characterisation of shrinkage in oriented pet films and containers by thermomechanical analysis...

P o l p e r l n r e r n ~ ~ ~ ~ i n a l 3 2 (1993) 325-335

Characterisation of Shrinkage in Oriented PET Films and Containers by -

Thermomechanical Analysis (TMA)

B. Haworth, 2. W. Dong* & P. Davidson Institute of Polymer Technology and Materials Engineering,

University of Technology, Loughborough, Leicestershire LE11 3TU, UK

(Received 14 May 1992; revised version received 1 February 1993; accepted 18 March 1993)

Abstract: In this paper, we report from an extensive series of tests in which the thermal shrinkage behaviour of poly(ethy1ene terephthalate) (PET) homopolymer and copolymer films and containers has been studied by thermomechanical analysis (TMA). Several independent variables have been investigated, and the accuracy of dynamic TMA scans of temperature-dependent shrinkage has been verified by parallel measurements made isothermally. Typical shrinkage levels of biaxially oriented films, of draw ratios between 2.7 and 4.1, are of the order of 5% at 100°C. The magnitude of shrinkage increases with draw ratio, and any evidence to suggest that PET copolymer is marginally more prone to shrinkage is thought to arise from differences in strain-induced crystallinity induced by elongational deformations during processing. Correlations between measurements made on films and commercial containers are good enough to suggest that the simulation technique used is sufficiently valid for further experimentation on processing behaviour.

Key words: poly(ethy1ene terephthalate), molecular orientation, thermal shrinkage, thermomechanical analysis (TMA).

I NTRO DU CTlO N

Poly(ethy1ene terephthalate) (PET) is widely used in the production of synthetic fibres, films and also for blow- moulded plastic products, because of its ability to be oriented by a drawing process and crystallised to yield a high-strength product. In its amorphous state, PET softens above its glass transition temperature (T,) of approximately 70°C to produce a material which is compliant, and highly elastic at low strains. The material generally crystallises quite slowly, except in a relatively restricted temperature range between about 120 and 220°C (the maximum rate of crystallinity occurs midway between T, and T,). Without the drawing process, in which PET is initially oriented, and then undergoes deformation-induced crystallisation at high strain, PET does not reach its full strength potential.

* Present address: Department of Packaging Engineering, Jilin University, Changchun, 130023 People’s Republic of China.

However, oriented PET, when heated above T,, tends to relax, resulting in shrinkage which creates some problems of dimensional instability with certain oriented products. In general, the shrinkage or dimensional instability is undesirable, except for shrink-wrapping in the packaging industry. Because of this inherent lack of dimensional stability at elevated temperatures, it has not been possible to use conventional biaxially oriented PET containers for pasteurised foodstuffs which require filling whilst still hot, or for re-fillable containers which must withstand hot washing processes carried out at temperatures close to the glass transition.

Therefore, there is considerable commercial interest in investigating how the temperature at which shrinkage (elastic recovery of oriented material) or ‘reversion’ occurs in oriented PET containers can be increased.

Conventionally processed biaxially oriented PET containers are dimensionally stable only to temperatures of approximately 60-65°C Cjust below the TJ. The dimensional stability of PET containers has recently been

325 Polymer International 0959-8 103/93/$06.oO 0 1993 SCI. Printed in Great Britain

326 B. Haworth, Z . W. Dong, P . Davidson

improved to permit hot-filling. This has been achieved by a heat-setting treatment,' in the temperature range of 160-220°C. If the relationship between shrinkage (at fill temperature) and heat setting temperature could be established, hot fill PET containers could be designed to allow for a specific amount of shrinkage during the filling operation.

In view of the technological importance of dimensional stability of polymer materials, previous investigations on the shrinkage, or reversion, of poly(viny1 chloride) (PVC), polypropylene (PP), polyethylene (PE) and PET have been With PET, high levels of frozen-in molecular orientation are produced by rapid cooling during processing, such as in thermoforming and stretch blow- moulding. On reheating, these oriented molecular seg- ments in the amorphous region start to relax (or release stored elastic energy) and return to the original random coil state. This results in longitudinal shrinkage, in the direction(s) appropriate to the deformation applied during processing. Therefore, thermal shrinkage is an excellent method of detecting the effects of structural changes in oriented polymer materials and products.

To characterise the processes and structural features giving rise to shrinkage in thermoplastic materials, various alternative measurement techniques have been developed, such as birefringence,'.' wide-angle and small- angle X-ray scattering (WAXS and SAXS),8- l o differential scanning calorimetry (DSC)". " and thermomechanical analysis (TMA).233.6 These techniques have often been used in combination, to reveal the relationships between processing conditions, polymer microstructure and orientation deformation or thermal shrinkage.

The purpose of the work reported in this paper is to investigate the following influences on the thermal shrinkage of PET at elevated temperatures:

-the effect of copolymerisation, in PET materials of relatively high intrinsic viscosity (IV); and

-the effect of stretch ratio, for a given thermomechanical history, in biaxially oriented PET films and containers.

In particular, the authors have attempted to improve upon previously reported methods of shrinkage measurement, by investigating the potential use of TMA in order to characterise the temperature-dependent nature of dilational changes in a single, dynamic scan.

2 BACKGROUND

Several researchers have investigated the structure and properties of oriented polymer fibres and films and some molecular models have been proposed to explain the mechanism of the shrinkage behaviour of oriented polymer^.^^'^- l 6 In order to define and to understand the various factors which influence the thermal shrinkage of polymers, it is helpful to examine the structural changes in an oriented amorphous polymer.

An unoriented amorphous thermoplastic may be pictured as a mass of randomly arranged, intricately intertwined spring-like molecules, frozen in place. When the material is stretched at orientation temperatures, two important changes occur.

The spatial arrangement of the molecules in the material is changed. A rotating force is exerted which pulls or turns the entangled molecules in a direction more nearly approaching the direction in which the material is stressed. The force required for this unravelling of entangled molecules is that needed to overcome the sum of the friction-like forces exerted by side-chain or segment entanglement, and intermolecular attraction. This energy is non-recoverable. The molecules are extended, and the force needed to achieve this deformation is that required to overcome the intrachain forces (that result in coiling) and shape-forming forces (interatomic forces at fixed bonding angles) within the back- bone of the molecules. As in a coiled spring, the resulting strain is an elastic deformation, and, as such, the force is stored as potential energy, and is generally recoverable when molecular strain is relieved.

According to the viscoelastic theory of polymers, the shrinkage on reheating occurs only in a free rotation phase, or in the oriented amorphous region at T > T,, since rotational freedom prevails only in the amorphous phase. The shrinkage depends upon such structural factors as molecular weight, molecular weight distri- bution, crystalline structure and the degree of chain entanglement which affects the orientation in the amorphous phase.

However, recent investigations have demonstrated that shrinkage may occur in both amorphous and crystalline phases of oriented polymers.16 A new model called 'taut- tie-molecules' was first proposed by Peterlin.' On the basis of this model, the shrinkage in oriented crystalline polymers is attributed to the tie molecules existing between small crystallites created by a cleavage of large crystallites during stretching deformation. Because these newly created tie molecules are located in a highly stretched state, the primary shrinkage takes place within the melting temperature region. Molecular interpret- ations of isotropic and oriented crystalline PET mor- phologies have been proposed by Peterlin.'

Another model developed to interpret the shrinkage in oriented semicrystalline polymers is referred to as a 'phase-transition' model, as suggested by Jordon et aL5

In order to understand the mechanism of thermal shrinkage in semicrystalline polymers such as PET, we have to distinguish two basic forms in which PET crystallises: (a) spherulitic crystallisation, and (b) strain- induced crystallisation. Spherulitic crystallisation exists in the form of 'isotropic' aggregates of folded chains in

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Shrinkuge in oriented PETJilms and containers 327

material slowly cooled in the region between T, and the crystalline melting point. It is generally undesirable (often leading to brittleness) and is opaque (due to light scattering effects), but, because it is thermally induced crystallisation, it makes no contribution to shrinkage on reheating. In contrast to spherulitic crystallisation, the strain- induced crystallisation is produced by stretching a previously quenched amorphous material, at a temperature just above T,. These strain-induced crystallites form network aggregates which are usually small, so that the material is still transparent to light but has a higher tensile strength and stiffness than isotropic PET. It is a desirable form of crystallisation for many applications. The network in oriented semicrystalline polymers may be disrupted upon subsequent heating at temperatures appreciably below the ‘normal’ melting point. According to the stress and/or draw ratio, the molecular structure of oriented polymers will change in different ways: if the applied stress is relatively modest, the molecular chain orientates weakly so that the PET remains amorphous. As a larger stress is applied, more network junctions form, so that the oriented molecules constitute a crystalline lattice. The shrinkage in this type of polymer is usually attributed to relaxation of the frozen-in orientation in the amorphous region, when the product is reheated towards a temperature in the vicinity of the glass transition.

In this paper, we have attempted to study the dynamic aspects of shrinkage in biaxially oriented films and containers, using TMA, in a manner which is able to describe both the onset and progression of heat reversion in a single experimental scan. Several TMA variables have been studied, relative to the physical properties of the PET samples, and the effects of copolymerisation and stretch ratio are highlighted.

3 EXPERIMENTAL

3.1 Materials and processing

Two materials were selected for this study. Melinar B90N PET homopolymer, and B95A ‘Laser’ PET copolymer, each of which is a commercial grade of PET manufactured by ICI Chemicals and Polymers Ltd (Wilton, Cleveland, UK) specifically for applications in the stretch blow-moulded containers market. Both these grades of PET have a nominal intrinsic viscosity (IV) of 0.82 0-02 dl/g.

The materials were supplied as quenched (amorphous) isotropic sheets of 500pm average thickness and with a measured intrinsic viscosity of 0.79 0.02 dl/g. TO prepare the biaxially oriented films, 60 mm2 samples were machined from the centre of the sheets, and were marked with lOmm grids using indelible ink, prior to stretching. Bi- axial deformation was induced on a ‘TM Longstretcher’ (at ICI’s Laboratories) at a temperature of 1 10°C using an extension rate of 10 in/s (254 mm/s). The biaxial drawing was simultaneous but not equal; there was a factor of 1-5 between the principal stretch ratios generated (nominally 4.1 and 2.7, in the major and minor axes, respectively). The actual (achieved) draws were deter- mined by measuring the new dimensions of the elements of material from the original (10 mm2) grid pattern.

3.2 Thermomechanical analysis (TMA)

3.2.1 Instrumentation used. A Mettler TMA-40 thermomechanical analyser with a TA processor and film and fibre attachment was used for shrinkage analysis in this study and is shown in Fig. 1. A rapid response in measurement

( A ) ( B )

Fig. 1. Schematic illustration of the mettler TMA-40 Thermomechanical Analyser. (A) Cross-section: a, measuring sensor; b, sample support; c, purge gas inlet; d, LVDT; e, linear motor coil; f, linear motor stator. (B) Film and fibre attachment: a, sample support; b, jaw;

c, sample; d, measuring sensor; e, quartz cover.


328 B. Haworth, Z . W. Dong, P. Davidson

of length is achieved by using a built-in linear variable differential transformer (LVDT). The core of the displacement transducer is directly attached to the measuring probe so that any change in length of the sample gives rise to a corresponding movement of the core relative to the coils, which results in an electrical signal proportional to the displacement. The apparatus shown in Fig. 1(A) is only suitable for bulk samples. With the film and fibre attachment fitted (Fig. l(B)), the TMA-40 can be used for film, foil or fibre samples. In the measurement mode, the sample is lightly tensioned between the sample support and the lower lying measuring probe. The probe force, which is defined by the experimental parameter 'load, acts as a tensile stress to the sample. By suitable choice of load, the sample can either be tensioned very lightly (in a shrinkage experiment) or subjected to a higher stress, as would be the case in a creep analysis.

The TA processor on the TMA unit is used for data manipulation, and to provide information on specific events associated with the experiment (e.g. thermal transitions).

3.2.2 Saniple preparation. Two types of sample were used in this work. One was from biaxially oriented PET films, the other from stretch blow-moulded PET bottles. These two types of sample were made from the same grades of PET material (Section 3.2.1).

For the biaxially oriented PET films, samples of size 15 mm x 6 mm were cut parallel to both draw axes, with a sample punch. After cutting out, the samples were fixed in the jaw, using an assembly jig; the samples had a free length of exactly 10mm. Finally, the samples were attached to the sample support and then to the film and fibre probe in readiness for the test.

For the PET bottle samples, identical specimens to those described above were cut from the sidewalls of as- selected PET bottles, from both the axial and hoop directions. All other preparation procedures were similar to those adopted for preparing film samples, except that a higher load (up to a maximum internal load of 50 g) was applied to the bottle samples when testing, to protect them from excessive distortion due to non-uniform through-thickness shrinkage.

3.2.3 Dj~nmnir and isothermal modes of testing. The TMA-40 allows two modes of experimentation:

(a) measurement of dilation with temperature for a fixed load; or

(b) measurement of dilation with time a t fixed temperature and load.

In the dynamic tests, the chosen temperature scan was between 30 and 230"C, for all the PET oriented film and bottle samples reported in this work. The effect of various heating rates (10, 20 and SO'C/min) on the shrinkage of samples has been studied in the dynamic mode, under a constant applied load of 1 g. Also, the effect of various

applied loads (0.25, 1, 10 and 50g) on the shrinkage of samples has been studied in the dynamic mode, at a constant heating rate of 20°C/min.

For the PET bottle samples, the temperature- dependence of thermal shrinkage has been investigated in the dynamic mode, using heating rates of 10 and 20"C/min, and applied loads of 20 and 50g.

In the isothermal tests, the time-dependence of thermal shrinkage has been investigated for the oriented PET film samples at different temperatures of 80, 100, 120 and 150°C, under a constant applied load of 1 g.

3.3 Differential scanning calorimetry (DSC)

Differential scanning calorimetry is widely applicable to the characterisation of polymer materials. Some thermal analysis experiments were designed as support experiments for the TMA studies and were carried out using a DuPont 200 Thermal Analyser. These experiments included :

Measuring the degree of crystallinity of isotropic PET materials, biaxially-oriented PET films, and PET bottle samples. The crystallinities ( X ) for different PET polymer materials were calculated according to the following relationship:

where A H represents enthalpy changes due to the melting or formation of PET crystallites; AH,,, = (AH,,, - AH,), in which AH,,, is the area under the melting peak and A H , is the area under the cold crystallisation peak; and AHo is the heat of fusion of 100% crystalline polymer. The value of AHo for PET is 126.4J/g.17 Experiments to simulate TMA conditions, using various heating rates: 10, 20 and 50°C/min. Heat/cool/reheat experiments.

During the thermal analysis experiments, a nitrogen blanket was applied to the test chamber, to protect PET samples from oxidative or hydrolytic decomposition.

4 RESULTS AND DISCUSSION

4. I TMA investigations

The thermal shrinkage in oriented PET films and containers has been systematically studied by TMA. The measurement accuracy of TMA, for all samples, was thought to be less than & 10%. Some of the results are reported in the following sections.

4.1.1 The general form of shrinkage in dynamic TMA tests. The general form of the TMA curve in dynamic tests for an ICI B90N (homopolymer) oriented PET film sample is shown in Fig. 2, which was obtained under test conditions

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Shrinkuge in oriented PETj lms and containers 329

Y I I I I

50 100 150 200 Temperature ('C)

Fig. 2. A typical, temperature-dependent shrinkage curve for biaxially oriented PET film (homopolymer, tested parallel to the draw ratio axis of4.1), derived by TMA. Test conditions: heating

rate lOC/min, with an applied load of 1 g.

of 10"C/min heating rate and using an applied load of 1 g. As demonstrated in Fig. 2, the apparent T, is about 85°C for this sample, which is higher than that of isotropic PET material (typically 70-75°C) due to the stretching ratio.I6 Some thermal expansion at T < T, has been observed, and the mean coefficient of linear expansion (a) is between 48 and 61 x 10-'/K, in the range 40-80°C. In the vicinity of the glass transition, the sample starts to shrink significantly; the percentage shrinkage reaches about 30% at 220°C for the sample of draw ratio 4.1.

Two maxima for the temperature-dependent shrinkage rate occur at temperatures of about 105 and 230"C, as shown in the first derivative TMA curves plotted in Fig. 3. The peak immediately above the glass transition temperature is associated with the release of molecular orientation in the amorphous phase.

For an ICI B95A (copolymer) oriented film sample with the same draw ratio, the temperature-dependence of shrinkage is substantially the same as that of the homopolymer sample, and is not illustrated in the paper. The main difference is the slightly higher overall shrinkage in the copolymer, which reaches 32% at 220°C Under the test conditions adopted, however, this is not considered to be a significant difference between the polymers.

10

9

8

7

n 6 < a :

- 5

.* c,

.d s 4 A 3

2

1

0

-1

0

II- r I 50 100 150 200

Temperature ('C)

Fig. 3. Shrinkage derivative versus temperature for biaxially oriented PET homopolymer films (from raw data shown in

Fig. 2).

4.1.2 Effect qf lieuring rutc. The effect of heating rate on the shrinkage of biaxially oriented PET films has been studied in the dynamic mode, at various heating rates between 10 and 5O"C/min. Some results are shown in Figs 4 and 5 for PET homopolymer and copolymer, respectively.

As shown in these figures, the percentage shrinkage in both materials increases at higher heating rates, i.e. the higher the heating rate, the greater the shrinkage observed at any given temperature. The temperature-dependence of shrinkage for samples at the heating rates of 10 and 20"C/min have almost the same form, although the shrinkage is typically 5% greater(between 100 and 200°C) a t the higher heating rate. However, for samples heated at rates of around 50°C/min, anomalous data (implying a much more rapid onset of shrinkage) were obtained, due to instrumental factors. These have not, therefore, been included in the figures.

The temperatures at which the shrinkage rate maxima occur appear to be around 95105°C and 210-230°C for both polymer materials, when tested at heating rates of 20 and 10"C/min, respectively. As is well the thermal shrinkage in drawn polymers generally increases with higher draw ratio. This trend was confirmed in this study, despite the fact that the c-axis of PET crystallites is likely to be parallel to the principal stretch axis.

The fact that the shrinkage varies with heating rate may

POLYMER INTERNATIONAL VOL. 32, NO. 3.1993

3 30 B. Haworth, Z . W. Dong, P . Davidson

40-

3 5 1

- 5 - A *B * C t D

-4 -10 I I I I I

50 100 150 200 Temperature E)

Fig. 4. Effect of heating rate on the temperature-dependent TMA shrinkage bchaviour of biaxially oriented PET homopolymer films (constant applied load, 18). A, Draw ratio 2.7, heating rate 10 C/min; B, draw ratio 4.1, heating rate lO"C/min; C, draw ratio 2.7, heating rate ZO-C/min; D, draw ratio 4.1,

heating rate 20"C/min.

405 35

-4 I" I I I I

50 100 150 200 Temperature ('C)

Fig. 5. Effect of heating rate on the temperature-dependent TMA shrinkage behaviour of biaxially oriented PET copolymer lilms (constant applied load, 1 g). Sample designation (A-D) as

in Fig. 4.

possibly be interpreted by a slow heating rate allowing for simultaneous relaxation of the sample, thereby reducing its shrinkage capacity. However, this argument would lead to the samples heated at 10"C/min showing a greater tendency to develop shrinkage in the earlier stages of the test, for which there is no supporting evidence in Figs 4 and 5. Alternatively, fast heating rates could promote non-uniform temperature distributions and localised melting of small, imperfect crystallites formed by stretching, which might lead to greater shrinkage.

Although the effect of applied load can influence the degree of shrinkage in TMA to a significant extent (Section 4.1.3), it appears unlikely that when using a load of only 1 g (Figs 4 and 5) any significant creep strain occurs. Therefore, notwithstanding the comments made earlier, the authors conclude that the differences in shrinkage attributed to the heating rate may be primarily associated with measurement lag effects in the instrument. It is important, therefore, when making direct comparisons of material behaviour, to ensure that experimental conditions are as closely matched as possible.

4.1.3 Effect of upplied loud. For conventional TMA measurements using the film and fibre attachment, a pre- specified weight is applied to the sample; thus, the film sample is held under tensile stress. The shrinkage behaviour of oriented PET films was therefore investigated under various applied loads. The results are plotted in Figs 6 and 7. In general, the percentage shrinkage decreases with increasing load, for both PET materials. There are, however, competing effects taking place simultaneously in these experiments: when orientation is released (which first occurs at temperatures immediately above the glass transition), shrinkage occurs, which is detected by TMA. However, if the load applied to the attachment is high enough, the shrinkage is counter- acted by an increase in sample dimensions, created by a gravitational tensile load in what is, in effect, a creep test. Since creep rate will also increase with increasing temperature, it is likely that the reliability of the measurement procedure may be affected by the structural changes which first become feasible around the glass transition temperature. In order to achieve shrinkage data which are free from the complicating influence of tensile creep, it appears that applied loads up to log yield accurate and reproducible results, for the PET film samples used in this study.

I t is noted that the creep effect dominates the response in TMA when testing with a maximum internal load of 50 ga t higher temperature. It may be appreciated that the retractive forces in the samples which give rise to shrinkage become too small to balance the tensile stress applied from this load. Therefore, the apparent shrinkage decreases with increasing weight, when creep effects start to dominate the response. The tensile stress may be calculated according to the mass applied, and the specimen dimensions used.


Shrinkage in oriented PET,films and containers 33 1

TABLE 1. Tensile stress applied to film samples for different loads in TMA

Applied load (kg)

Tensile stress (Pa)

0.25 x 10-3 9.08 x 10-3 0.001 3.63 x lo-* 0.01 0.36 0.05 1.82

For the PET film samples studied, the width was 6mm and thickness (typically) 45 pm, so that the tensile stresses applied were calculated and are listed in Table 1 . The variations in stress magnitude will account for the response differences evident from the TMA data in Figs 6 and 7.

In conclusion, the appropriate conditions for the dynamic TMA tests are a heating rate of 10"C/min and an applied load of 1 g, for the oriented PET film samples studied in this work.

4.1.4 Effect of copolymerisation. Only a few works on the effect of PET copolymerisation have been reported previously.6 In principle, the molecular chains in copolymers are less regular than in homopolymers and PET copolymers therefore have slightly depressed transition temperatures, and a reduced tendency to

+ A1 + A2 -A- 61 + 82 -n- Cl -b cz t Dl U D2

60 100 150 200 Temperature (OC)

Fig. 6. Effect of applied load (TMA) on the temperature- dependent shrinkage behaviour of biaxially oriented PET homopolymer films (constant temperature scan rate of 20"C/min). Al, draw ratio 2.7, load 0.25g; A2, draw ratio 4.1, load 0.25 g; BI , draw ratio 2.7, load 1.0 g; B2, draw ratio 4.1, load 1.Og; CI, draw ratio 2.7, load log; C2, draw ratio 4.1, load log; DI, draw ratio 2.7, load 50g; D2, draw ratio 4.1, load 50g,

40

35

%- A1 + A 2

+ 82 *a 4- C2 + Dl -R- 02

-+m

-51 -10 '

I I I 60 100 150 200

Temperature (OC)

Fig. 7. Effect of applied load (TMA) on the temperature- dependent shrinkage behaviour of biaxially oriented PET copolymer films (constant temperature scan rate of 2O"C/min).

Sample designation (Al-D2) as in Fig. 6.

crystallise. It may be expected that the thermal shrinkage in copolymers might be larger than that in homopolymers, for the same average molecular weight material, and the authors' practical data have verified this to some extent. The temperature-dependence of shrinkage in PET homopolymer and copolymer film samples is illustrated in Fig. 8. The T, and the degree of crystallinity measured by DSC for these materials are presented in a later section (Section 4.2.1). In particular, the lesser developed crystallinity for PET copolymer may be significant, in the context of shrinkage.

4.1.5 Isothermal testing. A typical isothermal test curve produced by TMA is shown in Fig. 9. The sample used is oriented PET homopolymer film, with a draw ratio of 2.7 (specimen from the minor draw axis). The constant test temperature was 150°C and the applied load was 1 g. As shown in Fig. 9, the sample shrinks very quickly in the first I0 min of testing, then the dimensional changes tend to an equilibrium level after about 30 min. The percentage shrinkage (10.45%) after 30 min is very close to the final degree of shrinkage (1057%) measured after 60 min.

As there are many different isothermal test temperatures of practical relevance, different equilibrium times should be taken, such that steady-state conditions are always achieved.

The isothermal shrinkage of biaxially oriented PET film has been measured for both polymer materials, and in each principal draw axis, in the temperature range of 80-1 50°C. Comparing the data from isothermal tests with


332 B. Haworth, Z . W. Dong, P. Davidson

-10 -54 50 100 150 zoo

Temperature ('C)

Fig. 8. The effect of copolymerisation on the temperature- dependent TMA shrinkage behaviour of biaxially oriented PET films (applied load 1 g, heating rate IO"C/min). A, Homopoly- mer, draw ratio 2.7; B, homopolymer, draw ratio 4.1; C,

copolymer, draw ratio 2.7; D, copolymer, draw ratio 4.1.

those from dynamic tests we find a high degree of agreement (see Table 2). Therefore, it can be concluded that the dynamic mode o f T M A can be used instead of the more time-consuming isothermal mode, in order to generate shrinkage data across a chosen temperature range (i.e. in a single test) with a high degree of accuracy.

4.1.6 T M A on PET bottle sanzples. In bottle processing, the biaxial orientation of preforms during stretch blow-

14' 12

0 : I I I I I

0 10 20 30 40 50 Time (min)

Fig. 9. A typical isothermal TMA shrinkage curve, for biaxially oriented PET homopolymer of draw ratio 2.7 (applied load 1 g

at a constant temperature 150°C).

moulding is usually conducted a t temperatures above T, but below the crystallisation onset temperature (typically 120°C). The amount of effective orientation achieved in the final bottle depends on the original dimensions of the preform, on the molecular weight and type of PET resin and on the blow-moulding conditions used. The bottle samples used in this study were cut from the sidewalls of PET homopolymer and copolymer bottles which had been manufactured under unspecified conditions, al-

TABLE 2. Comparison of shrinkage data for isothermal tests and dynamic TMA tests

Sample Draw 80°C 100°C 120°C 150°C Test mode ratio

ICI B90N 2.7 -0.30 0.80 7.80 10.7 homopolymer 2.7 -0.34 1.18 7.65 11.2

4.1 -0.17 2.00 11.6 15.5 4.1 -0.25 1.80 10.7 16.3

ICI B95A 2.7 -0.23 2,31 9.4 14.1 copolymer 2.7 -0.34 1.16 7.8 12.0

4.1 -0.08 4.10 11.2 16.1 4.1 -0.14 2.43 11.6 17.5

isothermala Dynamicb

Isothermal Dynamic

Isothermal Dynamic

Isothermal Dynamic

a Isothermal test conditions: applied load 1 g, test duration time 60 min at 1 OO"C, 30 min at 150°C and 120°C. 20 min at 80°C.

Negative sign indicates thermal expansion, i.e. below T,. Dynamic test conditions: heating rate 1 O"C/min, applied load 1 g.


Shrinkage in oriented PET.fi1m.y and containers 333

though the stretching conditions of oriented films (from PET isotropic sheet material) studied in this work were thought to be closely related to commercial manufacturing practice in stretch blow-moulding.

The shrinkage of PET container samples has been measured by TMA in the dynamic mode. Since the bottle samples were appreciably thicker than the films, a higher weight of 50 g was applied to the bottle samples during testing, to prevent them from warping due to non- uniform shrinkage. Typical bottle wall thicknesses lie in the range 0.2Ck0.25 mm, giving tensile stress values of 0.33-0.41 Pa (i.e. very similar levels to those calculated for PET film samples-see Table 1 ) when subjected to a load of 50g in TMA.

A distortion effect often seen with PET bottle samples is thought to be attributed to existence of a significant temperature gradient across the preform wall during the reheat period, which results in different morphological states (hence different shrinkage behaviour) through the thickness of stretch blow-moulded bottle samples.

The results are depicted in Fig. 10, from which it is evident that the magnitude of thermal shrinkage of PET bottles is very similar to that of biaxially oriented films reported earlier in this study. Therefore, it can be concluded that the shrinkage in bottles can be adequately simulated by the stretched films, if similar stretch ratios are used.

However, it is noted that the shrinkage behaviour of

Temperature ('C)

Fig. 10. TMA shrinkage behaviour of PET bottle samples (applied load 50 g, heating rate 10"C/min): homopolymer sample cut parallel to the axial (A) and hoop (B) directions; copolymer sample cut parallel to the axial (C) and hoop (D)

directions.

these two bottle materials is different, relative to each other. For the sample manufactured from PET homopolymer, the shrinkage in the axial direction is higher than that in the hoop direction, but the shrinkage behaviour is reversed for the sample from the PET copolymer bottle. This difference may possibly be due to different conditions used in bottle processing. Further work would be required to substantiate the origin of this effect, involving more closely controlled preform and bottle manufacturing phases.

The effects of heating rate and applied load on the shrinkage of bottles have the same overall trends as those of biaxially oriented films mentioned earlier.

4.2 DSC investigations of oriented PET

4.2.1 Crystallinity. All thermal analysis data, derived at a temperature scan rate of 20"C/min, for isotropic PET polymer and biaxially oriented PET films are listed in Table 3. For the bottle samples, the data were obtained at a scan rate of 10"C/min; these are also listed in Table 3.

These results show that the PET isotropic sheet material contains only 2-3% crystallinity, whereas the crystallinities of oriented PET films and blow moulded bottles range from 30 to 40%. The crystallinity of homopolymer samples is a little higher than that of copolymer, either for film samples or for bottle samples, indicating that the homopolymer crystallises (under strain) to a greater extent, under constant stretching conditions.

4.2.2 DSC simulated TMA conditions. In order to explore the nature of the shrinkage behaviour of the oriented films in TMA tested at various heating rates, a series of DSC scans (at identical heating rates) was carried out. At scan rates of 10 and 20'C/min, the results coincide well with those presented in Table 3, and there were no significant endothermic or exothermic changes detected, up to the onset of melting at around 240°C.

However, for the curves from tests at a scan rate of 50°C/min, an apparent endotherm just above T, has been

TABLE 3. Crystallinities ( X ) in various PET polymer materials

Isotropic ICI B90N 75.7 256.1 Very low ICI B95A 73.6 251.4 z:; } crystallisation

Oriented film ICI B90N 254.0 37.9 Temperature scan ICI B95A 248.5 36.7) 20"C/min

Bottle ICI B90N 250.7 33.8 Temperature scan ICI B95A 250.6 31.91 1 O"C/min


334 B. Haworrh, Z . W. Dong, P. Davidson

observed. Similar effects near to T, in the DSC curves in a polymer have been investigated by some authors previ- o ~ s l y . ~ ~ ~ ' * - ~ ~ An endothermal peak, called a superheating peak, has been observed for an annealed sample, and may occur immediately above T,, if the sample is heated at a rate faster than that at which it was cooled. The endothermal peak has been attributed to a sudden expansion of the polymer as molecules achieve mobility after superheating. In contrast, for a quenched sample, a peak (or depression) at the lower end of the glass transition range occurs, depending on the heating rate.

However, the behaviour in detecting a low melting point endotherm when using a scan rate of SO"C/min is partially attributed to an experimental measurement lag. Since the TMA data at this heating rate were also subject to an instrument-related error, the effect has not been pursued further.

Overall, the shrinkage of PET appears to be concen- trated in two distinct temperature ranges (Fig. 3). Having carried out some thermal analysis experiments, the evidence gained appears to support the assumption that the low-temperature shrinkage of PET (at temperatures just above T,) is due to the relaxation of orientation in the amorphous phase. Further orientation may be effectively locked into the semicrystalline structure by the restraining influence of the crystalline network, and subsequent relaxation may then be dependent upon the onset of melting (i.e. at temperatures above 220°C) or other forms of motion in the crystalline phase.

cooling 159.62 'C I

38.07 J /9

4658J /9 heating

254.32 O C '

Fig. 11. DSC thermogrdms for biaxially oriented PET homopolymer films using heating, cooling and reheating modes, at a

temperature scan rate of 20°C/min.

14887 O C cooling . -._ _______--- ----* .. 190.87 O C

heating W 4 9 J / 9

248.07 'C V I I I L 4 I L

XI 150 250 TEMPERATURE OC 1

Fig. 12. DSC thermograms for biaxially oriented PET copolymer films using heating, cooling and reheating modes, at a

temperature scan rate of 20"C/min.

4.2.3 Hrat/cool/reheat trsts. The results of heat/cool/ reheat tests at a DSC scan rate of 20"C/min for PET oriented films are shown in Figs 11 and 12. For the first heating cycle, the results are in close agreement to those in the simulation experiments described above (see Table 3 or Section 4.2 2).

For the cooling cycle, the crystallisation onset is 181.2 and 190.9"C for PET homopolymer and copolymer, respectively. It is noted that the delayed peak crystallisation temperature (T,) upon cooling is lower than T,,, by about lOO"C, and is 159.6 and 148.9"C for homopolymer and copolymer, respectively. This shows the significant supercooling characteristics, which are a feature of PET in several commercial processes.

For the reheat process, the melting point (T,) is a little lower than that measured in the first heating process. This is due to the very low crystallinity. Rapid quenching in cooling will reduce crystallinity and subsequent reheating will result in additional crystal formation. This is the case observed when reheating copolymer samples (Fig. 12(b)). In fact, the crystallinity of the samples after quenching is less than 3%, so we can conclude that they are virtually amorphous.

5 CONCLUSIONS

(1) TMA is a useful technique to characterise the temperature-dependent nature of dilational changes in oriented plastics, in a single dynamic scan. Comparisons


Shrinkuge in oriented PET,films and containers 335

with isothermal data have verified the accuracy of dynamic TMA scans. The most appropriate conditions for dynamic tests on oriented PET films (of thickness between 50 and 100pm) are a scan rate of 10"C/min and an applied (tensile) load of 1 g, corresponding to a tensile stress of -0.04 Pa.

(2) Changes in heating rate are not significant unless the rate is sufficiently high to induce measurement lag effects, as observed at 50°C/min in the authors' experiments. The overall levels of apparent shrinkage are reduced when significant tensile loads are used with the film and fibre attachment, if the stress is high enough to induce creep effects. Since the creep rate increases at elevated temperature, the reliability of the data may decrease as temperature increases, in circumstances where 'excessive' tensile loads are used in TMA.

(3) TMA can be used to measure thermal expansion coefficients, at temperatures below T,. Once above T,, significant shrinkage occurs due to the relaxation of orientation in the amorphous phase. Although shrinkage occurs continuously as temperature is increased further, the rate of shrinkage increases again above 200"C, which may be due to the release of further entropy-elastic strain which becomes feasible as the restraining crystalline network starts to melt.

(4) The magnitude of shrinkage, at a given measurement temperature, increases with stretch ratio, for both homopolymer and copolymer PET. There is marginal evidence that shrinkage in PETcopolymer may be higher, although the properties of each grade are remarkably similar in the majority of tests. The thermal analysis data have shown that the processing characteristics of homopolymer and copolymer may be different, since the degree of crystallinity measured in the latter material is generally lower; this observation might then be respon- sible for any differences detected subsequently by TMA.

( 5 ) The shrinkage characteristics of samples cut from PET bottles have also been analysed successfully by TMA, although a different load was necessary to allow for increased specimen thickness. The correlation between the shrinkage behaviour of films and bottles was

sufficiently good to conclude that an adequate simulation of commercial processing behaviour can be achieved using the TM Long stretcher.

ACKNOWLEDGEMENTS

The authors are grateful to the Melinar Research and Technical Service Group of Imperial Chemical Industries PLC (ICI) for providing stretching facilities and all Melinar" PET materials. One of the authors (Z.W.D.) would like to express his thanks to Dr D. R. Gabe for permission to study at IPTME, Loughborough Univer- sity of Technology, as part of this research programme.

The authors also acknowledge contributions made by Dr M. Gilbert, Dr D. J. Hitt, Mr R. Owens and Mr J. H. Zhao, for their practical assistance and contributions to technical discussions.

REFERENCES

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