, as described in an earlier section, the rate constant k' varies slightly with changes in the DMP content of the polymer.


PET, PEP, and their copolymers were synthesized from DMT, DMP, and EG monomers by conventional melt polycondensation. The compositions and molecular weights were determined: the DMT-rich polymers were obtained with reasonably high molecular weights, whereas the DMP-rich polymers were synthesized with relatively low molecular weights due to restriction of polymer growth by the highly kinked DMP monomer. The thermal properties of the polymers were also measured. The polymers containing DMP in amounts [less than or equal to] 21 mol% were crystallizable, whereas polymers with other compositions were amorphous. All the polymers exhibited single glass transition temperatures. The glass transition temperatures follow the Fox equation as well as the Gordon-Taylor equation, confirming that the DMT-rich polymers are random copolymers. Moreover, the equilibrium melting temperatures [T[degrees].sub.m] determined for the crystallizable polymers (i.e., DMT-rich polymers) follow the Flory equation, again confirming that the DMT-rich polymers are random copolymers.



The non-isothermal crystallization behaviors of the DMT-rich polymers were measured at various cooling rates by DSC and analyzed by the modified Avrami method. The Avrami exponents n were found to range from 2.7 to 3.8, suggesting that the copolymers crystallize via a heterogeneous nucleation and spherulitic growth mechanism, as observed for the PET homopolymer; the incorporation of DMP units as the minor component in the copolymers does not affect the growth mechanism. In addition, the activation energies of the crystallizations were estimated. The DMT-rich copolymers were found to have higher activation energies than the PET homopolymer.
Table 1. Compositions, Intrinsic Viscosities ([[eta]]s), and Weight
Average Molecular Weights ([bar.M.sub.w]) of the Poly(ethylene

 Sample Feed Ratio (a) Polyester Composition (b) [[eta]]
Designation (DMT/DMP) (DMT/DMP) (dL/g)

 PET 100/0 100/0 0.545
 DMP6 90/10 93.6/6.4 0.572
 DMP11 85/15 89.0/11.0 0.551
 DMP18 80/20 82.3/17.7 0.505
 DMP21 70/30 79.3/20.7 0.512
 DMP32 60/40 67.7/32.3 0.340
 DMP43 50/50 56.7/43.3 0.343
 DMP55 40/60 45.5/54.5 0.203
 DMP66 30/70 34.5/65.5 0.156
 DMP78 20/80 22.3/77.7 0.171
 DMP88 10/90 12.1/87.9 0.125
 PEP 0/100 0/100 0.175

Designation [bar.M.sub.w.sup.c]

 PET 36,000
 DMP6 38,000
 DMP11 37,000
 DMP18 32,000
 DMP21 33,000
 DMP32 18,000
 DMP43 18,000
 DMP55 8500
 DMP66 5700
 DMP78 6600
 DMP88 4100
 PEP 6800

(a) Molar ratio of DMP and DMT monomers fed into the polymerization.
(b) Measured by [.sup.1]H-NMR spectroscopy.
(c) Estimated from the measured intrinsic viscosity using the Mark-
Houwink-Sakurada equation with [K.sub.[eta]] = 4.33 X [10.sup.-4] and
a = 0.68 (16).

Table 2. Glass Transition Temperatures ([T.sub.g]), Melting Temperatures
([T.sub.m]), Heats of Fusion ([DELTA][H.sub.f]), and Crystallinities
([X.sub.c]) of the Poly(ethylene Phthalate-co-terephthalate)s.

 Sample [T.sub.g] (a) [T.sub.m] (b) [DELTA][H.sub.f] (b)
Designation ([degrees]C) ([degrees]C) (J/g)

PET 78.5 254.2 41.9
DMP6 73.9 226.6 35.8
DMP11 71.0 214.1 31.3
DMP18 69.2 211.2 29.6
DMP21 66.4 183.4 8.5
DMP32 59.6 -- (d) --
DMP43 55.5 -- --
DMP55 47.0 -- --
DMP66 35.4 -- --
DMP78 33.5 -- --
DMP88 32.3 -- --
PEP 38.2 -- --

 Sample [X.sub.c] (c)
Designation (%)

PET 35.7
DMP6 30.5
DMP11 26.7
DMP18 25.2
DMP21 7.2
DMP32 --
DMP43 --
DMP55 --
DMP66 --
DMP78 --
DMP88 --
PEP --

(a) Measured by DSC with a heating rate of 10.0[degrees]C/min after
quenching from the melt.
(b) Measured by DSC with a heating rate of 5.0[degrees]C/min after
crystallization with a cooling rate of 10.0[degrees]C/min from the
molten state.
(c) Estimated from the measured heat of fusion using the heat of fusion
(117.6 J/g) of the fully crystallized PET polymer.
(d) Not detected.

Table 3. Crystallization Peak Temperatures ([T.sub.p]), Heats of
Crystallization ([DELTA][H.sub.c]), and Crystallinities ([X.sub.c]) of
the Poly(ethylene Phthalate-co-terephthalate)s.

 Sample Cooling Rate [T.sub.p] [DELTA][H.sub.c]
Designation ([degrees]C/min) ([degrees]C) (J/g)

PET 20 193.7 -41.74
 10 203.6 -42.09
 5 212.2 -42.2
 2.5 218.9 -43.95
DMP6 20 171.5 -36.42
 10 185.1 -37.29
 5 194.6 -39.81
 2.5 201.7 -40.75
DMP11 20 160.7 -26.23
 10 171.2 -31.94
 5 180.8 -31.72
 2.5 189.8 -32.59
DMP18 20 155.7 -25.49
 10 169.3 -30.84
 5 180.7 -29.21
 2.5 188.9 -29.69

 Sample [X.sub.c] (a)
Designation (%)

PET 35.49
DMP6 30.97
DMP11 22.30
DMP18 21.68

(a) Estimated from the measured heat of crystallization using the heat
of fusion (117.6 J/g) of the fully crystallized PET polymer.

Table 4. Avrami Parameters of the Poly(ethylene Phthalate-co-

 Cooling Rate ([degrees]C/min)
 Sample Avrami
Designation Parameters (a) 20 10 5 2.5

 n 3.3 3.8 3.1 3.8
 k' 0.98 0.76 0.49 0.08
 [t.sub.1/2] 0.90 0.98 1.12 1.77
 n 2.9 2.9 3.0 3.0
 k' 0.91 0.74 0.42 0.092
 [t.sub.1/2] 0.91 0.98 1.18 1.95
 n 2.7 3.3 3.5 3.0
 k' 0.91 0.66 0.31 0.067
 [t.sub.1/2] 0.90 0.96 1.26 2.18
 n 2.9 2.9 3.2 3.4
 k' 0.91 0.75 0.42 0.086
 [t.sub.1/2] 0.90 0.98 1.15 1.86

(a) n is the Avrami exponent, k' the corrected rate constant of
crystallization, and [t.sub.1/2] the crystallization half-time.


This study was supported by the Center for Integrated Molecular Systems (Korea Science and Engineering Foundation) and by the Ministry of Education (BK21 Program).


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Department of Chemistry, Polymer Research Institute Center for Integrated Molecular Systems, BK21 Program, and Division of Molecular and Life Sciences Pohang University of Science and Technology San 31, Hyoja-dong, Nam-gu, Pohang 790-784, The Republic of Korea

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Author:Lee, Byeongdu; Lee, Jin Won; Lee, Seung Woo; Yoon, Jinhwan; Ree, Moonhor
Publication:Polymer Engineering and Science
Date:Sep 1, 2004
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