Comparison of structure development in injection molding of isotactic and syndiotactic polypropylenes.
Injection molding is one of the major processing operations for isotactic polypropylene. The morphology. crystallization and orientation development in injection molded isotactic polypropylene have been extensively studied in the literature (1-9). Kantz et al. (1) have subdivided the microstructure of injection molded isotactic polypropylene into three zones, i.e. skin layer, shear and core zones. Other researchers (2) have subdivided it into more layers. The skin layer possesses high orientation, which has been considered to be caused by high elongational deformation at the flow front (fountain flow) (10). Next to the skin layer, a highly oriented shear zone is developed by high shearing action at solidified wall during mold-filling stage. The core zone has generally low orientation and shows a spherulitic structure. Injection molded isotactic polypropylene exhibits polymorphic crystal structures (3-6), which include the Natta-Corradini monoclinic [alpha]-form crystal (11) and the hexagonal [beta]-form (12, 13).
There have been many studies of the crystal structures of syndiotactic polypropylene (14-18). However. there have been only a few studies (19-22) of structure development during processing because syndiotactic polypropylene has so far had only limited applications. Choi and White have investigated the crystallization and orientation behavior of syndiotactic polypropylene under uniaxial stress conditions in melt spinning (21) and under biaxial stress conditions in tubular film extrusion (22). In the melt spun fibers of syndiotactic polypropylene, a structural change from [(ttgg).sub.2] helix to trans-planar chain conformation has been found as spinline stress increases (21). In biaxially oriented tubular films of syndiotactic polypropylene, Form I unit cell is found with the (200) plane preferentially oriented parallel to the film surface (22).
There have been no published studies of injection molding syndiotactic polypropylene. In the present paper, we investigate the crystallization and orientation behavior of syndiotactic polypropylene in injection molding in comparison with those of isotactic polypropylene.
Two commercial grades of isotactic and syndiotactic polypropylenes, supplied by AtoFina Petrochemicals Inc., were used in this study. The isotactic polypropylene (iPP) is a general Ziegler-Natta catalyst grade with the typical melt flow rate of 10 g/10 min (ASTM D1238L), and its isotacticity is about 0.95 ~ 0.97. The syndiotactic polypropylene (sPP) was synthesized using a metallocene catalyst and had a melt flow rate similar to that of iPP. According to the material supplier, the sPP material possessed a racemic pentad content [rrrr] about 78%. This is lower in syndiotacticity than some of the materials described in early studies (14-18). However, injection molding experiments require a significant supply of material, and we must work with what is commercially available.
The shear viscosities for the materials were measured at a melt temperature of 200[degrees]C using a parallel-plate rheometer (Rheometrics RMS-800) at low shear rates (below 10 [s.sup.-1]) and an Instron capillary rheometer at high shear rates (above 10 [s.sup.-1]). Bagley end corrections and Welssenberg shear rate corrections were made on the capillary data.
The heating and cooling thermograms for the materials were obtained by differential scanning calorimetry (DSC) using a Perkin Elmer DSC-7. The polymer pellets were melted at 200[degrees]C, then cooled to 0[degrees]C at a cooling rate of 20[degrees]C/min, and then heated again to 200[degrees]C at a heating rate of 20[degrees]C/min.
The isothermal crystallization rates of iPP and sPP were measured at various crystallization temperatures using the same DSC instrument. The polymer melts at 200[degrees]C were quenched to a specified crystallization temperature, and then kept at the temperature until the heat flow became steady. Crystallization half-time was determined as the time interval required for the cumulative area of exothermic crystallization peak to reach 50% of the total peak area in a time-varying thermogram.
A reciprocating-screw injection molding machine (Van Dorn 55 HP) was used for the sample preparation. The materials were molded into dumbbell-shaped bars that had cross-sectional dimensions of width 12.7 mm and thickness 3.2 mm at the center.
The injection molded samples were prepared under various processing conditions as listed in Table 1.
In injection molded samples. we distinguish the machine direction (MD) as 1, thickness direction (ND) as 2, and the transverse direction (TD) as 3.
A small piece was cut from the middle part of each molded sample as shown in Fig. 1. For iPP samples, a thin-sliced sample (thickness ~50 [micro]m) was prepared by cutting this piece perpendicularly to TD using a Microtome (Reichert). For sPP samples, thick sheets (~1 mm thick) were used for better accuracy in measuring birefringence.
Birefrigence was measured using a Leitz cross-polarized optical microscope with a Berek (tilting) compensator and a white light source. When light passes through the MD-ND plane as shown in Fig. 1, [DELTA][n.sub.12] is measured. The birefringence was measured at many positions (z/H) along a continuous black fringe line on each sample. Here, z/H represents the normalized distance from the center (0 at center and 1 at surface).
WAXD Patterns and 2u Scanning
Thin-sliced samples of thickness ~50 [micro]m (for 20 scanning) and ~100 [micro]m (for WAXD patterns) were prepared by cutting the molded samples perpendicularly to the ND using the Microtome as shown in Fig. 1. Crystalline forms in the sliced samples were characterized through both WAXD patterns and 20 scanning experiments.
WAXD patterns were taken using an X-ray flat-film camera (G.E.) with a CuK[alpha] X-ray source (30 kV, 30 mA) and a Ni X-ray filter. The incident X-ray beam was perpendicular to the surface of sliced samples.
WAXD 2[theta] scanning experiments were carried out using another X-ray diffractometer (Rigaku), which has a CuK[alpha] X-ray source (40 kV, 150 mA), a graphite monochromator, and a slit detector. The incident X-ray beam was perpendicular to the MD. Intensity data were obtained for 30 seconds at every 0.02[degrees] increment of 2[theta]. We used a [theta]/2[theta] reflection scan mode in which only the crystallographic planes parallel to the sample surface can contribute to the diffraction intensity.
Several thin-sliced samples were stacked together into one multi-ply sample with dimensions of about 1 mm (TD) X 1 mm (ND) X 10 mm (MD]. The samples were mounted on a goniometer so that the MD coincided with the sample rotation axis (1-axis in Fig. 2).
Pole intensities were measured using the above G.E X-ray diffractometer with a graphite monochromator. For each crystallographic plane, pole intensities were measured for 10 seconds at every 50 increment of [phi] (0[degrees] ~ 90[degrees]) and at every 10[degrees] increment of [beta](0[degrees] 360[degrees]). Here, [phi] is defined as the azimuthal rotation angle and [beta] as the sample rotation angle.
The pole-figure experiments were carried out for the [(040).sub.[alpha]] and [(110).sub.[alpha]] planes of iPP samples and the [(200).sub.Form I] and [(002).sub.Form I] planes of sPP samples. For better understanding, every pole-figure was displayed in a TD-centered polar coordinate.
CRYSTALLINE ORIENTATION FACTORS
From the pole-figure intensity data, the crystalline orientation factors were determined. In injection molded samples. the principal orientation directions of samples do not necessarily coincide with the machine principal directions. In such a case, it is more desirable, as suggested by White arid Cakmak (23), to represent the level of orientation with respect to the orientation symmetry axes rather than the machine principal directions.
As represented in Fig. 2, the orientation of a crystallographic plane (hkl) with respect to an arbitrary direction Z ([[phi].sup.*], [[beta].sup.*]) can be calculated using the generalized equation
([cos.sup.2] [[phi].sub.hkl,z] = [[integral].sup.[pi].sub.0] [[integral].sup.2[pi].sub.0] I ([phi], [beta]) sin[phi] [cos.sup.2] [[phi].sub.hkl,z] d[beta]d[phi]/[[integral].sup.[pi].sub.0] [[integral].sup.2[pi].sub.0] I([phi], [beta]) sin[phi] d[beta]d[phi] (1)
[cos.sup.2] [[phi].sub.hkl,z] = [(P * Z).sup.2] =
= [(cos[phi] [cos[phi].sup.*] + sin[phi] cos[beta] [sin[phi].sup.*] [cos[beta].sup.*]
+ sin[phi] sin[beta] [sin[phi].sup.*] [sin[beta].sup.*]).sup.2] (2)
In the particular cases that the Z-direction coincides with the 1-, 2- or 3-directions, Eq 1 becomes
([cos.sup.2] [[phi].sub.hkl,1]) = [[integral].sup.[pi].sub.0][[integral].sup.2[pi].sub.0]I([phi], [beta]) sin[phi] [cos.sup.2][phi]d[beta]d[phi]/[[integral].sup.[pi].sub.0][[integral]. sup.2[pi].sub.0]I([phi], [beta]) sin[phi] d[beta]d[phi] (3a)
([cos.sup.2] [[phi].sub.hkl,2]) = [[integral].sup.[pi].sub.0][[integral].sup.2[pi].sub.0]I([phi], [beta]) [sin.sup.3][phi] [cos.sup.2][beta]d[beta]d[phi]/[[integral].sup.[pi].sub.0][[integral] .sup.2[pi].sub.0]I([phi], [beta]) sin[phi] d[beta]d[phi] (3b)
([cos.sup.2] [[phi].sub.hkl,2]) = [[integral].sup.[pi].sub.0][[integral].sup.2[pi].sub.0]I([phi], [beta]) [sin.sup.3][phi] [sin.sup.2][beta]d[beta]d[phi]/[[integral].sup.[pi].sub.0][[integral] .sup.2[pi].sub.0]I([phi], [beta]) sin[phi] d[beta]d[phi] (3c)
The orientations for the injection molded samples were represented in terms of the White-Spruiell biaxial orientation factors (24), defined as
[f.sup.B.sub.I J] = 2 < [cos.sup.2][[phi].sub.J,I] > + < [cos.sup.2] [[phi].sub.J,II] > - 1 (4a)
[f.sup.B.sub.II J] = 2 < [cos.sup.2][[phi].sub.J,II] > + < [cos.sup.2] [[phi].sub.J,I] > - 1 (4b)
where j is a crystallographic axis (j = a, b or c). In this study, we considered the reference direction I as the tilted symmetry axis in each pole figure and II as an orthogonal axis on the MD-ND plane.
For the iPP samples, the c-axis orientation factors in the monoclinic [alpha]-form unit cell were calculated using the (110) and (040) reflections according to Wilchinsky's method (25). We defined an a'-axis orthogonal to the b- and c-axes in the monoclinic unit cell to represent the a-axis orientations.
In our previous papers (21, 22), crystalline orientation factors of syndiotactic polypropylene were determined using both (200) and (002) reflections. For the injection molded sPP samples, however, the (002) pole intensities were too weak to be useful for orientation determination. So we used only (200) reflection for determining the crystalline orientation factors by assuming that [f.sub.b] [approximately equal to] [f.sub.a].
The shear viscosities of the two polymers are shown in Fig. 3. The shear viscosities of the two polymers were quite similar to each other at high shear rates. At low shear rates, sPP had much lower viscosity than iPP. This suggested that sPP had a lower molecular weight and a narrower molecular weight distribution (MWD) than iPP (26). The latter feature might be attributed to the characteristics of metallocene catalysts.
The DSC thermograms for the polymer materials are shown in Fig. 4. iPP showed a melting peak at ~163[degrees]C and the heat of fusion of ~93 J/g. sPP showed two melting peaks at ~115[degrees]C and ~130[degrees]C and a heat of fusion of ~35 J/g. The DSC crystallinities were ~0.52 for iPP and ~0.27 for sPP. Upon cooling, sPP crystallized at much lower temperature than iPP.
The isothermal crystallization half-times, presented in Fig. 5, showed a large difference between iPP and sPP. For iPP, crystallization rate increased steeply with decreasing crystallization temperature up to ~110[degrees]C, below which it crystallized very fast. For sPP, crystallization rate increased steeply with decreasing the crystallization temperature up to ~80[degrees]C, but even below the temperature, sPP did not crystallize as fast as iPP.
From the experimental results of both non-isothermal and isothermal crystallization, it is concluded that sPP crystallizes at much lower temperature and more slowly than iPP. The slow crystallization behavior of syndiotactic polypropylene has also been observed in melt spinning (21) and in tubular film extrusion (22).
Crystalline Forms in Injection Molded Parts
WAXD 2[theta] scans are shown in Fig. 6. The iPP samples exhibited both monoclinic [alpha]- and hexagonal [beta]-form peaks. The characteristic peak of the [beta]-form is the (300) reflection at 2[theta] angle 16.1[degrees] (CuK[alpha]). The relative peak heights (k) of [beta](300) reflection were determined according to the method of Turner-Jones et al. (27).
k = [I.sub.[beta](300)]/[I.sub.[beta](300)] + [I.sub.110] + [I.sub.040] + [I.sub.130] (5)
Figure 7a shows the k value profiles determined for the iPP samples prepared at the packing pressure of 7 MPa and various injection speeds. k had very low values in the skin layer and exhibited maximum peaks in the shear zone and in the core zone. With decreasing injection speed, k changed little in the core zone, but decreased considerably in the shear zone.
Figure 7b shows the effects of packing pressure on k. By reducing the packing pressure, k changed little in the shear zone, but decreased considerably in the core zone.
The sPP injection molded samples all exhibited typical diffraction peaks of the disordered Form I structure (18). The (211) reflection peak at 2[theta] 18.8[degrees] (CuK[alpha]), the characteristic peak of the ordered Form I, was not observed. Nor were the peaks of the Form III planar zigzag structure (16) observed.
Figure 8 shows birefringence profiles for the samples prepared at the packing pressure of 7 MPa and various injection speeds.
For the iPP samples, the three micro-structural layers (skin layer, shear and core zones) were clearly distinguished in the birefringence patterns. The skin layer had a thickness of about 0.1 mm and showed very high gradient of birefringence. The surface exhibited the highest birefringence in most samples. The skin layer birefringence was highly dependent upon the injection speed. The shear zone also exhibits relatively high orientations and shows some maximum peak(s). The birefringence in the shear zone increased a little with increasing injection speed. The shear zone thickness, which was about 0.3 ~ 0.5 mm, decreased with increasing injection speed. The fringe line in the core zone were not clearly seen but very diffuse owing to the scattering of light by large spherulites. The birefringence in this zone exhibited a small maximum peak and was little affected by the injection speed.
For the sPP samples, both skin layer and core zone were observed, but no distinct shear zone was found. The skin layer birefringence in the sPP samples showed high dependency upon the injection speed as in the iPP samples, but had much lower values than the iPP samples. The core zone birefringence in the sPP samples was quite high and showed maxima.
Injection molded samples were also prepared under low packing pressure (0.7 MPa) to investigate the effects of packing pressure on the orientation development. The results are shown in Fig. 9. For both iPP and sPP samples, the skin layer and the shear zone exhibited almost the same birefringence as found in the samples prepared under the packing pressure of 7 MPa. However, the core zone birefringence was greatly decreased at the lower packing pressure. The sPP sample showed near zero birefringence in the core zone.
The pole-figures for the injection molded samples are presented in Fig. 10. The pole-figures for both iPP and sPP samples showed orientation patterns close to uniaxial symmetry about certain directions. Some departures from the ideal uniaxial symmetry might be attributed mainly to the experimental errors caused by the asymmetric cross sections of pole-figure samples. The (002) pole-figures in the sPP samples may include relatively large errors resulting from the very low diffraction intensities.
In most samples, it is seen that the chain axis orientation directions are not parallel to the MD, but tilted by several degrees from the MD toward to the ND. The tilting angles were about 5[degrees] ~ 15[degrees] for the iPP samples and about 0[degrees] ~ 30[degrees] for the sPP samples.
Crystalline Orientation Factors
The biaxial crystalline orientation factors determined from the pole-intensity data are shown in Fig. 11. Both iPP and sPP samples showed orientations close to uniaxial orientation about the lilted chain axis direction (axis-I). Subsequently, the Hermans-Stein uniaxial orientation factors were determined using Eq 1 and Eq 2 with respect to the tilted chain axis direction of each sample, and the results are presented in Fig. 12. Both iPP and sPP samples exhibited [f.sup.H.sub.c, I] profiles very similar to the birefringence profiles.
The iPP samples exhibited high orientations in the skin layer and the shear zone, but relatively low orientations in the core zone. [f.sup.H.sub.a, I] has small positive values in most layers and negative values only in the skin layer, whereas [f.sup.H.sub.b, I] 1 large negative values in every layer.
The sPP samples exhibited higher crystalline orientation levels than the iPP samples. The skin layer had high orientations, and the core zone also had relatively high orientation levels.
As seen in Fig. 12, the WAXD patterns are in good correspondence with the Hermans-Stein crystalline orientation factors determined.
Isotactic polypropylene exhibits polymorphic behavior in the injection molded samples. The [alpha]- and [beta]-form crystals coexist, and their relative amounts vary with position in sample and processing conditions. This has also been observed by earlier investigators (3-6). In a general sense, both cooling rate and applied stress may control the competitive non-isothermal crystallization kinetics of the two crystalline forms. This is in agreement with Dragaun et al. (28) in that [beta]-form crystals are formed by high shear stresses. Such stresses are applied during the mold-filling stage and perhaps the packing stage.
The injection molded sPP samples had the disordered Form I structure. This is the same crystalline form that has been found in melt-crystallized samples of syndiotactic polypropylene copolymers by De Rosa et al. (29) and in other melt processing operations such as melt spinning (21) and tubular film extrusion (22) by the authors.
Effect of Crystallization Rate
As mentioned above, isotactic and syndiotactic polypropylenes possess much different crystallization rates. This results in significant differences in the orientation behavior in injection molding, as shown in Fig. 8 and Fig. 9.
Isotactic polypropylene is a fast crystallizing material. Oriented melt at the fountain-flow front (10) is quickly frozen on the cold mold wall and forms a highly oriented thin skin layer. Onto this skin layer, the flowing melt crystallizes under high shear stresses during the mold-filling stage and forms another highly oriented layer (shear layer). The core zone, which remains uncrystallized after the mold-filling stage, crystallizes slowly during the packing stage. The core zone has generally a low orientation level, owing to the low stresses during the mold-filling stage and subsequent stress relaxation. However, the core zone orientation can be increased by applying high packing pressure.
Syndiotactic polypropylene exhibits much different crystallization behavior and orientation development during injection molding than isotactic polypropylene. Owing to its slow crystallizing characteristics, syndiotactic polypropylene seems not to crystallize during the mold-filling and subsequently exhibits significant stress relaxation after the mold-filling stage. Thus. the orientation developed during the mold-filling stage is frozen-in, to some extent, only in the skin layer and near the skin layer, but becomes almost fully relaxed in other zones after the mold-filling stage.
Effect of Packing Pressure
Packing pressure gives rise to additional flow to compensate volume shrinkage during the packing stage. Because this additional flow passes mainly through the unsolidified core zone in mold, it has little effect on the structures in the skin layer or shear zone, but significantly changes the crystal structures and orientations in the core zone. Since the core zone experiences relatively low stress during the mold-filling stage and large stress relaxation after the mold-filling stage owing to the slow cooling rate, the orientation level in the core zone is generally very low. However, high packing pressure can considerably increase the orientation in the core zone because of the high shear stress developed during the packing stage.
Crystalline Orientation and Birefringence
We have sought to relate the crystalline orientation and the birefringence in injection molded parts. Both orientation factors and birefringence represent second moments of the orientation distribution (23, 24), and so these may reasonably be compared.
The birefringence ([DELTA][n.sub.12]) is plotted versus the Hermans orientation factor of the crystalline c-axis in Fig. 13. For both iPP and sPP samples, birefringence increased almost linearly with the crystalline orientation factors. The sPP samples, however, exhibit much lower birefringence than the iPP samples when compared at the same crystalline orientations.
For uniaxially orientated materials, the birefringence and the crystalline and amorphous orientation factors are inter-related by
[DELTA]n = [Xf.sub.cryst] [[DELTA][degrees].sub.cryst] + (1 - X)[f.sub.amor][[DELTA][degrees].sub.amor] + [DELTA][n.sub.form] (6)
where the form birefringence is generally neglected.
Using the DSC crystallinity data (X) and intrinsic birefringences of [[DELTA][degrees].sub.cryst] (0.029) and [[DELTA][degrees].sub.amor] (0.060) for isotactic polypropylene (30) and [[DELTA][degrees].sub.cryst] (0.023) and [[DELTA][degrees].sub.amor] (0.064) for syndiotactic polypropylene (31, 32), we may estimate [f.sub.amor]. The results are presented in Fig. 14. Compared to the crystalline orientation factors, the amorphous orientation factors exhibit much lower values. This is more marked in sPP samples. Similar results were found in melt-spun fibers (21).
Generally, [[DELTA][degrees].sub.amor] is considered much larger than [[DELTA][degrees]cryst] because [[DELTA][degrees]cryst] corresponds to helical chains in the crystalline state and [[DELTA][degrees].sub.amor] to zigzag chains in the amorphous state. Polymer chains in the amorphous phase thus have to be much more oriented than those in the crystalline phase to give the same orientation factors. Thus, low amorphous orientation in the samples is not surprising.
We have investigated crystallization and orientation development in injection molding isotactic polypropylene and syndiotactic polypropylene.
The molded iPP samples had both [alpha]- and [beta]-form crystal structures, and their relative ratios were very dependent upon local stress level and cooling rate. The birefringence profiles of iPP samples exhibited distinct three sublayers, i.e., the skin layer and shear zone with high orientations and the core zone with relatively low orientation.
The sPP samples had the disordered Form I crystal structure in our processing conditions. Compared to isotactic polypropylene, syndiotactic polypropylene crystallizes at lower temperature and much more slowly. Because of this characteristic, syndiotactic polypropylene does not form an observable shear zone during the mold-filling stage and apparently experiences severe relaxation of stress or orientation after the mold-filling stage. This results in very low orientations in the core zone when there is no significant contribution to the orientation by packing pressure.
After the mold-filling stage, further orientation was developed in the samples during the packing stage. This significantly affected the core zone orientation, especially for syndiotactic polypropylene.
Both iPP and sPP pole-figures exhibited almost uniaxial symmetry, but the chain axis orientation directions did not coincide with any machine principal direction. Despite the higher crystalline orientation factors, the sPP samples had relatively lower birefringence than the iPP. This indicates that the amorphous orientation factor in injection molded syndiotactic polypropylene is much lower than that in isotactic polypropylene.
[FIGURE 3 OMITTED]
[FIGURE 4 OMITTED]
[FIGURE 5 OMITTED]
[FIGURE 6 OMITTED]
[FIGURE 7a OMITTED]
[FIGURE 7b OMITTED]
[FIGURE 8 OMITTED]
[FIGURE 9 OMITTED]
[FIGURE 12 OMITTED]
[FIGURE 13 OMITTED]
[FIGURE 14 OMITTED]
Table 1 Processing Conditions of the Injection Molding Experiments. Barrel temperature ([degrees]C) 200 Mold temperature ([degrees]C) 40 Injection speed ([cm.sup.3]/s) 2, 6, 12, 28 Packing time (s) 22 Packing pressure (MPa) 0.7, 7
We are grateful to Dr. Joe Schardl of AtoFina Petrochemicals Inc. for his kind provision of the polymer materials used in this study.
(1.) M. R. Kantz, H. D. Newman, and F. H. Stigale, J. Appl. Polym. Sci., 16, 1249 (1972).
(2.) D. R. Fitchmun and Z. Mencik, J. Polym. Sci.: Polym. Phys., 11, 951 (1973).
(3.) M. Fujiyama, T. Wakino, and Y. Kawasaki, J. Appl. Polym. Sci., 35, 29 (1988).
(4.) P. Zipper, A. Janosi, E. Wrentschur, and P. M. Abuja, J. Appl. Cryst., 24, 702 (1991).
(5.) P. Zipper, A. Janosi, E. Wrentschur, W. Geymayer, E. Ingolic, W. Friesenbichler, and F. Eigl, Int. Polym. Process., 12, 192 (1997).
(6.) Y. Yu and J. L. White, Polym. Eng. Sci. (in press).
(7.) G. Menges, G. Wubken, and B. Horn, Colloid Polym. Sci., 254, 267 (1976).
(8.) E. Fleischmann, Int. Polym. Process., 4, 158 (1989).
(9.) A. I. Isayev, T. W. Chan. M. Gmerek. and K. Shimojo. J. Appl. Polym. Sci., 55. 821 (1995).
(10.) Z. Tadmor, J. Appl. Polym. Sci., 18, 1753 (1974).
(11.) G. Natta and P. Corradini, Nuovo Cimento Suppl., 15, 40 (1960).
(12.) H. D. Keith, F. J. Padden, N. M. Walter, and H. W. Wyckoff, J. Appl. Phys., 30, 1485 (1959).
(13.) S. Bruckner, S. V. Meille, V. Petraccone, and B. Pirozzi, Prog. Polym. Sci., 16, 361 (1991).
(14.) P. Corradini, G. Natta, P. Ganis, and P. A. Temussi, J. Polym. Sci., C16, 2477 (1967).
(15.) B. Lotz. A. J. Lovinger, and R. E. Cais, Macromolecules, 21, 2375 (1988).
(16.) Y. Chatani, H. Maruyama, K. Noguchi, T. Asanuma, and T. Shiomura, J. Polym. Sci., C28. 393 (1990).
(17.) Y. Chatani, H. Maruyama, T. Asanuma, and T. Shiomura, J. Polym. Sci., B29, 1649 (1991).
(18.) C. De Rosa, F. Auriemma, and V. Vinti. Macromolecules, 30, 4137 (1997).
(19.) M. Gownder, SPE ANTEC Tech. Papers, 44, 1511 (1998).
(20.) R. K. Sura, P. Desai, and A. S. Abhiraman, J. Appl. Polym. Sci. 81, 2305 (2001).
(21.) D. Choi and J. L. White, Int. Polym. Process., 15, 398 (2000).
(22.) D. Choi and J. L. White, Polym. Eng. Sci., 41, 1743 (2001).
(23.) J. L. White and M. Cakmak, Int., Polym. Process., 2, 48 (1987).
(24.) J. L. White and J. E. Spruiell, Polym. Eng. Sci., 21, 859 (1981).
(25.) Z. W. Wilchinsky, J. Appl. Phys., 30, 792 (1959).
(26.) W. Minoshima, J. L. White. and J. E. Spruiell, Polym. Eng. Sci., 20, 1166 (1980).
(27.) A. Turner-Jones, J. M. Aizlewood, and D. R. Beckett, Makromol. Chem., 75, 134 (1964).
(28.) H. Dragaun, H. Hubeny, and H. Muschik, J. Polym. Sci.: Polym. Phys., 15, 1779 (1977).
(29.) C. De Rosa, F. Auriemma, V. Vinti, A. Grassi, and M. Galimberti, Polymer, 39, 6219 (1998).
(30.) R. J. Samuels, J. Polym. Sci., A3, 1741 (1965).
(31.) M. F. Vulk, Optica Spektroskopia, 2, 494 (1957).
(32.) K. Nagaya, Y. Hatano, S. Hibi, T. Kawakami, Y. Kakizaki, and Y. Sakamoto, Kobunshi Ronbunshu, 56, 445 (1999).