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Morphology Development during Biaxial Stretching of
Polypropylene Films
L.Capt(1)(2), M. R. Kamal(1), H. Münstedt(2), K. Stopperka(3) and J. Sänze(3)
(1) Department of Chemical Engineering, McGill University
3610 University Street, Montreal, Quebec, Canada H3A 2B2
(2) Institute of Polymer Materials , Department of Materials Science,
University Erlangen-Nuremberg, Martensstrasse 7, D-91058 Erlangen, Germany
(3) Brückner Maschinenbau GmbH
PO box 1161, 83309 Siegsdorf, Germany
Abstract
The present work attempts to investigate the biaxial stretching behavior of isotactic polypropylene in
the partly molten state. A novel laboratory biaxial extension device, which simulates closely
commercial biaxial extension equipment, is used. The effect of biaxial extension under various
conditions (effect of temperature and deformation rate) on the morphology and properties of stretched
films is investigated. The effect of the cast film morphology on the biaxial stretching behavior and the
stretched morphology is also discussed.
Introduction
Biaxially oriented films, especially based on polyolefins, represent a major component of the film
packaging industry. The trend towards using wider and faster production lines makes it necessary to
develop new polymeric compositions that resist the stresses encountered during processing without loss
of mechanical and optical properties. Biaxially oriented polypropylene films are mostly produced through
a sequential biaxial stretching process, in which films are cold drawn in two consecutive steps at two
different temperatures. Recently, a novel simultaneous biaxial stretching (LISIM©) process was
developed. This commercial one-step-biaxial-stretching technique is supposed to allow the production of
uniform and highly oriented films at high speed while minimizing energy and production line breaks
occurring during deformation. Moreover, the one-step-stretching takes place in the partly molten state. A
laboratory biaxial stretching device, which simulates closely the above novel commercial biaxial
extension equipment, especially in terms of strain rate, was newly developed.
Biaxially oriented polypropylene films have been already extensively investigated. In previous
studies, biaxially oriented samples were prepared by rolling[1, 2], uniaxial compression, and tenter-frame
stretching[3-5]. However, few studies have reported the use of laboratory equipment that can
simultaneously biaxially stretch samples at high strain rates[6]. Therefore, the present work attempts to
investigate the simultaneous biaxial stretching behavior of isotactic polypropylene (i-PP) by means of the
novel laboratory extension device.
The orientation process can be optimized by varying the starting morphology, stretching
conditions and polymer structure. For isotactic polypropylene, the effects of temperature, draw ratio, and
annealing on the resulting morphology and properties of the uniaxially[4, 7, 8] and sequentially
biaxially[3, 9] oriented films have already been studied. However, there is little information about the
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effect of cast film morphology on deformation behavior, resulting morphology and physical properties of
simultaneously equi-biaxially stretched films.
In this work, wide angle x-ray pole figure measurements were used to yield information about
crystal texture. The thermal behavior and crystallinity of cast and oriented films were studied by means of
differential scanning calorimetry (DSC). Furthermore, wide angle x-ray scattering (WAXS) in reflection
was used to follow the evolution of crystal structure and mean crystallite size.
Experimental
Materials
Two commercial grades of i-PP with a similar molecular structure but a different crystallization
behavior were investigated in this work. Characteristics of both homopolymers are listed in Table 1.
Resin
Isotacticity*
Mw (kg/mol)
Mw / Mn
Tc (°C)
Crystallinity (%)
Tm (°C)
PP1
95.4
330
8.0
117.1
49.6
163.8
PP2
95.5
321
8.8
112.1
46.3
163.9
Table 1
Properties of the two polypropylene resins. (Melting point, Tm, crystallization
temperature, Tc, and Crystallinity were determined by DSC); *determined by dissolution in Xylene
Preparation of the cast films
The two polypropylene resins were extruded through a slit die at 250°C followed by cooling on a
combined chill roll/water bath unit, whose temperature was set to 20°C. Additionally, the cooling
conditions were changed by varying the temperature of the chill-roll and water bath in order to obtain
various cast film morphologies. The following temperatures were chosen: 20°C (CF1), 54°C (CF2), and
80°C (CF3 and CF4). In the case of the casting conditions CF4, the water bath was taken away.
Preparation of the biaxially stretched films
These cast films were then drawn in the partly molten state at temperatures between 140°C and
160°C on a Brückner biaxial stretching device. This machine allows simultaneous and sequential biaxial
stretching along two perpendicular directions using a pantograph system. In particular, the apparatus is
capable to achieve the high strain rates and temperature change seen under process conditions. Before the
stretching process, the samples were rapidly pre-heated (40s) to the desired temperature in a hot-air oven.
Samples were then drawn and finally quickly cooled to room temperature. Forces in two perpendicular
directions, air and sample temperature, and displacement were recorded. Area stretching ratio, λa, strain
rate and temperature were varied.
Measurements
A Dupont-2200 differential scanning calorimter was used to study the melting behavior of
the specimens. The heating rate was 10°C/min. The melting point Tm and the onset of melting Tom were
assessed from these measurements.
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A Siemens D500 Diffractometer mounted with a θ-2θ goniometer using the CuKα
wavelength was utilized to carry out the diffraction measurements in reflection. The apparatus was
equipped with aperture diaphragms of 1° and of 0,15° for the detector diaphragm, respectively. A
monochromator was used to suppress the Kβ reflections.
The presence and amount of the different crystal forms of i-PP (i.e., α-, β- and smectic
phases) were estimated from theses measurements. Crystallinity was also calculated using the method
from Hermans and Weidinger[10]. This method was slightly adapted for the measurements made on
stretched samples. Although calculation of crystallinity index on stretched samples is influenced by the
orientation of the crystallites, it still yields important information about crystallinity when comparing
similarly prepared samples.
Additionally, calculations of crystallite size were made from measurements of the full width at
half maximum ∆(2θ) of the diffraction profiles using a Pseudo-Voigt curve fitting procedures. The mean
crystallite size Dhkl in the direction normal to the (hkl) crystal plane was then estimated using the Warren
approximation and the Scherrer relation[11]:
Dhkl =
Kλ
∆(2θ ) cosθ
where K is a crystallite form coefficient and was taken equal to 1. Although this equation is not
accurate because it neglects the broadening due to lattice distortions, it is sufficient for the present
comparison purposes. Moreover, in the present paper the word “crystallite” will refer to the fundamental
unit of both lamellar and fibrillar structures.
Results
Effect of temperature and stretching ratio
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PP1
-1
Ts = 150°C / ε = 1s
MPa
4
Stress σN
3
2
λa=24
λa=4
(2x2)
1
λa=9
λa=16
(3x3)
λa=32
(5.7x5.7)
(4.9x4.9)
(4x4)
MD
TD
0
Figure 1
0
1
2
Strain εN
3
4
5
Stress-strain curves for the PP1 cast film stretched up to different ratios at 150°C
with a Hencky strain rate of 1s-1 (l0=70mm).
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Figure 1 shows the nominal stress-strain curves for the PP1 cast film samples that were
simultaneously and equi-biaxially stretched at 150°C with a Hencky strain rate of 1s-1 up to different
stretching ratios. Each curve represents an average curve of five experiments. First, it can be seen that the
stretching experiments showed good reproducibility (standard deviation of yield stress = 2%). The small
difference existing between the machine direction (MD) and the transverse direction (TD) stresses is only
due to a mechanical artifact. The same difference between MD and TD was obtained when the cast film
sample was rotated by 90°, which suggested that the difference can not be related to an eventually preoriented cast film.
The load-extension curves in both directions follow the typical ductile deformation behavior - a
yield point followed by strain hardening – well known for semi-crystalline polymers in tensile
experiments. Uniaxial experiments for the same cast film were also carried out under the same drawing
conditions. The effects of temperature and strain rate on uniaxial deformation are qualitatively
comparable to the ones observed for the simultaneous biaxial deformation. Thus the nominal stress level
increased for all strains with either decreasing temperature or increasing strain rate. A homogeneous
deformation (no yield point) at temperatures very close to the melting point was also observed for both
uniaxial and biaxial experiments.
The stress-strain curves of Figure 1 suggest that the residual crystalline phase, together with the
“molten” amorphous phase, undergo the morphological transformation from spherulites to fibrils known
for cold drawing of semi-crystalline polymers[13]. This can be said for the temperature range 140°C –
155°C. Yet the stress-strain curves at 160°C did not exhibit any yield point but a quasi-rubber-like
deformation behavior.
0.80
X-Ray Crystallitinity Index
0.75
0.70
0.65
0.60
160°C
155°C
150°C
145°C
140°C
0.55
0.50
Error = 0.002
0.45
Figure 2
Cast
0
10
20
30
40
Area Stretching Ratio
50
60
70
X-ray crystallinity index versus area stretching ratio for the PP1 films biaxially
stretched at different temperatures
The results from the morphological investigations on the simultaneously equi-biaxially stretched
films are now discussed. The overall crystallinity of the stretched films, as determined by DSC, was not
found to be strongly affected by the stretching temperature and drawing ratio, as already reported[6].
However, crystallinity, as determined by X-rays, shows a clearer dependence on drawing ratio and
temperature, as seen in Figure 2. WAXS measurements seem to be more sensitive than DSC to any
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morphological changes occurring during stretching. Figure 2 shows that the X-ray crystallinity index
increases with increasing stretching ratio and temperature.
From a certain area stretching ratio (λa = 9 – 16, depending on temperature), the plastic flow
behavior starts to be dominated by the deformation of fibrils. Indeed, the lamellar structure present in the
cast film has almost been totally transformed. Furthermore, the dominating texture obtained from pole
figure measurements on the biaxially stretched films was found to be the uniplanar mode (according to
the terminology of Heffelfinger and Burton[14]), in which the majority of the crystallites are
schematically oriented as shown in Figure 3.
ND
b
b
TD
MD
c
Figure 3
c
Schematic representation of the main crystallite orientation in a biaxially stretched
polypropylene films.
Dkhl crystallite size was also calculated from the X-ray measurements. The evolution of the
different Dhkl with drawing ratio for simultaneously biaxially stretched films has already been
investigated[6]. It was shown that the crystallite size decreases with increasing stretching ratio. In this
paper, only D040 will be discussed. The D040 crystallite size corresponds to an average size of the
crystallites that are oriented parallel to the film plane. Figure 4 shows the evolution of D040 with stretching
ratio in the temperature range between 140°C and 160°C.
240
160°C
155°C
150°C
145°C
140°C
220
D040 (Å)
200
180
160
140
Error = 5Å
120
Cast
0
Figure 4
6
12
18
24
30
36
42
Area Stretching Ratio
48
54
60
66
72
X-Ray mean crystallite size, D040, versus area stretching ratio for the PP1 films
biaxially stretched at different temperatures.
The decrease of D040 with increasing ratio suggests that the crystallite size is reduced in a
direction parallel to the b crystallographic axis, as seen in Figure 5. It can thus be deduced that, during
simultaneous biaxial stretching of i-PP in the partly molten state, the (010) [001] crystallographic slip
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system (cf. Fig. 5) is probably dominating the crystal deformation mechanism. This slip system was
already observed to be the dominating one for sequentially drawn films[5].
b
(010) [001]
D040
deformation
direction
Figure 5
Schematic representation of the (010) [001] slip system.
Effort was made for estimating the order of magnitude of the temperature effect on D040. For
example, a temperature change of 5°C from 150°C will produce an approximately 8% change in D040 for
a film stretched to an area stretching ratio of 24.
Effect of cast film morphology
The morphology of the PP2 cast films was then characterized with DSC, WAXS, and polarized
light microscopy. The results are listed in Table 2.
Cast Film
CF1
CF2
CF3
CF4
PP1
Table 2
Tm
(°C)
163.7
164.9
165.5
165.4
164.6
Spherulite Crystal Phase CIWAXS
Size (µm) α / β / smectic (%)
unhomog.
57 / 0 / 43
37.8
17
96 / <1 / 3
48.7
40
98 / 2 / 0
57.0
60
90 / 10 / 0
62.3
<10
100 / 0 / 0
51.1
D110 D040
(Å) (Å)
125 139
160 181
169 187
211 219
165 187
Morphology characteristics of the different cast films obtained with PP2 and the
reference resin PP1.
It should at least be noticed that with increasing the chill roll and water bath temperatures, the
smectic phase content decreased and the β-phase content increased as well as crystallinity and spherulite
size. Additionally, removing the water bath had the effect of reducing the cooling efficiency which
induced a higher degree of crystallinity and spherulite and crystallite sizes.
PP1 and PP2 (CF1), which have a similar molecular structure, were extruded and cast under the
same conditions. However, great differences in all morphological characteristics of the cast film
morphology can be seen in Table 2. This shows the importance of the crystallization kinetic on the cast
film morphology. The difference in crystallization behavior between both i-PP grades is probably due to
the different amounts of residual radicals coming from the peroxide treatment, which was applied to the
resins to improve their processability.
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4
CF1
CF2
CF3
CF4
MPa
Stress σN in MD
3
2
1
0
Figure 6
Ts=150°C
λ=4.9x4.9
-1
.
ε = 1s
0
1
2
Strain εN
3
4
Deformation behavior of the different cast films biaxially stretched at 150°C.
Figure 6 shows the nominal stress-strain curves of these cast films simultaneously equi-biaxially
stretched at 150°C up to an area stretching ratio of 24. The expected increase of yield stress with
increasing crystallinity can clearly be seen in Figure 6. Uniaxial experiments on the same cast films
confirmed this dependence of yield stress on crystallinity. Furthermore, an equivalent correlation between
the yield point and the crystallite size of the initial cast film can also be deduced, in agreement with
uniaxial tests [12]. Besides, the same homogeneous deformation behavior (no yield point) for the cast
film CF1 was observed for uniaxial deformation [12]. Yet, this cannot merely be explained by the high
amount of smectic phase contained in CF1. Indeed, the smectic phase is known to be transformed into the
monoclinic form upon heating above 70°C [15, 16] and was found to be in the melt state at 150°C, as
confirmed with our DSC measurements. Therefore, the homogeneous deformation behavior comes from
the lower level of thermally stable crystalline structure present in CF1 in comparison with the other cast
films. Indeed, the thermally unstable lamellae and the small-range order structure will melt during the
thermal pretreatment. Consequently, the residual crystallinity present in the partly molten state just before
drawing is greatly reduced and leads to a reduced yield stress or even to a homogeneous deformation.
The higher nominal stress level after the yield point region for CF1 is due to the higher thickness
profile of the stretched film obtained for CF1. This difference will probably be erased if the true stresstrue strain curves were plotted.
The DSC and WAXS measurements on these stretched films showed that the film crystallinity is
independent of the cast film morphology. This can be explained by separating the effect of the thermal
treatment from that of the stretching. This was obtained by preparing samples that were preheated to the
drawing temperature but without being stretched. WAXS measurements revealed that crystallinity for all
preheated cast films was within a 2% crystallinity range. In addition, the crystallinity difference between
the stretched film and the preheated sample, which can be related to the stretching effect on crystallinity,
was found to be almost constant for all cast films. Therefore, it can be deduced that the degree of
crystallinity of the stretched film is principally controlled by the thermal treatment applied to the cast film
before stretching, more than by the cast film processing conditions.
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148
TS=150°C
λ=4.9x4.9
-1
ε=1 s
1=CF1; 2=CF2; 3=CF3; 4=CF4
PP1
1
Α
4
D040 of Stretched Film
144
4%
140
Α
3
2
136
5%
132
128
Figure 7
1
130
140
150
160
170
180
190
D040 of Cast Film
200
210
220
230
D040 of stretched film versus D040 of cast film for the different cast films stretched at
150°C up to 4.9x4.9.
Nevertheless, the cast film morphology seems to have an effect on the morphology of the
stretched films. D040 of the stretched film increases with increasing cast film crystallinity or D040 of cast
film, as seen in Figure 7. This suggests that the lamellar structure present in the initial cast film, which
will be transformed into fibrils, affects the size of the crystallite building the fibrils.
However, it should be noted that the main differences are for CF1 and CF4, which contain nonnegligible amounts of smectic and β-phase, respectively. It is well known that the β-phase is
mechanically and thermally unstable and is transformed into α-form between 125°C – 140°C[17]. Under
the present preheating conditions (heating rate >150°C/min), the β-phase and the smectic phase have
probably not enough time to be transformed into a thermally stable α-form and are thus in the melt state
before the stretching. Consequently, it can be concluded that the amount of residual crystallinity and the
lamellar structure still present after the thermal pretreatment and before the stretching seem to be the
decisive morphological properties that will affect the stretching behavior and consequently the resulting
morphology of the biaxially stretched films.
Conclusion
In this work, it was found that:
1 – Simultaneous equi-biaxial stretching of i-PP in the partly molten state follows a ductile deformation
behavior – i.e., yield point followed by strain hardening – for the temperature range 140°C – 155°C. This
suggests that spherulites are being transformed into fibrils according to Peterlin’s model for cold-drawing
of semi-crystalline polymers. At ∆T=Tm-Tstretching < 5°C, i-PP exhibits a quasi-rubber-like deformation
behavior and does not probably follow the same deformation model.
2 – The X-ray crystallinity index of the biaxially stretched films increases with area stretching ratio and
temperature.
3 – The mean crystallite size, D040, was found to decrease with increasing stretching ratio and decreasing
temperature. This suggests that the fibril deformation mechanism is dominated by the (010) [001]
crystallographic system.
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4 – Different cast film processing conditions will affect the cast film morphology, which in turn will
affect the stretching behavior, the morphology, and consequently the properties of the biaxially drawn
films:
4a) – The level of yield stress was found to be dependent on the X-ray crystallinity index and D040
crystallite size of the cast film. The quasi-homogeneous deformation observed for CF1 at 150°C was
attributed to the high amount of thermally unstable crystalline structure present in the initial cast film
morphology.
4b) – The D040 mean crystallite size of biaxially stretched films is significantly affected by the cast film
morphology.
Acknowledgement
The authors wish to thank the Bavarian Research Foundation for financial support.
References
1.
2.
3.
4.
5.
6.
7.
8.
9.
10.
11.
12.
13.
14.
15.
16.
17.
Magill, J.H., et al., Int. Polym. Process. (1987) , 66-76.
Prud'homme, R.E. and C.-P. Lafrance, Polymer (1994) 35, 3927-3935.
Vries, A.J.d., Pure and Appl. Chem. (1981) 53, 1011-1037.
Ward, I.M., A.K. Taraiya, and G.A.J. Orchard, J. Appl. Polym. Sci. (1990) 41, 1659.
Bartczack, Z. and E. Martuscelli, Polymer (1997) 38, 4139-4149.
Capt, L., M.R. Kamal, and H. Münstedt. in XIIIth Int. Congress on Rheology. 2000. Cambridge,
UK.
Samuels, R.J., Polym. Eng. Sci. (1979) 19, 66-76.
Flood, J.F., et al., J. Plast. Film & Sheet. (1987) 3, 171 -197.
Vries, A.J.d., Pure & Appl. Chem. (1982) 54, 647-670.
Weidinger, A. and P.H. Hermans, Makromol. Chem. (1961) 50, p98-115.
Warren, B.E., J. Appl. Phys. (1941) 12, 375-383.
Rettenberger, S., et al. in 17th PPS annual meeting. 2001. Montreal.
Peterlin, A., J. Mater. Sci. (1971) 6, 490-508.
Alexander, L.E., "X-Ray Diffraction Methods in Polymer Science", Wiley, New York, (1969).
O'Kane, W.J., R.J. Young, and A.J. Ryan, J. Macromol. Sci.-Phys. (1995) B34, 427-458.
Zannetti, R., et al., Makrom. Chem. (1969) 128, 137-142.
Varga, J., J. Mat. Sci. (1992) 27, 2557-2579.
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