Printer Friendly

Analysis of row nucleated lamellar morphology of polypropylene obtained from the cast film process: effect of melt rheology and process conditions.


A row nucleated lamellar structure as an initial morphology has been a matter of concern for developing microporous membranes based on a three-stage process (melt-extrusion/annealing/uniaxial-stretching) [1, 2]. In this process illustrated in Fig. 1 a precursor film with an appropriate row nucleated structure is first prepared, then by stretching at room temperature (cold stretching) voids are created due to lamellae separation and finally stretching in high temperature enlarge the voids resulting in a fairly regular distribution of the pores [1]. The preparation of a precursor film with the appropriate lamellar morphology is the most important issue. This structure is believed to be obtained by crystallization of the chains during melt stretching, which in literatures is called stress induced crystallization [1, 3]. Nogales et al. [3] studied the shear-induced crystallization of polypropylene. They concluded that a certain amount of high molecular weight chains is necessary to develop an oriented structure in the melt state. They determined this critical molecular weight for polypropylene using gas permeation chromatography (GPC) and small angle X ray scattering (SAXS) experiments. They showed that this "critical molecular weight for orientation" was dependent on shear rate and temperature up to a certain shear rate, then, it would be unchanged at higher shear rates. They speculated that the high molecular weight chains (with long relaxation times) were the main cause for the formation of shear-induced nuclei that acted as the initial nucleation sites for crystal lamellae growth. It was shown that the development rate of the oriented structure morphology was much faster for a high molecular weight resin. It has also been pointed out that a certain shear rate was necessary for the extension of the chains and this critical shear rate shifted to lower values as the molecular weight increased.

In polymer melt processing chain orientation in the melt state affects significantly the subsequent crystallization process. Chain orientation in the melt state improves the crystallization kinetics, as shown in the detailed study of Kornfied et al. [4]. They found that the time for crystallization under stress is orders of magnitude faster than under quiescent conditions. The temperature dependence of the crystallization time was much less pronounced in shear-induced crystallization while for quiescent crystallization it rose drastically with temperature. They considered the turbidity of the melt as a criterion of crystallization initiation, and found that in quiescent state, the crystallization time was 10,000 s (at a temperature of 141[degrees]C) while under stress this reduced to 16 s. Under quiescent crystallization conditions over a wide range of temperature, the lower molecular weight materials crystallized faster owing to the higher crystal growth rate (higher mobility), but under shear-induced conditions crystallization commenced with longer chains and the overall rate was higher. The increase in the crystallization rate was mainly due to the increase in the flow-induced primary nuclei. In the other words, the nucleation stage controlled the process, and the total crystallization rate was higher because of the faster nucleation rate.


In our previous paper Sadeghi et al. [5], we have shown that a high chain orientation in the melt state can result in the generation of a row nucleated lamellar structure for polypropylene. In the current work, we analyze key factors that may control the orientation of the lamellae in the row nucleated structure for polypropylene. These factors include rheological characteristics, molecular weight, and process conditions. We also make use of an FTIR method to detect the amorphous orientation between the lamellae. Finally, the influence of the orientation on the solid state tensile response of the samples is investigated. The effect of the lamellae orientation on the properties of membranes and their performance prepared by cold and hot drawing will be the topic of another paper.



The five different polypropylene grades used in this study are listed in Table 1.

PDC1280 is a general purpose grade and Pro-fax 6823 possess the lowest melt flow rate (MFR) in the homo polypropylene list introduced by Basell. Pro-fax 6823 is a high molecular weight homo polypropylene of Basell. Pro-fax 814 (Pf814) is a branched polypropylene with long chain branches. It was selected because of the effect of branched chains on the relaxation time of molten polyolefins. In this article, we refer to these materials as PP1280, PP6823, and Pf814. PP4292E1 and PP4612E2 are the grades used for the production of oriented films according to ExxonMobil. The manufacturer reports these two resins as homopolymers, an FTIR test was carried out and no trace of ethylene was found in the absorption spectrum of both resins.

Rheological Characterization

The linear viscoelastic behavior was obtained using a Rheometric Scientific SR-5000 controlled stress rheometer with 25-mm-diameter parallel plates. Stress sweep tests were performed to determine the linear viscoelastic region and the frequency sweep tests were carried out within the linear region at 230[degrees]C.

Measurements of the elongational viscosity were made with an elongational rheometer (Rheometrics RER 9000). The measurements were performed at 200[degrees]C and constant strain rate. The deformation of the molten specimen, immersed in a bath of hot silicone oil, was accomplished by exponentially varying the velocity of one end of the sample with time. The density of the oil was close to the density of the polymer to avoid buoyancy effects. The samples were carefully prepared so that all the residual stresses could be removed, to avoid deformation when the sample was melted. The granules were transfer-molded under vacuum using the Rheometric RSV-2100 sample preparation kit in conjunction with a Yokogawa programmable temperature controller model UP25. The uniformity of the resulting cylinder-shaped specimens was verified, and the average diameter was found to be 6 mm. The sample length used in this study varied from 22 to 24 mm. The end surfaces of the molded specimens were cleaned and then fixed to aluminum ties using a high temperature epoxy (24-h curing).

Film Preparation

The films were prepared by cast film processing with a slit die at the temperature of 220[degrees]C. To improve cooling, a fan was installed to supply air to the film surface right at the exit of the die. The key parameters were the fan speed, take-up roll speed, and air temperature. We worked at the maximum speed of the fan and constant die temperature, so the only variable was the take-up speed. The cast film samples were prepared at different draw ratios (roll speed to die exit velocity, thus implying different thicknesses) as shown in Table 2.

Orientation Features

Crystalline orientation measurements were carried out using a Bruker AXS X-ray goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper Cu K[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. Poles figures of (110) and (040) crystalline reflections were measured and orientation factors determined from the measurements.

Tensile Tests

Tensile tests were performed using an Instron 5500R machine, based on the D638-02a ASTM standard.


Infrared spectra were recorded on a Nicolet Magna 860 FTIR instrument from Thermo Electron Corp. (DTGS detector, resolution 4 [cm.sup.-1], accumulation of 128 scans). The beam was polarized by means of a Spectra-Tech zinc selenide wire grid polarizer from Thermo Electron Corp.


The elasticity of the resins in the melt state is taken as the ratio of the storage modulus to the product of the loss modulus with frequency [lambda] = G'/G"[omega] that is also called a characteristic elastic time [6]. It should be pointed out that, generally in the low frequency region, the elastic response of a melt is more pronounced largely because of higher degree of entanglements. The small amplitude oscillatory data (SAOS) in terms of the complex viscosity and the characteristic elastic time of the resins as functions of frequency are shown at 230[degrees]C, in Fig. 2a and b, respectively. Only the low frequency data for the elastic time are reported in Fig. 2b. The molecular weight of the linear homo polypropylenes can be evaluated through the zero-shear viscosity dependency of polymer melts to molecular weight [7]:

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

For PP1280, [[eta].sub.0] was determined from the frequency sweep test at 230[degrees]C and [M.sub.w] was obtained from GPC. The K factor was then calculated from Eq. 1 to be 1.21 x [10.sup.-15] (Pa.s/(g mol)[.sup.3.4]) and this is in agreement with results of the literature [7]. The molecular weight of the linear resins as listed in Table 3 based on the zero shear viscosity obtained at 230[degrees]C. For branched polypropylenes the zero shear viscosity would show a stronger dependence on molecular weight than 3.4 [8]. For this reason, no attempt was made to estimate the molecular weight of Pf814.The rheological behavior of these resins is typical of that of molten polyolefins. However, it is interesting to note that the branched polypropylene (Pf814), which exhibits the lowest viscosity, shows the longest elastic characteristic time at low frequency. The high elastic character of this resin is clearly related to its long chain branching structure and likely to its melt strength. The elasticity for the three resins, PP6823, PP4292E1, and PP4612E2, is quite similar at 230[degrees]C and this suggests that the effect of molecular weight at high temperature is less important for these linear polypropylenes. The larger values for PP4292E1 compared with PP1280 might come from a very small fraction of high molecular weight chains, although its weight average molecular weight is slightly lower than that of PP1280. This portion of high molecular weight chains could be detected through the relaxation spectrum of the resin [5].

Figure 3 shows the transient elongational behavior of four polypropylene resins at 200[degrees]C for two elongational rates, 0.1 and 0.3 [s.sup.-1]. The elongational behavior of PP4292E1, PP1280, and PP6823 is quite similar and no strain hardening is observed. No elongational tests have been carried out for PP4612E2, as its behavior was expected to be quite close to that of the first three linear resins. The larger elongational viscosity for PP6823 is attributed to its larger molecular weight (see Table 3). In contrast for Pf814 a highly strain hardening for both elongational rates is observed and starts at short times or small Hencky strains. The strain hardening for this material is very similar to that observed for low density polyethylene [9] and caused by the long branched chains. It should be noted that shear stress growth data were also obtained at very low shear rate and a good correspondence with the transient elongational viscosity data of Fig. 3 was observed, except for the strain-hardening portion of Pf814. However, the transient shear data are not presented to avoid crowding (Fig. 3).


The films obtained by cast film extrusion were examined by WAXD (wide angle X-ray diffraction) to analyze the orientation of the crystal reflections. To calculate the orientation functions we used the Herman equation:


f = 3 [cos.sup.2][phi] - 1/2 (2)

The details of the method have been mentioned in our previous paper [5] and the data collected from the WAXD apparatus give directly the Herman functions, f. To obtain a better perspective of the orientation feature, we report in Fig. 4 the [cos.sup.2][phi] values for all the samples where [phi] is the angle between the axes of the crystal blocks (a, b or c-axes) with respect to machine (MD), transverse (TD), or normal direction (ND). The determination of chain orientation in polymers is of great interest because it specifies the arrangement of crystal blocks in the bulk of the precursor film. This arrangement in turn is one of the important factors, which controls the physical and mechanical properties of the film. The high orientation of the crystal blocks is evidently a characteristic of a row nucleated lamellar structure and of a more significant order in the lamellae structure [5]. The higher molecular weight resin shows a larger orientation function of the c-axis along the machine direction. In all cases with increasing draw ratio, the c-axis moves toward the machine direction while the b-axis position almost stays in the TD-ND plane and the a-axis moves towards the TD-ND plane. A higher draw ratio will extend more chains in the machine direction, which means more nucleating sites for lamellar crystallization. A higher draw ratio also creates longer extended chains. It has been shown that a cluster of these long chains forms fibril crystals, which act as a base or support for the lamellar growth [4]. The effect of the draw ratio is more effective when the molecular weight is larger. The tendency of the a-axis to move towards the TD-ND plane increases for the higher molecular weight resins. The slightly higher orientation value for PP1280 compared with PP4292E1 is probably due to its larger molecular weight. For Pf814 and PP4612E2 the values don't show a large orientation and this is likely due to their lower molecular weight.


The effect of molecular weight on PP shear-induced crystallization has been already studied in a series of experiments with different grades of polypropylene [3, 10]. High, moderate, and low molecular weight linear PPs were used to investigate the nucleation process in shear induced crystallization. It was found out that the nucleation of the crystal entities for the lower molecular weight resin occurred after a longer shearing time [10]. However, the effect of long chain branching complicates the shear induced crystallization mechanism. In Pf814 chain interlocking probably provides more stretched chains in the molten state, Moreover, it is not certain that at very high stretching rates (much beyond the maximum Henky strain recorded by the extensional viscometer as reported in Fig. 3) applied in practical processing, the behavior will still remain the same or interlocking will be diminished. The second issue is lamellae defects and twisting that we believe are more severe for the lower molecular weight resins. The extended chains in the nucleation stage form extended fibrils that act as nucleating sites for the growth stage. It is speculated that Pf814 creates more stretched fibrils due to its long chain branches, but because of its lower molecular weight the fibrils are much shorter. This could result in more units that can easily twist or tilt. Figure 5 shows schematically the idea: the long fibrils in the case of the high molecular weight resin form a longer and stronger scaffold for the lamellae to grow on them.


Annealing is one of the main processing stages in developing microporous membranes by stretching as it first removes defects in the crystalline structure and second increases the lamellae thickness [11]. We annealed one of the samples (PP6823-2) for which the original orientation is illustrated in Fig. 4a to see the effect on the crystal orientation. The annealing was carried out under an extension of 2% with respect to the initial length at 145[degrees]C for 35 min. Figure 6 shows the orientation results for the annealed sample. As can be observed, annealing pushes the a-axis towards the TD-ND plane (a random distribution of a and b in the TD-ND plane), and this is an approach to the formation of a planar morphology. This kind of morphology is also desirable for improving the elastic behavior of the films. By conducting tests with other PPs it was observed that annealing is more effective for higher molecular weight resins.

We didn't study the effect of isotacticity (steroregularity of the chains) of PP on its crystallization. This important issue in film and fiber production processes has been discussed by other authors [12]. The rate of crystallization and crystalline chain-axis orientation for resins with higher isotacticity was significantly greater than for lower isotacticity resins. It was shown that, for the resin with high tacticity, an increase in the spinline stresses increases [f.sub.c] while [f.sub.a] and [f.sub.b] decrease and all reach finally a plateau. For resins of lower tacticity the trend was not simple and a maximum for [f.sub.c] was observed [12].

Fourier transform infrared spectroscopy was carried out to detect the amorphous phase orientation. The method is based on differences in the magnitude of infrared beam absorption along the machine and transverse directions. For an isotropic sample the absorbance is the same in both directions, but as the draw ratio increases the chains are likely more oriented in the machine direction and this causes higher density of the chains in that direction, which leads to higher absorbance. The Herman orientation function is used to describe the orientation of given molecular axis with respect to the sample direction [13]:


[f.sub.i.j] = ([[S.sub.j]/[S.sub.0]] - 1)1/[3 [cos.sup.2][[alpha].sub.i] - 1] (3)

where [S.sub.j] is the intensity of the peak in the spectrum obtained in MD, TD, or ND, and [S.sub.0] = (1/3)([S.sub.MD] + [S.sub.TD] + [S.sub.ND]), [alpha] is an angle related to the configuration of the bonds. For polypropylene the absorption values at the wave numbers of 998 and 972 [cm.sup.-1] are considered as the contribution of the crystalline part for the first and the combination of both crystalline and amorphous parts for the second. For both states the angle, [alpha], will be equal to 90[degrees] and in the case of the uniaxial orientation [S.sub.TD] = [S.sub.ND] the above formula can be modified as:


[f.sub.i.MD] = ([([S.sub.MD]/[S.sub.TD]) - 1]/([[S.sub.MD]/[S.sub.TD]) + 2]) = ([D - 1]/[D + 2]) (4)

where D is the ratio of the absorbance in the machine (parallel) to the transverse (vertical) direction. For the orientation function of the crystal phase, [f.sub.c], the band 998 [cm.sup.-1] is considered so D will be ([A.sub.II]/[A.sub.[perpendicular to]])[.sub.998], where A is the absorbance. For measuring the total orientation that includes both the crystalline and the amorphous phase orientations the band at 972 [cm.sup.-1] is selected and [f.sub.av] (average orientation) is calculated based on ([A.sub.II]/[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] (5)

where [X.sub.c] is the degree of crystallinity determined through DSC (differential scanning calorimetry). The results for the orientation functions vs. draw ratio are plotted in Fig. 7 for all the linear PPs. The trend for the orientation of the crystalline phase is in accordance with the WAXD results. However, the orientation function values are slightly lower than the WAXD values. As expected, the orientation increases with draw ratio, but the increases for the crystalline phase are much more significant, with a clear influence of the molecular weight. Almost the same order for PPs is observed for amorphous orientation but less influence of the draw ratio is seen and the values are closer to each other compared with the crystalline orientation functions.

Tensile tests have been carried out to analyze the response of the solid films to extension and the results are presented in Fig. 8. It is likely that for the larger draw ratio all the resins develop a lamellar structure while for smaller draw ratio the structure is not a fully oriented with a tendency to form twisted lamellae and also a high density of cross hatching [5]. All film samples prepared at the smaller draw ratio, except PP6823, show a clear yielding behavior (Fig. 8a), whereas the samples prepared at the larger draw ratio show no significant yielding (Fig. 8b). The higher yield strength for PP6823-2 in comparison with PP6823-1 can be attributed to the higher orientation function. A part of this also could be due to the larger number of active tie chains between the crystal blocks. A direct relationship between the yield strength and the fraction of the tie chains has been assumed by Nitta and Takayanagi [14]. For highly oriented samples that may possess more active tie chains the stresses are mostly concentrated in the space between the crystal blocks and are spent to stretch the tie chains. For low oriented samples the stresses are mostly spent to deform the crystal blocks (shear yielding) and orient them in the stretching direction. Hence, we believe that for highly oriented samples the role of tie chains is more critical as their number and orientation can influence the strength significantly [15].


Effect of the Cooling Process

Quenching the melt film at the die exit is an important issue for generating a lamellar structure [5]. Quenching prevents the chains to relax from their stretched state and stretched chains develop shishs upon crystallization. Although solidification takes place rapidly, enough time should also be provided for crystallization of the rest of the chains to form stable kebabs perpendicular to the shishs. We believe that a very fast quenching will result in more shish fibrils and much thinner kebabs so that the overall crystallinity does not change, as verified by DSC. Figure 9 shows the tensile results for two samples both produced with a draw ratio of 65. Sample No. 1 was obtained under moderate cooling (80% of the maximum speed of the fan used to blow air to the film at the die exit) and Sample No. 2 was prepared for a maximum speed of the cooling fan. We also examined these two samples by wide angle X-ray scattering to reveal the orientation. As it is shown in Table 4, the orientation functions of the samples are in the same range. We believe that the differences in the tensile response for Sample No. 2 are due to the formation of thinner lamellae but in larger population (as crystallinity remains almost unchanged).

Effect of Annealing

Annealing under low tension is the first stage for making a microporous membrane from the precursor film. As it was shown in the orientation results, annealing has a significant effect on the orientation and thickness of the crystal lamellae [11]. The improvement of orientation and thickness may be the result of the crystallization of the trapped end chains between the lamellae and also the melting of very thin lamellae and their recrystallization in the form of the thicker ones. Annealing at a temperature above the [alpha] transition will allow the crystalline blocks to move locally. This local sliding removes some defects in the crystalline structure of the lamellae and reorients them in the direction of the applied tension. This may explain the increase in orientation as shown in Fig. 6. The tensile behavior of one of the precursor films obtained from PP1280 (PP of moderate molecular weight) under a medium draw ratio of 48 is compared with the annealed sample (performed at 145[degrees]C and for 25 min.). As the test proceeds the sample passes the initial elastic region that is believed to be due to the stretching of the short tie chains. Following the elastic response region the second part coincides with voids creation as a result of chain scissions of the already stretched short chains. Pores are created between the lamellae and are enlarged as the stress increases linearly. With scission of short tie chains the stress is transferred to the longer tie chains causing them to be stretched out. Stretching of the longer tie chains will result in a local crystallization, which explains the strain hardening behavior observed in Fig. 10. Local crystallization creates bridges between the crystal blocks as it is shown in Fig. 11 (micrograph for the sample before the breaking point). Stretching will continue until the stress reaches a value large enough to breakdown the bridges and then rupture happens. The rupture for the annealed sample occurs at much shorter strain than for the nonannealed sample. It should also be noted that the initial orientation of the crystal blocks for the nonannealed precursor film is far from that of a planar morphology. Instead of a series stacked aligned lamellae there will be some lamellae tilting. For the nonannealed sample, in the second region of Fig. 10, a part of the stress is consumed for reorientation of the crystal blocks. This contribution is the largest for the less oriented samples and the smallest for the annealed and highest molecular weight PP sample (compare Figs. 4b and 6). This sample possesses a planar morphology associated with the highest orientation of the c-axis along the machine direction and the stress is largely concentrated on the crystal blocks and spent to separate them. That may also explain a sharper rise of the stress with strain. The larger strength for the annealed sample is likely due to the presence of thicker lamellae and also a direct concentration of the stress on the crystal blocks. The shoulder in the tensile curve for the annealed sample is probably due to the load transition to the long tie chains and also to voids formation.





Five different polypropylene resins have been characterized by rheology and the row-nucleated lamellar structure development during the cast film process has been investigated. The orientation of crystal blocks appears to be controlled by the molecular weight of the resins, as indirectly deduced from the melt index and the rheological properties (complex viscosity and elastic characteristic time) for the linear PPs. The behavior of the branched Pf814 was found to be totally different with a much longer elastic characteristic time at low frequency with respect to its melt index and strain-hardening effects under elongational flow. A row-nucleated morphology for polypropylene can be achieved via cast film extrusion under efficient cooling. The orientation of the c-axis of crystalline lamellae along the machine direction improved with draw ratio and this was more pronounced for the higher molecular weight resin. The strain-hardening and the larger elasticity for the branched Pf814 did not favor the orientation of the crystalline phase. The film samples showed different responses to tensile tests. The tensile properties were shown to increase with increasing draw ratio of the cast films. Except for the high molecular weight resin, a necking behavior was observed for samples prepared under a low draw ratio. For samples prepared under a high draw ratio nonsignificant yielding was observed and the tensile strength values were much larger than those for the low draw-ratio samples. Annealing the precursor film under a slight tension improved the orientation and significantly affected the tensile response. Two distinct regions were detected: the first was the typical elastic region attributed to the stretching of the short tie chains, while the second region was due to the contribution of the stretching of the longer tie chains, voids creation, and reorientation of crystal blocks.


1. T.H. Yu, PhD Thesis, Virginia Polytechnic Institute and State University (1996).

2. R.W. Callahan, R.W. Call, K.J. Harleson, and T.-H. Yu. U.S. Patent 6,602,593 (2003).

3. A. Nogales, B.S. Hsiao, R.H. Somani, S. Srinivas, A.H. Tsou, F.J. Balta-Calleja, and T.A. Ezquerra, Polymer 42, 5247 (2001).

4. J.A. Kornfield, G. Kumaraswamy, and A. Issaian, Ind. Eng. Chem. Res. 41, 6383 (2002).

5. F. Sadeghi, A. Ajji, and P.J. Carreau, J. Plastic Film Sheet 21, 199 (2005).

6. P.J. Carreau, Daniel C.R. De Kee, and R.P. Chhabra, Rheology of Polymeric Systems: Principles and Applications, Hanser, Munich (1997).

7. M. Fujiyama and H. Inata, J. Appl. Polym. Sci., 84, 2157 (2002).

8. F.A. Morrison, Understanding Rheology. Oxford University Press, New York (2001).

9. M.H. Wagner, S. Kheirandish, and M. Yamaguchi, Rheol. Acta, 44, 198 (2004).

10. F. Jay, J.M. Haudin, and B. Monasse, J. Mater. Sci. 34, 2089 (1999).

11. L.E. Alexander, X-ray Diffraction Methods in Polymer Science, Wiley, New York (1969).

12. D. Choi and J. White, Polym. Eng. Sci. 44, 210 (2004).

13. I.M. Ward, P.D. Coates, and M.M. Dumoulin, Eds., Solid Phase Processing of Polymers, Hanser, Munich (2000).

14. K.H. Nitta and M. Takayanagi, J. Polym. Sci. Part B: Polym. Phys 38, 1037 (2000).

15. M. Tayaanagi, K.H. Nitta, and O. Kojima, J. Macromol. Sci. B 4, 1049 (2003).

Farhad Sadeghi, (1) Abdellah Ajji, (2) Pierre J. Carreau (1)

(1) Center for Applied Research on Polymers and Composites, CREPEC, Ecole Polytechnique, Montreal, QC, Canada

(2) CREPEC, Industrial Materials Institute, CNRC, Boucherville, QC, Canada

Correspondence to: Farhad Sadeghi; e-mail:

Contract grant sponsors: NSERC (Natural Science and Engineering Research Council of Canada); FQRNT (Fonds Quebecois de Recherche en Nature et Technologies).
TABLE 1. Polypropylene grades used.

Resin code Company MFR 230[degrees]C/2.16 kg

PDC1280 Basell 1.2
Pro-fax 6823 Basell 0.5
Pro-fax 814 Basell 2.8
PP4292E1 ExxonMobil 2.0
PP4612E2 ExxonMobil 2.8

TABLE 2. Summary of the conditions for the cast films and their

No. Sample Draw ratio Thickness ([micro]m)

 1 PP6823-1 23 55
 2 PP6823-2 65 23
 3 PP1280-1 23 55
 4 PP1280-2 65 23
 5 PP4292E1-1 23 56
 6 PP4292E1-2 65 22
 7 PP4612E2-1 23 56
 8 PP4612E2-2 65 22
 9 Pf814-1 23 55
10 Pf814-2 65 23

TABLE 3. Zero-shear viscosity at 230[degrees]C and calculated molecular
weight of the PPs.

Resin Code [[eta].sub.0] (Pa.s) [M.sub.w] (kg/mol)

PP1280 16,000 420.5
PP 6823 31,000 510.8
PP4292E1 15,500 416.5
PP4612E2 8600 350.3

TABLE 4. Orientation for precursor PP1280 films.

f (Herman orientation a-axis b-axis c-axis
function) 1 2 1 2 1 2

MD 0.215 0.207 -0.475 -0.467 0.690 0.674
TD 0.136 0.083 0.239 0.219 -0.375 -0.302
ND 0.079 0.123 0.235 0.253 -0.314 -0.376

Sample No. 1 prepared under moderate cooling and Sample No. 2 prepared
under maximum cooling.
COPYRIGHT 2007 Society of Plastics Engineers, Inc.
No portion of this article can be reproduced without the express written permission from the copyright holder.
Copyright 2007 Gale, Cengage Learning. All rights reserved.

Article Details
Printer friendly Cite/link Email Feedback
Author:Sadeghi, Farhad; Ajji, Abdellah; Carreau, Pierre J.
Publication:Polymer Engineering and Science
Geographic Code:1CANA
Date:Jul 1, 2007
Previous Article:A novel method to prepare zinc hydroxystannate-coated inorganic fillers and its effect on the fire properties of PVC cable materials.
Next Article:Omastova special issue.

Terms of use | Privacy policy | Copyright © 2021 Farlex, Inc. | Feedback | For webmasters