Morphology Development during Biaxial Stretching of Polypropylene Films

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1 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 Erlangen, Germany (3) Brückner Maschinenbau GmbH PO box 1161, 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 1

2 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 PP 1 PP 2 Isotacticity* M w (kg/mol) M w / M n T c ( C) Crystallinity (%) T m ( C) Table 1 Properties of the two polypropylene resins. (Melting point, T m, crystallization temperature, T c, 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 (CF 1 ), 54 C (CF 2 ), and 80 C (CF 3 and CF 4 ). In the case of the casting conditions CF 4, 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,, 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 T m and the onset of melting T om were assessed from these measurements. 2

3 A Siemens D500 Diffractometer mounted with a θ-2θ goniometer using the CuK α 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 K β 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 D hkl in the direction normal to the (hkl) crystal plane was then estimated using the Warren approximation and the Scherrer relation[11]: Kλ D hkl = (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 5 MPa PP 1 T s = 150 C / ε = 1s Stress σ N 2 1 =4 (2x2) =9 (3x3) =16 (4x4) =24 (4.9x4.9) =32 (5.7x5.7) MD TD Strain ε N Figure 1 Stress-strain curves for the PP 1 cast film stretched up to different ratios at 150 C with a Hencky strain rate of 1s -1 (l 0 =70mm). 3

4 Figure 1 shows the nominal stress-strain curves for the PP 1 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 X-Ray Crystallitinity Index Error = Cast Area Stretching Ratio 160 C 155 C 150 C 145 C 140 C 70 Figure 2 X-ray crystallinity index versus area stretching ratio for the PP 1 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 4

5 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 ( = 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 MD TD c c Figure 3 Schematic representation of the main crystallite orientation in a biaxially stretched polypropylene films. D khl crystallite size was also calculated from the X-ray measurements. The evolution of the different D hkl 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 D 040 will be discussed. The D 040 crystallite size corresponds to an average size of the crystallites that are oriented parallel to the film plane. Figure 4 shows the evolution of D 040 with stretching ratio in the temperature range between 140 C and 160 C C 155 C 150 C 145 C 140 C D 040 (Å) Error = 5Å 120 Cast Area Stretching Ratio Figure 4 X-Ray mean crystallite size, D 040, versus area stretching ratio for the PP 1 films biaxially stretched at different temperatures. The decrease of D 040 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 5

6 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 deformation direction (010) [001] D 040 Figure 5 Schematic representation of the (010) [001] slip system. Effort was made for estimating the order of magnitude of the temperature effect on D 040. For example, a temperature change of 5 C from 150 C will produce an approximately 8% change in D 040 for a film stretched to an area stretching ratio of 24. Effect of cast film morphology The morphology of the PP 2 cast films was then characterized with DSC, WAXS, and polarized light microscopy. The results are listed in Table 2. Cast Film T m Spherulite Crystal Phase CI WAXS D 110 D 040 ( C) Size (µm) α / β / smectic (%) (Å) (Å) CF unhomog. 57 / 0 / CF / <1 / CF / 2 / CF / 10 / PP < / 0 / Table 2 Morphology characteristics of the different cast films obtained with PP 2 and the reference resin PP 1. 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. PP 1 and PP 2 (CF 1 ), 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. 6

7 MPa 4 3 CF1 CF2 CF3 CF4 Stress σ N in MD T s =150 C λ=4.9x4.9 ε. = 1s Strain ε N Figure 6 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 CF 1 was observed for uniaxial deformation [12]. Yet, this cannot merely be explained by the high amount of smectic phase contained in CF 1. 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 CF 1 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 CF 1 is due to the higher thickness profile of the stretched film obtained for CF 1. 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. 7

8 148 T S =150 C λ=4.9x4.9 ε=1 s =CF 1 ; 2=CF 2 ; 3=CF 3 ; 4=CF 4 Α PP D 040 of Stretched Film % 2 Α 3 4% D 040 of Cast Film Figure 7 D 040 of stretched film versus D 040 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. D 040 of the stretched film increases with increasing cast film crystallinity or D 040 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 CF 1 and CF 4, 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=T m -T stretching < 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, D 040, 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. 8

9 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 D 040 crystallite size of the cast film. The quasi-homogeneous deformation observed for CF 1 at 150 C was attributed to the high amount of thermally unstable crystalline structure present in the initial cast film morphology. 4b) The D 040 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. Magill, J.H., et al., Int. Polym. Process. (1987), Prud'homme, R.E. and C.-P. Lafrance, Polymer (1994) 35, Vries, A.J.d., Pure and Appl. Chem. (1981) 53, Ward, I.M., A.K. Taraiya, and G.A.J. Orchard, J. Appl. Polym. Sci. (1990) 41, Bartczack, Z. and E. Martuscelli, Polymer (1997) 38, Capt, L., M.R. Kamal, and H. Münstedt. in XIII th Int. Congress on Rheology Cambridge, UK. 7. Samuels, R.J., Polym. Eng. Sci. (1979) 19, Flood, J.F., et al., J. Plast. Film & Sheet. (1987) 3, Vries, A.J.d., Pure & Appl. Chem. (1982) 54, Weidinger, A. and P.H. Hermans, Makromol. Chem. (1961) 50, p Warren, B.E., J. Appl. Phys. (1941) 12, Rettenberger, S., et al. in 17 th PPS annual meeting Montreal. 13. Peterlin, A., J. Mater. Sci. (1971) 6, Alexander, L.E., "X-Ray Diffraction Methods in Polymer Science", Wiley, New York, (1969). 15. O'Kane, W.J., R.J. Young, and A.J. Ryan, J. Macromol. Sci.-Phys. (1995) B34, Zannetti, R., et al., Makrom. Chem. (1969) 128, Varga, J., J. Mat. Sci. (1992) 27,