Properties of uniaxially stretched polypropylene films.

Sadeghi, Farhad ; Carreau, Pierre J.

INTRODUCTION

Improving the strength of polymer films through stretching is a well-established technology. Film producers mostly rely on biaxial stretching to strengthen the films in both machine and transverse directions (TD) while the applications for uniaxial stretching (MDO for machine direction orientation) are more limited. Some of the applications are found in polyethylene terephthalate (PET) strapping, polyvinylchloride (PVC) food wraps, fibrous high-density polyethylene (HDPE) ribbons for weaving sacks, breathable hygienic films in diaper liners, self-adhesive labels, and polyolefin packaging and lamination (Schut, 2005). MDO is widely used for polyethylene and polypropylene (PP) films because of their numerous applications in packaging. A MDO unit normally includes a cast extrusion line, a slit die, a big drum roll for cooling, and multiple stretching units in sequence. Stretching is carried out in these units that can operate at different temperatures and stretching ratios. The MDO unit available at the Center for Applied Research on Polymers (CREPEC) contains two stretching zones as presented in Figure 1.

[FIGURE 1 OMITTED]

The MDO unit can be operated either in-line or off-line with extrusion and is controlled via variables such as: DR, drawing speed, drawing times (a film can be stretched many times), drawing temperature, and heat-setting conditions. The distance between the draw rollers was set at 5 cm.

The use of MDO units is growing, mostly for specific polymers. The process usually improves the strength and barrier properties, but there are some disadvantages for the produced films. For example, the MDO process improves the barrier properties of nylon and ethyl vinyl alcohol (EVOH) polymer, but it also makes the films brittle and of very low tear resistance in machine direction (MD) (Schut, 2005). For EVOH, MDO is applied selectively since any biaxial stretching deteriorates the barrier properties.

The stretching is usually performed at high temperature and results in a very highly oriented film that causes anisotropy. Nie et al. (2000) studied the morphology development during uniaxial and biaxial stretching of PP films using atomic force microscopy (AFM). They reported the formation of a fibrillar structure with thicker fibrils in the MD while much thinner fibrils connected the thick fibrils to create an overall strong crystalline network. Dez et al. (2005) investigated the influence of the stretching on the crystallinity and structure of biaxially oriented PP films. Their results confirmed the formation of the oriented crystalline fibrils during stretching. They also observed that the stretching ratio did not have a significant effect on the melting point of the stretched samples. Koike and Cakmak (2004) showed that an initial cross hatched of lamellae (typical morphology of PP films produced under low cooling conditions as shown by Sadeghi et al. (2005)) was firstly broken down into small crystal pieces during stretching and then joined to form microfibrils oriented in the MD. Koike and Cakmak (2004) observed an intermediate shish-kebab-like crystalline structure during stretching. The transformation from a shish-kebab-like into a fibrillar structure took place at a strain of 1.2 and the reported macrofibril size was 3-5 [micro]m in diameter.

Rettenberger et al. (2002) studied the effect of temperature on the stress-strain behaviour during stretching of PP. They found a typical ductile behaviour with a yield point, neck propagation, and strain hardening up to a temperature of 155[degrees]C. At higher temperatures instead of yielding the deformation was a quasi-rubber-like. This behaviour was also reported when the films were stretched at high strain rates (over 750 mm/s). They also observed that the non-homogeneity of the deformation reduced with increasing strain rate.

Srinivas et al. (2003) studied the cold stretching of a metallocene linear low-density polyethylene (LLDPE) and reported that the modulus and tensile strength were almost linearly improved with increasing DR (for DRs above 5). This DR dependence was also observed by Gould (1988) during the cold drawing of some polymers. For nylon and polyester, the tensile strength obeyed a strong linear relationship with DR, while for PP the graph revealed two regions: one with a steep slope in the small DR range of 5-7 followed by a second region with a smaller slope up to a DR of 10.

Bheda and Spruiell (1986) analyzed the light transmission properties of oriented PP films. Samples with higher orientation showed a greater overall light transmission. Bheda and Spruiell (1986) claimed that the light transmission was controlled first by surface roughness and second by the internal morphology of the films. Taraiya et al. (1993) examined the permeability of oriented PP films. They proposed that the primary factor that control the permeation of gas molecules and, hence, the barrier properties was the orientation of the amorphous phase. The flow of gas through the amorphous phase of a polymer film can be presented as P=DS, where P is the permeability coefficient, D is the diffusivity, and S is the gas solubility (adsorption of gas molecules on the film surface). The orientation of the amorphous phase influences both parameters, in particular the diffusivity. Taraiya et al. (1993) showed also that the oxygen permeability was reduced with decreasing draw temperature.

In spite of the importance, few results have been reported on the study of the MDO process in terms of film microstructure and the development of mechanical properties during stretching. The objective of this work was to investigate the effect of different processing conditions on the properties of the films of two different PP resins obtained using an industrial MDO line.

EXPERIMENTAL

Materials

Two different PP grades were used in this study. PDC1272 (PP1272) with a melt flow index (MFI) of 0.8 (230[degrees]C/2.16 kg) is a high molecular weight grade from Basell and PP4612E2 (PP4612) with MFI of 2.8 is an ExxonMobil grade proposed for the production of oriented films. The manufacturers reported these two resins as homopolymers and PP4612 was introduced as a bimodal molecular weight distribution resin. An FTIR test was carried out and no trace of ethylene was found in the absorbance spectrum of both resins.

Rheological Characterization

The rheological behaviour of the resins was determined using an ARES rheometer (TA Instruments, New-Castle, DE). Frequency sweep tests were performed at 230[degrees]C and time sweep tests were carried out for the samples at 230[degrees]C to evaluate the decrease of the elastic modulus with time. The results showed a thermal degradation less than 5% during the time required for running a frequency sweep test. The extensional viscosity was measured via the new SER universal testing platform from Xpansion Instruments (Tallmadge, OH). The model used in our experiment was SER-HV-A01, which is a dual windup extensional rheometer and has been specifically designed for the ARES rheometer platform.

Film Preparation

The filmswere prepared using two extrusion film processing lines: The first one was a lab-scale unit equipped with a slit die (20 cm width and 0.19 cm gap) and the extrusion was carried out at 220[degrees]C. To cool the extruded films a fan was installed to supply air to the film surface right at the exit of the die. The cooling was very fast allowing us to obtain a very highly oriented film (see Sadeghi et al., 2007). Since the extrudate velocity at the exit of the die was constant the take-up speed determined the DR for the film production. The second is an industrial MDO line from the Davis Standard Company (Pawcatuck, CT) specifically designed for production of oriented films. A sketch of the MDO line is presented in Figure 1. It includes single and twin screw extruders (we worked with a single screw extruder) equipped with a 122 cm width slit die, cooling drums and a stretching unit composed of two zones each composed of two 1370 mm width rolls. It is possible to operate at two different temperatures and different DRs consequently. The extrusion was carried out at 220[degrees]C and the distance from the die exit to the chill roll was 150 mm. We chose a constant roll temperature of 140[degrees]C for the first stretching zone and operated at different DRs using only the first stretching zone. No cooling was applied at the die exit, and the temperature drop in the extrudate between the die exit and the chill roll was estimated to be negligible.

Film Characterization

Crystal orientation measurements were carried out using a Bruker AXS X-ray (Karlsruhe, Germany) goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper CuK[alpha] radiation ([lambda]=1.542 [Angstrom]) was selected using a graphite crystal monochromator. The sample to detector distance was fixed at 8 cm. Prior to measurements, careful sample preparation was required to get the maximum diffraction intensity. This consisted in stacking several film layers in order to obtain the optimum total thickness of about 2.5 mm. Pole figures of the (1 1 0) and (0 4 0) crystalline reflection planes were measured and the orientation factors determined from the measurements.

Field emission scanning electron microscope (FE-SEM-Hitachi S4700) was employed for observation of the films. The samples were etched with a solution of 60% [H.sub.3]P[O.sub.4] and 40% [H.sub.2]S[O.sub.4] mixed with approximately 0.5 wt.% potassium permanganate. The etching time was 25 min and the samples were rinsed off and washed with a dilute solution of sulphuric acid, hydrogen peroxide, and distilled water.

For FTIR measurements, infrared spectra were recorded on a Nicolet Magna 860 FTIR instrument from Thermo Electron Corp. (Waltham, MA) using A DTGS detector with a resolution 4 [cm.sup.-1] and accumulation of 128 scans. The beam was polarized by means of a Spectra-Tech zinc selenide wire grid polarizer from Thermo Electron Corp. The crystalline and amorphous orientations were measured based on the method explained in detail by Sadeghi et al. (2007).

The differential scanning calorimetry (DSC) tests were carried out on a TA Instruments (New-Castle, DE) Q1000 using a heating rate of 20[degrees]C/min.

The oxygen transmission rates were determined using a modification of ASTM Standard Method D 3985-81 with an Ox-Tran Model 2/21 apparatus (Mocon Inc., Minneapolis, MN) at 25[degrees]C.

Tensile tests were performed using an Instron (Instron Structural Testing Systems Corporation, Novi, MI) 5500R machine. The procedure used was based on the D638-02a ASTM standard.

Haze and clarity were measured in accordance with the procedures specified in the ASTM D 1003-97. The measurements were carried out on a Haze Guard Plus[TM] instrument (Model 4725) made by the BYK-Gardner Inc. (Columbia, MD).

RESULTS AND DISCUSSION

The linear viscoelastic data for both PP resins are reported in Figures 2 and 3. As it is observed from Figure 2, PP1272 showed a larger complex viscosity at the very low frequency range, indicative of a larger zero-shear viscosity compared to that of PP4612.

[FIGURE 2 OMITTED]

[FIGURE 3 OMITTED]

Considering the dependency of the zero-shear viscosity of polymer melts to molecular weight as done by Sadeghi et al. (2007):

[[eta].sub.0] = K[M.sup.3.4.sub.w] (1)

the molecular weight of PP1272 is found to be 1.3 greater than that of PP4612. Since we have calculated the molecular weight of PP4612 in Sadeghi et al. (2007) as 350 kg/mol, the molecular weight for PP1272 is estimated to be 455 kg/mol. Figure 3 shows the weighted relaxation spectra for both resins. The relaxation spectrum is particularly sensitive to the molecular structure and can likely be employed to investigate the presence of high molecular weight chains with long relaxation times. The relaxation spectrum was calculated from the modulus data (G', G'', [omega]) using the NLREG (non-linear-regularization) computer software package of Honerkamp and Weese (1993). As observed in Figure 3 the weighted spectrum for PP4612 is different in shape and slightly broader than that for PP1272. This could be attributed to the bimodal molecular distribution of PP4612.

The transient elongational viscosities for the resins at two different elongational rates measured at 200[degrees]C are presented in Figure 4. Due to experimental difficulties and limitations of the device, no reliable data could be obtained at 230[degrees]C. For low Hencky strain rate of 1 [s.sup.-1] both resins show similar trends, although PP1272 is more viscous. However, for PP4612 slight differences in the curve for the Hencky strain rate of 5 [s.sup.-1] are observed, indicative of a very slight strain hardening, again attributed to its bimodal molecular weight distribution.

[FIGURE 4 OMITTED]

To address the orientation and arrangement of the microstructure of the produced films, wide-angle X-ray diffraction measurements (WAXD) and FTIR experiments were carried out. Figure 5 shows the WAXD diffraction patterns for PP1272 films obtained from two processes (lab unit under high cooling conditions and MDO line).

The crystalline orientation can be analyzed from the pole figures of the (1 1 0) and (0 4 0) reflection planes. The position of the crystalline planes with respect to the machine production direction (MD) is depicted in the sketch of Figure 5e. The normal to the 1 1 0 plane is the bisector of the a and b axes and 0 4 0 is along the normal of the b-axis of unit crystal cells. For the PP1272 film obtained from lab unit under high cooling conditions, the intensity (or orientation) of the 1 1 0 plane is concentrated almost equally in the TD and normal direction (ND), while less intensity (smaller orientation) is observed in the MD direction. For the PP1272 film obtained using the MDO unit the orientation of the 1 1 0 plane is mostly in TD and that of the 0 4 0 plane (b-axis) is mostly in the ND direction. The behaviour of the pole figures for the MDO sample of PP4612 was similar with slightly higher orientation intensities, and for that reason is not presented here. The orientation of the amorphous and crystalline phases (c-axis with respect to MD) of the samples has been measured by Fourier transform infrared spectroscopy and the results are presented in Table 1. The orientation function was calculated as proposed by Lamberti and Brucato (2003):

[f.sub.i,MD] = (D - 1/D + 2) (2)

where D is the ratio of the absorbance in the machine (parallel) to that in the TD. For the orientation function of the crystal phase (c-axis with respect to MD), [f.sub.c], the band 998 [cm.sup.-1] is considered so D will be ([A.sub.[parallel]]/[A.sub.[perpendicular to]])998, where A is the absorbance. To measure the total orientation that includes both the crystalline and the amorphous phase orientations the band at 972 [cm.sup.-1] was selected and [f.sub.av] (average orientation) was calculated based on [([A.sub.[parallel]]/[A.sub.[perpendicular to]]).sub.972]. The orientation of the amorphous phase, [f.sub.a], can be determined from these two values by:

[f.sub.av] = [X.sub.c][f.sub.c] + (1 - [X.sub.c])[f.sub.a] (3)

where [X.sub.c] is the degree of crystallinity, calculated from the DSC results. As expected from the WAXD results, Table 1 clearly shows that the orientation functions of both crystalline and amorphous phases increase with DR. However, as expected the MDO line is more efficient than the lab unit. Similar results have also been reported by Ratta et al. (2001) for a HDPE.

The orientation values for the precursor MDO films cooled under normal conditions (cooling drum) and not stretched (DR=1) are very low, indicative of a spherulitic structure. The important point is about the PP4612 film obtained from the MDO process that shows higher orientation in comparison to PP1272. It is speculated that the effect of the bimodality of the molecular weight distribution (that is represented by a broader relaxation spectrum in Figure 3) is more effective for the MDO process than the molecular weight by itself. At this point we also believe that the slight strain hardening observed in the elongational viscosity compared to PP1272 (see Figure 4) is due to its bimodal molecular weight distribution. This leads to a higher chain orientation in the MDO process where chains are stretched in the semi molten state. This is interesting since in our previous study (Sadeghi et al., 2008) we found that for films obtained from the high cooling lab extrusion unit the most influential parameter for the orientation was the molecular weight.

Figure 6 reports the orientation functions for PP1272 films obtained using the MDO unit as a function of DR. The orientation of the crystalline phase is always above that of the amorphous phase and both dependencies on DR (slope of the curves) reduce with increasing DR. It is worth mentioning that the difference between crystalline and amorphous orientation factors increases with DR.

The DSC results for PP1272 films produced by the MDO process are compared in Figure 7 to those of the lab-scale films and initial PP granules. The results for PP4612 are not reported because of similar trends.

The film prepared using the MDO line shows a much higher melting point (about 10[degrees]C) due to thicker crystals with a small shoulder in lower temperature (bimodal crystal lamellae thickness). This is somehow in contrast to what was reported by Dez et al. (2005) as MDO has a considerable impact on DSC results. For the film sample prepared using the lab unit under high cooling rate a lower (in comparison to the MDO), but sharper melting peak is observed in Figure 7 indicating a better uniformity for the crystal lamellae. Comparing the area under the curves, the crystallinity was increased for the oriented film samples to 49.7% and 49.1% (for the Lab unit and the MDO sample, respectively) compared to the initial polymer granules that was 46.1%.

[FIGURE 5 OMITTED]

To reveal the morphology of the films, SEM micrographs were taken from the film surfaces and presented in Figure 8 (the machine direction is shown in Figure 8a by the arrow). The samples were etched and stretched a little (3%) to have a better quality for the pictures. Figure 8a shows a stacked lamella structure for the film prepared from the lab unit under high cooling rate while for the MDO film sample with the same thickness a fibrillar structure is observed (Figure 8b). The formation of a stacked lamellae structure for the first sample has been studied in our previous paper (Sadeghi et al., 2007). However, for the second sample the observation confirms the results of the literature review (Nie et al., 2000; Koike and Cakmak, 2004). Our understanding is that stretching at a temperature of 140[degrees]C deforms the initial crystalline structure and breaks down most of the lamellae into tiny pieces that consequently will be joined and reoriented to form thick elongated fibrils. This is in accordance with the results on morphology development presented by Koike and Cakmak (2004). The DSC curves for the PP1272 film samples (Figure 7) could be a reflection of this mechanism, with the main peak likely attributed to the elongated crystal fibrils and the shoulder corresponding to the connecting tiny lamellae.

[FIGURE 6 OMITTED]

[FIGURE 7 OMITTED]

Most pronounced differences appear in the tensile response for the samples as shown in Figures 9 and 10. The results for tensile tests made in MD are presented in Figure 9. The PP1272 MDO film samples for DR=1 (i.e., the original precursor film with normal cooling and thickness of 150 [micro]m) and film drawn to a thickness of 35 [micro]m with low cooling conditions show a typical behaviour of PP of a spherulitic structure with a clear yielding and a strain hardening. The PP1272 film prepared by the lab unit under high cooling conditions and of thickness of 35 [micro]m shows a distinctive behaviour that has been extensively studied in our previous work (Sadeghi et al., 2007). The behaviour is due to a highly oriented lamellar structure (Figure 8a). A continuous strain hardening following the elastic region is indicative of such a structure. This is a result of tie chains stretching and reorientation of crystal blocks (Sadeghi et al., 2007). An almost perfectly elastic (solid-like) behaviour is observed for the MDO films prepared at DR of 3, 5, and 6. The strength is very high and the behaviour is attributed to the formation of strong, long and thick fibrils.

[FIGURE 8 OMITTED]

[FIGURE 9 OMITTED]

[FIGURE 10 OMITTED]

[FIGURE 11 OMITTED]

The response to the tensile test of films in the TD is another indication of differences in the crystalline structure of the film samples. As illustrated in Figure 10, PP1272 MDO films showed a clear yielding behaviour whereas PP4612 MDO films broke apart immediately after applying stresses indicating a very poor strength in the TD direction. The same behaviour was observed for the PP4612 obtained using the lab unit under high cooling conditions. These poor properties are probably due to the bimodal distribution in the molecular weight of PP4612 and fibril orientation as previously discussed. Among the PP1272 films the PP1272 MDO sample for DR = 1 (MDO = 1) shows a strain hardening. This is expected because of the very low orientation in both directions leading to similar stress-strain behaviour in both directions.

One of the main interests for MDO films is the reduction in gas permeability and improvement of barrier properties. That is most probably a result of orientation of amorphous phase (Taraiya et al., 1993). Figure 11 reports the oxygen permeability for PP1272 films obtained using the MDO process as a function of DR (the results have been normalized with respect to the film thickness). We observe a marked decrease of the permeability with DR, from about 150 x [10.sup.-6] to less than 60 x [10.sup.-6] Lm/[m.sup.2] day as the DR is increased from 1 to 6. This drop in oxygen permeability is more significant at lower DRs.

Light transmission (transparency) of packaging films always has been a concern for producers. For this matter, two parameters: haze and clarity are introduced. In film production, the quantitative assessment of clarity is just as important as that of haze. Clarity depends upon the linearity of the passage of light rays through the material. Small deflections of the light caused by scattering in the material cause a deterioration of the image. Haze is defined as the percentage of light that is scattered from a film sample for an incident beam of more than 2.5[degrees]. In other words haze measurements are made from wide-angle scattering, while clarity is determined by small-angle scattering. Figure 12 shows these two parameters as functions of the DR for PP1272 film samples prepared using the MDO unit. Haze decreases with DR and reached a plateau for DR>5 while clarity continuously increases with DR. One explanation is the following: with increasing DR the deformation of crystals increases and more are transformed into fibrils. These long fibrils are more oriented and, hence, the clarity increases. The other effective factor in reduction of haze and improving clarity is the decrease in thickness of samples with DR (the corresponding thicknesses are reported in Figure 10). We also performed tests with PP4612 films (MDO = 6) and obtained similar values for haze (4.2) and for clarity (96.8). This probably implies that the effect of molecular weight and bimodality on the light transmission properties of MDO films produced at high DRs is not very significant.

[FIGURE 12 OMITTED]

CONCLUSION

In this work we have investigated the effects of the processing parameters on the properties of two different PP resins. The processing units were uniaxially film stretching MDO industrial-scale machine and a lab-scale line under high cooling rate. Two PP grades of different molecular weight and molecular weight distribution were investigated.

The MDO generated a highly oriented fibrillar crystalline structure as a result of deformation of the initial crystalline structure. The fibrillar structure was getting stronger with increasing DR and, hence, the film strength increased showing a solid-elastic behaviour in tensile tests. The resin with a bimodal molecular weight distribution showed slightly greater orientation functions for films produced by the MDO line in comparison to the other linear PP of a narrower distribution although the molecular weight was lower. A very poor tensile strength in the TDwas measured for the lab unit samples. This test on MDO samples revealed a strain hardening behaviour before failure with an exemption for the bimodal resin PP4612 (DR = 6). The permeability to oxygen was greatly reduced with increasing DR for the MDO films. Finally, the haze property for the MDO samples decreased with DR reaching a plateau over DR of 5 whereas clarity rises continuously with DR.

NOMENCLATURE

D diffusivity ([m.sup.2]/s)

D ratio of the absorbance

[f.sub.a] orientation function of amorphous phase

[f.sub.av] average orientation function (Equation 3)

[f.sub.c] orientation function of crystals

[f.sub.i] orientation function (Equation 2)

G' elastic modulus (Pa)

G'' loss modulus (Pa)

H relaxation spectrum (Pa)

[M.sub.w] weight averaged molecular weight (g/mol)

P permeability coefficient=DS ([m.sup.2]/(s kPa))

S gas solubility ([m.sup.3]/([m.sup.3] kPa))

[X.sub.c] degree of crystallinity

Greek Symbols

[[eta].sup.*] complex viscosity (Pa s)

[[eta].sup.+.sub.E] transient elongational viscosity (Pa s)

[[eta].sub.0] zero-shear viscosity (Pa s)

[tau] relaxation time (s)

[omega] frequency (rad/s)

Abbreviations

DR draw ratio

EVOH ethyl vinyl alcohol polymer

HDPE high density polyethylene

LLDPE linear low density polyethylene

MD machine direction

MFI melt flow index

ND normal direction

PET polyethylene terephthalate

PVC polyvinylchloride

PP polypropylene

TD transverse direction

ACKNOWLEDGEMENTS

Financial support from NSERC (Natural Science and Engineering Research Council of Canada) and from FQRNT (Fonds Quebecois de Recherche en Nature et Technologies) is gratefully acknowledged. We also acknowledge the large infrastructure grant received from the Canadian Foundation for Innovation (Governments of Canada and Province of Quebec), which allowed us to build the unique POLYNOV facility. We are also thankful to Messrs. P. Cigana, L. Parent and P.M. Simard for their technical help.

Manuscript received April 6, 2008; revised manuscript received June 4, 2008; accepted for publication June 15, 2008

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Farhad Sadeghi ([dagger]) and Pierre J. Carreau * Department of Chemical Engineering, Center for Applied Research on Polymers and Composites, CREPEC, Ecole Polytechnique, Montreal, QC, Canada H3C 3A7

* Author to whom correspondence may be addressed. E-mail address: pcarreau@polymtl.ca

([dagger]) Present address: Chemical Engineering Department, University of Isfahan, Isfahan, Iran.

Table 1. Orientation factors for the films produced under different
process conditions

Sample Precursor film [f.sub.c] [f.sub.a]
 thickness
 ([micro]m)

PP1272-lab unit (with 35 0.60 0.32
rapid cooling), DR=6

PP1272-MDO line, cooled 35 0.10 0.05
under normal conditions,
DR=1

PP1272-MDO, DR=6 34 0.84 0.49

PP4612-MDO, DR=6 34 0.91 0.56