Aging Character is t ics of Oriented Poly(Ethy1ene Terephthalate)

SUMIT MUKHERJEE and SALEH A. JABARIN*

Department of Chemical Engineering and the Polymer Institute

College of Engineering The University of Toledo

Toledo, Ohio 43606-3390

The physical aging characteristics of oriented polytethylene terephthalate (PET), have been studied as functions of storage time and temperature below the glass transition temperature (T,) of PET. The free volume relaxation, associated with aging, has been characterized by the enthalpy of relaxation at Tg, as measured by differential scanning calorimetry. The effects of the free volume relaxation on mechanical properties and the mode of failure have been investigated. I t has been determined that a correlation exists between the enthalpy of relaxation and the ductile-to-brittle failure transition. Molecular orientation reduces significantly the enthalpy of relaxation, resulting in the disappearance of the ductile-to-brittle transition when highly oriented samples are aged over time. I t has been estab lished that a minimum amount of orientation is required to reduce or eliminate the effects of PET aging. Molecular orientation has also been found to reduce craze formation when oriented PET is exposed to a stress-cracking medium at a con- stant stress.

INTRODUCTION

hysical aging behavior, ie the gradual relaxation P of polymer chains, is usually observed for poly- mers in their glassy state. Aging of several amor- phous and semi-crystalline glassy polymers has been studied by various authors and has been the subject of reviews, papers, and books. Polymeric materials, when quenched to temperatures below their glass transition temperature, Tg, do not attain thermody- namic equilibrium. These materials can be referred to as super cooled liquids having excess free volume and excess enthalpy. The term excess is with refer- ence to the equilibrium state. The non-equilibrium nature of the glassy state results in physical aging or time dependent behavior.

Tant and Wilkes, in a review article (1) documented the various property changes that accompany the aging of polymers. The various temperature depen- dent properties which undergo sudden changes at the Tg have been found to change with time during aging. It has been observed that density, yield stress, and elastic modulus increase with increasing aging. Decreases have been observed in the tensile impact energy and creep rate. As the aged samples become stiffer and more dense, they have been found to un-

To whom correspondence should be addressed.

POLYMER ENGINEERING AND SCIENCE, JULY 1995, Yo/. 35, No. 14

dergo transitions from ductile to brittle modes of failure (2-4).

Stmik ( S ) studied the creep behavior of various aged polymers. I t was observed that with increased annealing time, the creep rate was reduced due to densiflcation. When plotted on a logarithmic time scale, the tensile compliance was found to be super- imposable on a master curve indicating the time de- pendence and predictability of the relaxation process. Many of the studies done on aging have been directed at predicting the long-term properties and perfor- mance of a plastic product from its short term char- acteristics; hence, the relevance of an aging study lies in predicting the service life of a fabricated product.

Dynamic mechanical loss studies (5, 6) have been performed with a torsion pendulum, over a wide tem- perature range below the Tg of various polymers. Results have shown that damping decreases with annealing time. Changes in the logarithmic decre- ment were found to correlate with changes in the enthalpy of relaxation. Dielectric property studies on annealed samples were also found to correlate well with the enthalpy data.

Minnini, Moore, Fick, and Petrie (7) examined the stress strain behavior of annealed amorphous PET samples. They found a correlation between the amount of free volume trapped in a sample and the mode of tensile failure. Craze formation studies indi- cated that the voids and inhomogeneities of an amor-

1145

Surnit Mukhejee and Saleh A. Jabarin

phous material possibly manifest as crazes in a stress-cracking environment (8).

Tensile studies (5, 9) on various polymers such as polytethylene terephthalate) (PET), polyvinyl chloride (PVC), polymethyl methacrylate (PMMA), polystyrene (PSI, and acrylonitrile butadiene styrene terpolymers (ABS) have shown decreases in impact strength with increasing aging times. Krezewski, Labovitz, and Sieglaff (4) studied the effects of aging injection molded PVC below its T,. The impact strength was observed to decrease with increased aging time. In- creasing the mold temperature resulted in an in- crease in the impact strength.

Tant and Wilkes (6) evaluated the aging of semicry- stalline PET films. Amorphous PET films were ther- mally crystallized and their stress-strain, stress r e laxation, and thermal behaviors were analyzed. It was found that the extent and rate of aging decreased with increasing crystallinity. This was explained by attributing aging to the relaxation process in the amorphous phase. The crystalline phase acts as physical cross-links which retard relaxation. I t was predicted that 100% crystalline material would not undergo the process of physical aging.

Following the structural changes in annealed PET films, Siegman and Turi (3) observed that increasing sub-T, annealing time caused an increase in T, and the crystallization temperature, (Tc), however, no evi- dence of formation of crystalline domains has been observed (10) during the process of physical aging. Spectroscopic analyses (2, 11) of sub-T, annealed PET samples have indicated changes in structural arrangements of the polymer chains toward more random conformations.

Attempts have been made to relate the free volume of relaxation accompanying the glassy state to the relaxation endotherm at the T, and changes in the mechanical properties of the polymer. Matsuoka (12) developed mathematical models to correlate the free volume with the excess enthalpy and excess entropy. The relationships, which incorporate parameters such as thermal history, pressure, and strain rate, help predict the nonlinear viscoelastic behaviors of polymeric solids under stress. Muller and Wendorf (13) compared the volume and relaxation enthalpy of aged glassy polycarbonate samples. Cold drawn sam- ples were observed to have higher rates of relaxation, compared to isotropic amorphous samples. It was suggested that the process of cold drawing increased the free volume resulting in higher rates of relaxation.

Light and Seymour (14) studied the effects of sub-T, relaxations on the gas transport properties of polyes ters. They modified the backbone chain structure of amorphous PET by copolymerization using different acids. This resulted in a stiffer chain backbone, which restricted sub-T, molecular movement. The molecu- lar mobility was measured by calculating the area under the &relaxation peak during dynamic me- chanical testing of the samples. Permeability studies of these samples showed that reduction in p-relaxa- tion peak areas were accompanied by decreasing oxy-

gen and carbon dioxide diffusion coefficients and per- meabilities.

Jabarin and Lofgren (2) investigated the effects of aging and environmental stress cracking on amor- phous unoriented PET samples. The effects of physi- cal aging were distinguished from those of stress cracking conditions. It was found that for PET h e mopolymer and copolymer sheets aged at tempera- tures below the glass transition, free volume relax- ation leads to an increase in the yield stress and a decrease in the percent elongation to break. There appeared to be a critical AH (enthalpy of relaxation value) above which the samples failed in the brittle mode. These critical relaxation values were different for the homopolymer and copolymer samples. Sam- ples exposed to stress-cracking conditions were also found to undergo transitions to brittle modes of fail- ure when they achieved this critical relaxation en- dotherm value. There seemed to be no correlation, however, between craze development and the mode of tensile failure. Samples with greater numbers of crazes did not necessarily undergo brittle failure.

Physical aging has been studied using various types of polymer systems including glassy amorphous poly- mers and semicrystalline polymers (6). The objective of the present work has been to study the effects of aging on oriented PET sheet. I t has been well doc- umented that proper stretching conditions and or- ientation of PET have resulted in a strain induced crystalline morphology (15, 16). Orientation imparts order to both the amorphous and crystalline polymer phases. Increasing orientation results in a decrease in the amorphous content making it significant to determine if increasing orientation reduces the ef- fects of aging. In this study, various PET oriented homopolymer sheets (uniaxially constrained and equibiaxially stretched) were aged at temperatures below T,. Biaxial orientation was performed both se- quentially and simultaneously. Changes in thermal, mechanical, and physical properties of the variously stretched samples were examined as a function pf aging time. The effect of aging temperature on these properties was also studied. An attempt has been made to correlate the mode of tensile failure and the enthalpy of relaxation at the glass transition. This was done in order to estimate the minimum orienta- tion required to reduce the effects of aging. It has been observed that stressed polymer samples develop crazes when placed in contact with a stress-cracking medium. The effect of orientation on craze formation and environmental stress cracking on oriented PET was also studied.

EXPERIMENTAL

Extruded PET homopolymer sheets (9663 grade) provided by Eastman Chemicals Division of Eastman Kodak Co. were used to prepare uniaxially and biaxi- ally oriented samples. The material had an intrinsic viscosity of 0.8 as measured in a 60/40 pheno/tetra- chloroethane solvent (by weight) at 25°C. The corre- sponding weight average molecular weight is 57,000.

1146 POLYMER ENGINEERING AND SCIENCE, JULY 1995, Vol. 35, No. 14

Aging Characteristics of Oriented PolyfEthylene Terephthalate)

Starting samples had negligible initial orientation and crystallinity, as verified by density measure- ments. Samples cut from extruded sheets were condi- tioned for a t least 24 h at 23°C and 50% relative humidity, to give all samples similar moisture con- tents, before they were stretched.

Samples were stretched at 100°C both uniaxially and biaxially using the Long Extensional Tester (L.E.T.), which essentially consists of two mutually perpendicular stretch heads. Cut and conditioned samples of dimension 6.05 cm X 6.05 cm (2.38 in x 2.38 in) were clamped to the stretch heads inside the L.E.T. and maintained at a temperature of 100°C for 1 min in order to attain a uniform temperature throughout. They were then stretched at a relatively fast rate of 200%/s (4 in/s) to prevent any relax- ations in the rubbery stage.

Oriented samples were prepared in the following stretch ratios:

Constrained Uniaxialmode: l . l X 1, 1.2X 1 , 2 X 1,

3 X 1 , 4 X 1. Biaxialmode: 1 . 5 ~ 1.5.2.5~2.5.3.5X3.5.

Unoriented samples (1 X 1) were prepared by anneal- ing the PET sheet at 80°C for 30 min to erase all previous low temperature thermal and stress histo- ries. The samples thus prepared were aged in ovens maintained at 35 or 40°C. These two temperatures correspond to the glassy state of PET as they are below the measured Tg of PET which is around 77°C. Since oriented PET undergoes shrinkage even below the measured Tg, aging temperatures well below the onset of shrinkage (around 55°C) were chosen.

Samples were analyzed in terms of their stress crazing behavior in a solution of 20% Diversey Sure Lube maintained at 30°C. The solution was stirred continuously to maintain uniform temperature throughout. Samples were clamped both at the top and bottom. A pre-determined load of 13.8 MPa was applied to the samples through cables attached to the top of the sample holders. The applied load was much less than the yield stress of the samples. Sample dimensions are the same as that followed in ASTM D 1708.

Thermal properties were measured using a Perkin- Elmer differential scanning calorimeter (DSC-2) at a heating rate of 10"C/min. A nitrogen purge was used to prevent oxidative degradation of the samples. For each determination a single thickness of sample sheet was cut to size and placed in an uncrimped alu- minum sample pan. In this manner baseline instabili- ties, often noted with the use of multiple sample pieces, were avoided in the temperature range near Tg. The relaxation endotherm was taken as the area enclosed by the peak occurring after Tg and a post Tg baseline, corresponding to that of an equivalent un- aged sample. Values for the relaxation endotherm were calculated using Perkin-Elmer standard TADS 3600 computer software and reported as A H(kJ/kg).

Tensile properties were measured in the direction of orientation, using a table model Instron (TM 1 10 1). at room temperature, with a crosshead speed of 0.0423 cm/s (1.90 %/s) , according to ASTM D 1708. For thin biaxially stretched films ASTM D 882 was used. The crosshead speed, in this case, was 0.846 cm/s (16.7 % / s ) for samples having elongation to break above 100% and 0.0846 cm/s (1.67%/s) for elongation to break values of less than 100%.

Density was measured at 25°C according to ASTM D 1505, using a density gradient column prepared from aqueous calcium nitrate solutions. The volume percent crystallinity of PET was calculated from mea- sured densities, assuming the amorphous density to be 1.333 g/cm3 (17) and the density of completely crystalline polymer to be 1.455 g/cm3 (18).

RESULTS AND DISCUSSION

The properties of oriented and unoriented PET sheets, aged at 40 or 35"C, were monitored as func- tions of aging time. Changes in Tg and enthalpies of relaxation were monitored using differential scanning calorimetry. The effect of aging on sample density was also investigated as were two aspects of mechanical properties, specifically the yield stress and the strain- to-break.

Thermal Properties Aging is basically a non-equilibrium behavior. When

polymer samples are heated above their Tg, they at- tain a state of high free volume and energy. If they are subsequently cooled rapidly to a temperature below Tg, the polymer chains are frozen in that state of high free volume and energy. With increasing annealing time the polymer chains relax to their equilibrium lower energy state (19). The phenomena of aging is illustrated in Fig. 1. The enthalpy us. temperature plot of a polymer sample, quickly cooled from above

Fast Cooling / Equilibrium Glassy Slate

I / I

He 1.' Ta Tg

Temperature (T)

Fig. 1. Schematic illustration of the relationship between PET aged in the glassy state and the relaxation endotherm ob tained at Tg , in terms of enthalpy us. temperature changes.

POLYMER ENGlNEERlNG AND SCIENCE, JULY 7 995, Yo/. 35, No. 14 1147

Sumit Mukhejee and Saleh A. Jabarin

Tg, is represented by the thin solid line. The initial enthalpy, after cooling to temperature Ta, is H a If the sample is aged at Ta for a time t, the enthalpy decreases to Ht. This change in enthalpy ( H r ) is equal to Ha-Ht The equilibrium enthalpy (He) that can be attained at Ta is represented in the Agure by the dashed line. With increased time and tempera- ture (below Tg) of aging, Ht approaches He. When a sample aged at Ta is reheated during a DSC run (heavy solid line), the change in enthalpy ( H r ) that occurred during aging is absorbed at Tg. This energy absorption is seen as an endothermic peak at Tg and is known as the relaxation endotherm.

Figure 2 gives typical DSC scans illustrating the effects of various aging times on PET relaxation en- dotherms, observed during subsequent heating at lO"C/min. These samples had been stretched se- quentially to equibiaxial extension ratios of 1.5 x 1.5 prior to aging at 40°C. The curve obtained for the unaged sample shows no peak at Tg. All aged sam- ples, however, do exhibit relaxation endotherms. The areas measured under these peaks are related to the enthalpy of relaxation, reflecting the amounts of re- laxation that occurred during aging. With increasing aging the areas under the peaks increase since sam- ples aged for longer periods of time undergo more relaxation.

Glass transition temperatures have been taken as the midpoints of the extrapolated portions of the glassy and rubbery regions of the DSC scans shown in Fig. 2 and are denoted by the circles. These scans indicate an increase in the Tg with increasing aging times. The unaged sample has a Tg of 78°C. With 41 days of aging this temperature increases to 80°C and after 65 days it is observed to be 81°C. This shift in Tg with increasing sub-Tg aging time results from the decreased free volume related to molecular mobility and indicates the increased temperature at which molecular rearrangement may occur. Similar in- creases have been observed for unoriented PET as well as other amorphous polymers such as polycar-

0.6 -J

Y 0.5

E f 0.4

0

6 0 . 3

E 0 0.2

m Y

.- +

2 0.1 a

0

4 65 days

I , I

40 50 60 70 ao 90 100 110 120

Temperature ("C)

Rg. 2. DSC scans of (1.5 X 1.5) sequentially stretched PET, illustrating changes in the relaxation endothem as a result of various aging times at 40°C.

bonate (PC), polyvinyl chloride (PVC), (20) polyvinyl acetate (PVAc), (21), and polystyrene (PSI, (22).

Figure 3 shows the variation of the relaxation en- dotherm as a function of aging time for uniaxially stretched samples aged at 40°C. The relaxation en- dotherms for the uniaxially stretched samples in- crease with increased aging times. Such increases have been reported for various annealed, unoriented, amorphous polymers (4). After equivalent aging time, the 4 X 1 sample exhibits a lower enthalpy of relax- ation than the 2 x 1 sample. This is due to the orien- tation process, which causes strain induced crystal- lization of the amorphous polymer phase. The 4 X 1 samples have 27% crystallinity compared to 4% in the 2 X 1 samples. A similar trend is observed for the sequentially equibiaxially stretched samples, as seen in Fig. 4. In general, at equivalent aging conditions, a smaller relaxation endothem is seen with increasing sample orientation. This is to be expected as oriented samples have reduced amorphous fractions because of increased crystallinity (observed from density), and the amorphous fraction is that portion of the polymer known to experience the aging phenomena. Levels of relaxation, therefore, show smaller increases with in- creasing orientation. Exceptions to the rule are sam- ples with very low stretch ratios. At low stretch ratios, samples are stretched to levels below that of yield strain (elastic region) resulting in an increase in the amount of residual stress. This stress may result in an increase in the relaxation rate. A case in point is the 1.1 X 1 sample which has a higher enthalpy of relaxation than the 1 X 1 unoriented sample.

At a higher aging temperature the relaxation p r e cess is faster due to greater free volume relaxation and higher segmental mobility of the polymer chains. This is observed in Fig. 5, where uniaxially stretched

1.1 x 1 " I - 0 Y 4 - .

2 x 1

4 x 1

0 , . , . , . , . , . I .

0 25 50 7 5 100 1 2 5 150

Aging Time (Days)

Rg. 3. Variations of the relaxation endotherms obtained for uniaxially stretched PET samples. as functions of aging times at 40°C.

1148 POLYMER ENGlNEERlNGAND SCIENCE, JULY 1995, VoI. 35, No. 14

Aging Characteristics of Oriented PolyfEthylene Terephthalate,)

c 0 U c - W

r 0 .-

2.5 x 2.5 @ 35°C

0 20 4 0 60 80

Aging Time (Days)

Fig. 4. Variations of the relaxation endothem obtained for sequentially stretched equal biaxial PET samples, as func tions of aging times at 35 and 40°C.

I I

0 2 5 50 7 5

Aging Time (Days)

Flg. 5. Comparisons of relaxation endothem obtained for unoriented and uniaxially stretched PET samples aged at 35 and 40°C.

PET samples aged at 40°C have larger relaxation en- dotherms after a given number of days compared to similarly stretched samples aged at 35°C. Similar trends have been reported (22) for polystyrene sam- ples aged at different temperatures. The effect of tem- perature is similar for the biaxially stretched samples as shown in Q. 4.

A comparison of the relaxation endotherms was made for equibiaxially sequentially and simultane- ously stretched samples. For the same stretch ratio and the same aging temperature, there was not much difference between the relaxation endotherms o h

served for the sequentially us. simultaneously stretched samples. This is shown in Fig. 6. The relax- ation process, therefore, appears to be independent of the mode of stretching for the ratios and conditions investigated.

From the above plots it is evident that the rate of stress relaxation is high initially and decreases with time. Highly oriented, uniaxially (4 X 1) and biaxially (2.5 x 2.5) stretched samples show relatively smaller changes in their enthalpy of relaxation values than their less oriented counterparts.

Figure 7 shows the variation of density with planar extension for both uniaxially and biaxially stretched samples. Both the uniaxially and biaxially oriented samples initially show no appreciable change in den- sity. Beyond a planar extension of two, the densities increase sharply and finally level off. Samples with planar extensions close to two have high amorphous fractions (- 96%) and hence are expected to show greater aging behavior than more highly oriented, more crystalline samples. The curve for the uniaxially stretched samples has a steeper rise than that o h tained for the biaxially stretched samples. The right side axis of Fig. 7 gives the volume percentage crys- tallinities corresponding to the density values.

The decreasing free volume of the polymer with increasing aging time causes the gradual densifica- tion of the samples. Frgure 8 shows the change in density with aging for biaxially stretched samples. Similar trends are seen for the uniaxially stretched samples. For samples with low levels of orientation, the curves show an initial increase in density fol- lowed by a gradual leveling off. More highly oriented samples show no change in density with increasing aging time. Since the change in free volume due to chain relaxation at sub-T, temperatures is very small,

5

1.5 x 1.5 : Simultaneous cn $ 4

E 5.

Q

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- 1

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2.5 x 2.5 : Sequential

9 0 0 -

0 2.5 x 2.5 : Simultaneous

0 20 40 60 80 100 120

Aging Time (Days)

Flg. 6. Comparisons of relaxation endotherms obtained for sequentially and simultaneously stretched equal biaxial PET samples aged at 40°C.

POLYMER ENGINEERING AND SCIENCE, JULY 1995, Vol. 35, No. 14 1149

Surnit Mukherjee and Saleh A. Jabarin

1.37 . . . . I . . . I " . 1 . " 1 ' " 1 . ' . I . . ' 4 0

1.365 i'-- uniaxiallv slratchad

1.36 - " . 5 1.355

.? 1.35

- C

1.345

1.34

1.335

1 3 5 7 9 1 1 13 15

Planar Extension

Fig. 7. Variations in density and crystallinity values, ob tained for uniaxially and equibiaxially stretched samples, as functions of planar extension.

1.365

1.360

.g 1.350 i2 B

1.345

A 2.5x2.5

0 3.5x3.5

z 1.340

1 . 3 3 5 4 . I . I . I . I .

0 30 60 90 120 1

Aging Time (Days) 3

Fg. 8. Densities of the biaxially stretched PET samples, as functions of aging times at 40°C.

the observed changes in density are not large. They are of the order of 0.1 %. Changes of the same order of magnitude have been reported in volume relaxation studies on PS and rigid PVC (1).

Mechanical Properties

It is generally recognized that changes in mechani- cal properties can occur as a result of aging polymer samples below their glass transition temperatures. With increasing aging time, the decrease in free vol- ume causes the polymer to become more compact and results in an increase in the yield stress (2,9,20, 23) values. For both uniaxially and biaxially oriented samples, yield stress increases with increasing stretch ratios. This is because of the greater align- ment of the polymer chains in the orientation direc- tion, resulting in enhanced mechanical properties in

1150

that direction. Figure 9 shows that yield stress in- creases with increasing aging time for uniaxially stretched samples aged at 40°C. Very highly oriented samples (5 X 1) do not however, show any appreciable change in their yield stress values, consistent with their relaxation endotherm and density values. This is due to the fact that strain induced crystallization, which occurs during orientation, converts some of the amorphous content of the polymer into a crystal- lized phase. At a lower aging temperature, the relax- ation process is slower, hence less increase in yield stress is observed after equivalent aging times. Figure 10 gives examples of this behavior recorded for sam- ples with orientation ratios of 2 x 1 and 4 x 1, aged at 35 and 40°C.

The reduction in free volume and the increase in enthalpy of relaxation, with increasing aging time, causes the onset of brittle mode of failure during mechanical testing. This behavior has been reported for W C (20) and rubber modified epoxy systems (23). The decreasing polymer free volume allows less space for the dissipation of the applied stress and results in brittle failure or failure during which elongation to break occurs at or below the yield point (2). This can be seen in Fig. 11, where the strain to break is plotted against aging time.

The elongation to break values plotted in the Fig ures are averages of brittle and ductile modes of failure and do not reflect a decreasing elongation to break for individual samples. From the plot, it can be seen that elongation to break values decrease less rapidly for 1.2 X 1 stretched samples than for unori- ented (1 x 1) samples. Brittle failure has not been observed in highly oriented 3 x 1 and 4 x 1 samples as that may require more volume relaxation, which would require longer aging times and corresponding higher relaxation enthalpies than were achieved dur-

200

3 150 n I Y

In In e

E P)

>

100

.-

50

0

T T 5 x 1

4 x 1 ..

&-++--3x1 2 x 1

, . , . I

25 5 0 7 5 1 DO 125

Aging Time (Days)

Fig. 9. Yield stress values, obtained for uniaxially stretched PET samples, asfunctions of aging times at 40°C.

POLYMER ENGINEERING A N D SCIENCE, JULY 1995, Vol . 35, N o . 14

Aging Characteristics of Oriented PolyfEthylene Terephthalate)

150

-125 m P I Y

u) u) z 5 100

E al > -

75

50

4 x 1 @ 40°C 2 4 x 1 0 35%

,z x 1@ 40'C

€ 2 x 1 fa 35°C

I 25 50 75 100 125 150 175

Aging Time (Days)

Fg. 10. Comparisons of yield stress values, obtained for uniaxially stretched PET samples, asfunctions of aging times at 35 and 40°C.

0 1 . , . , . , . , \ . , . 0 25 5 0 7 5 100 125

Aging T h e (Days)

Flg. 1 1 . Changes in elongation to break values, obtained for uniaxially stretched PET samples, asfunctions of aging tlmes at 40°C.

ing these experimental conditions. Hence such orien- tation levels would be suitable for packaging applica- tions requiring storage for prolonged periods. For the biaxially stretched samples also, brittle failure occurs at low levels of orientation as seen in Fig. 12. The 2.5 x 2.5 and 3.5 x 3.5 samples do not exhibit brittle failure during tensile testing. I t should be noted that measured mechanical properties are sensitive to the strain rate and temperature utilized during testing. In the present study, the mechanical properties were measured at room temperature with strain rates from

400

h

300

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iii 100

0

2.5 x 2.5

A

5 k

3.5 x 3.5 , . I . , . , . ,

2 5 5 0 7 5 100 125 150

Aging Time (Days)

Flg. 12. Changes in elongation to break values, obtained for equibiaxially stretched PET samples, as functions of aging times at 40°C.

1.67 to 16.7%/s. The measured percent elongation to break would be expected to decrease if much higher strain rates were employed during testing.

A previous study (2) successfully correlated the onset of brittle tensile failure with the relaxation en- dotherm, present at the glass transition temperature of PET. Unoriented, amorphous PET homopolymer samples were found to exhibit onsets of brittle failure corresponding to enthalpy of relaxation values of 3 kJ/kg. Results of current investigations, utilizing ori- ented PET samples, indicate similar behaviors. Fig ures 13 and 14 show normalized average elongation to break values, plotted as functions of relaxation endotherms, obtained for uniaxially and biaxially ori- ented aged samples. Transitions from ductile to brit- tle characteristics are noted for samples with en- dotherms of about 3 kJ/kg, with almost all samples exhibiting brittle failures as endotherms approach 4

Average elongation to break values shown in the above Figures have been normalized to a scale of one hundred, with this value indicating that all samples failed in a ductile (non-brittle) manner. The transition zone from ductile to brittle modes of failure appears to have been broadened as a result of increased sam- ple orientation. Figure 13 shows this relationship for uniaxially stretched samples. As can be seen, sam- ples with enthalpy of relaxation values below 3 kJ/kg do not generally exhibit onsets of brittle failure. The 3 x 1 and 4 x 1 samples have enthalpy of relaxation values less than 3 kJ/kg and hence have ordinate values of 100. Thus, samples with higher stretch ratios are not expected to attain complete brittle characteristics until they have been aged for much longer times than similar samples with lower stretch ratios. Samples aged at a lower temperature will re-

kJ/Q

POLYMER ENGINEERING AND SCIENCE, JULY 1995, Vol. 35, No. 14 1151

Sumit Mukherjee and Saleh A. Jabarin

1 oa

80

60

40

20

C

1 x 1

1.1 x 1

1.2 x 1

2 x 1

3 x 1

4 x 1

0.0 1 .o 2.0 3.0 4.0 5.0

Relaxation Endotherm (kJlkg)

Fig. 13. Normalized average ebngation to break values, ob- tained for uniaxially stretched PET samples, as .functions of relaxation endothem.

2.5 x 2.5

1.5 x 1.5

2.5 x 2.5

1.5 x 1.5

0 1 2 3 4 5

Relaxation Endotherm (kJlkg)

Flg. 14. Normalized average elongation to break values, ob tained for equibiaxially stretched PET samples, as functions of relaxation endothem.

quire longer times to attain brittle characteristics than equivalent samples aged at higher tempera- tures.

The annealing time, after which the onset of brittle mode of failure occurs, is an important parameter for various application of PET and many other polymers. The Williams-Landel-Ferry equation (24-27) is one semi-empirical equation which has been successfully found to describe the temperature dependence of the relaxation process. Although this equation has been most frequently employed to explain relaxation p r e cesses above the glass transition temperature, it has

sometimes found application at lower temperatures; indeed down to temperatures of Tg - 52°C (1, 28).

The original equation has the form

where c1 and c, are constants dependent on the reference temperature, Trep chosen. The parameters, T and T~~~ are measures of the relaxation time corre- sponding to T and Tref respectively. For the above elongation to break data, by choosing T as represen- tative of the time required for the onset of brittle mode of failure and using a reference temperature of 40°C. the constants c1 and c, are determined to be 1.1 16 and 19.57 respectively. A time-temperature su- perposition fit using the WLF equation helps predict the failure mode at different temperatures after vari- ous periods of aging. The time required for the onset of brittle mode of failure for unoriented samples an- nealed at 40°C is 24 days compared to 101 days for samples annealed at 35°C. The predicted curve for an annealing temperature of 65°C is found to be close to that of the actual data obtained in a previous study (2). This is shown in Fig. 15. By extending the above results to aging conditions at 30°C, the transition to brittle mode of failure is predicted to be around 350 days. At room temperature the transition to brittle characteristics is expected to take several years.

Using an Arrhenius type expression for the time to embrittlement, an activation energy for the onset of brittle failure can be calculated by plotting the loga- rithmic of the time to embrittlement as a function of the reciprocal temperature as shown in Fig. 16. The calculated activation energy of 135.6 kJ/mol is com- parable to the value of 167.4 kJ/mol arrived at in a previous work (7). The slight difference may be due to

400 1 n 6

300 X m e 0

c 0

m m c 0

c. 200

.- c.

i 100

0 0 5 0 1 0 0 1 5 0 2 0 0

Aging Time (Days)

Q. 15. Tim-temperature superposition of elongation to break data obtained for unoriented PET samples.

1152 POLYMER ENGINEERING AND SCIENCE, JULY 1995, Val. 35, NO. 14

Aging Characteris tics of Oriented Poly(Ethy1ene Terephthalatd

Activation Energy = 135 kJ/rnol

Point

/ 00029 0.0030 00031 0 0032 00033 0 14

l/(Temperature) , ( K - I )

Flg. 16. Semilogarithmic plot of the time to embrittlement CIS

afunction of the reciprocal of temperature, for unoriented PET samples.

different molecular weight PET used and other exper- imental variables.

The volume relaxation process was also evaluated in relationship to stress-cracking medium exposure. Stretched PET samples were exposed to a stress- cracking solution of 20% Diversey Sure lube and maintained at a constant stress of 13.79 MPa, which is below the yield stress of these stretched samples. The stress-cracking medium was maintained at 30°C. Over a period of three weeks the craze development was monitored. There was wide sample-to-sample variation in the number of crazes formed. However, in general it was observed that at equivalent exposure times, the number of crazes formed decreased with increasing orientation as shown in Fig. 1 7. After three weeks the samples were removed from the stress- cracking apparatus, washed with deionized water, dried, conditioned, and subjected to mechanical test- ing. As reported previously (2). the samples having greater numbers of crazes did not necessarily fail in the brittle mode. In Fig. 18, the percent brittle failure is plotted against the planar extension. With increas- ing planar extension, the percent brittle failure is observed to decrease, indicating that increasing ori- entation possibly reduces the free volume relaxation and inhomogeneities in the samples. Thus, orienta- tion has been found to reduce the percentages of brittle failures as well as the formation of crazes among samples exposed to stress cracking condi- tions.

CONCLUSIONS

The effect of physical aging and environmental stress-cracking on oriented PET has been investi- gated. In addition to imparting strain induced crys- tallinity, the process of orientation leads to the order-

6 . /

0 5 10 1 5 2 0

Exposure Time (Days)

FYg. 17. Dependence of solution stress craze development upon sample stretch ratio, monitored asfunction of stressed exposure time at 30°C.

20 - biaxial

1 .o 1.5 2.0

Planar Extension 5

FYg. 18. Percent brittle failure us. planar extension behauior, recorded for samples after removal from three weeks oj exposure to stress cracking conditions.

ing of both the amorphous and crystalline phases. I t has been shown that levels of orientation have a major impact on the relaxation process.

In summary, the following conclusions can be made:

1. Both oriented and unoriented samples undergo free volume reduction upon physical aging, as re- flected by relaxation endotherms obtained, while samples are heated in a DSC. The rates and ex- tents of relaxation, decrease with increasing orien-

POLYMER ENGlNEERlNG AND SCIENCE, JULY 1995, Vol. 35, No. 14 1153

Surnit Mukhejee and Saleh A. Jabarin

tation. Reducing the amorphous content in the polymer as a result of orientation, reduces the rate of physical aging, as reported in other studies per- formed using thermally crystallized PET.

2. Decreasing the annealing temperature decreases the rate of relaxation.

3. Changes in the relaxation endotherm, yield stress, and density as a result of aging are reduced by increased degree of orientation.

4. The onset of brittle failure for oriented samples is observed near an enthalpy of relaxation value of 3 kJ/kg, as reported for unoriented sheets. Above relaxation endotherm values of 3 kJ/kg, brittle modes of failure predominate during tensile test- ing.

5. The minimum orientation required to reduce the effects of aging depends not only on the stretch ratio but also the aging temperature and environ- ment.

6. Stress-crazing results show that the rate of craze formation decreases with increasing orientation. The percentage of brittle failure under stress- cracking conditions also decreases with increasing orientation.

REFERENCES

1. M. R. Tant and G . L. Wilkes, Polym Eng. Sci., 21. 874 (198 1).

2. S. A. Jabarin and E. A. Lofgren, Polym Eng. Sci, 32, 146 ( 1992).

3. A. Siegmann and E. Turi. J. Macromol Sci, BlO(4). 689, (1974).

4. R. J. Krzewski, M. Labovitz, and C . L. Siealaff, Macromol Proc. Second Symp., 10. 67 (1978).

5. L. C. E. Stuik, in Physical Aging of Amorphous Polymers and Other Materials, Elsevier, New York and Amsterdam (1978).

6. M. R. Tant and G. L. Wilkes, J. Appl Polym Sci. 26,

7. R M. Mininni, R S. Moore, J. R. Flick, and S. E. B. Petrie,

8. Polymers: An Encyclopedic Source Book of Engineering

9. A. Aref-Azar, F. Biddlestone, J. N. Hay, and R. N. Howard,

10. A. Aref-Azar and J .N . Hay, Polymer, 23. 1129 (1982). 11. R. S. Moore, J. K. OLoanne, and J. C. Shearer, Polgm

12. S. Matsuoka, Polym Eng. Sci, 21, 907 (1981). 13. J. Muller and J. H. Wendorf, J. Polym Sci Polym Lett

14. R R Light and S. W. Seymour, Polym Eng. Sci.. 22, 857

15. S. A. Jabarin, Polym Eng. Sci., 32, 1341 (1992). 16. A. Misra and R Stein, J. Polym Eng. Sci., Polym Phys.

17. E. W. Fischer and S. Fakirov, J. Mater. Sci, 11. 1041

18. C. W. Bunn and R. P. Daubeny, Proc. R. Soc. A, 226.

19. M. J. Richardson and N. G. Savill, Br. Polym J., 11. 123

20. K. H. Illers, Makromol Chem, 127, 1 (1969). 21. H. E. Bair, G. E. Johnson, E. W. Anderson, and S . Mat-

suoka, Polym Eng. Sd, 21, 930 (1981). 22. A. Agrawal, J. Polym Sci Polym Phys. Ed.. 27, 1449

(1989). 23. Z. H. Ophir, J. A. Emerson, and G. L. Wilkes, J. Appl

Phys., 49. 5032 (1978). 24. J. D. Ferry, in Viscoelastic Properties of Polymers. Wiley

& Sons, New York (1971). 25. M. L. Williams, J. Phys. Chem. 39, 95 (1955). 26. M. L. Williams, R F. Landel, and J. D. Ferry, J. A m

27. A. K. Doolittle, J. Appl Phys., 22, 1471 (1951). 28. J . J. Aklonis and W. J. MacKnight, in Introduction to

Polymer Viscoelasticity. 2nd Ed., Ch. 3 & 4, John Wiley and Sons, New York (1983).

2813 (1981).

J. Macromol Sci.. B8, 343 (1973).

Properties, John Wiley and Sons. New York (1987).

Polymer, 24, 1245 (1983).

Eng. Sci, 21. 903 (1981).

Ed.. 26, 421 (1988).

(1982).

Ed, 17, 235 (1979).

(1976).

531 (1954).

(1979).

Chem Soc.. 77, 3701 (1955).

Reuised July 1994

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