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Crystallization and orientation development in fiber and film processing of polypropylenes of varying stereoregular form and tacticity.

INTRODUCTION

Stereoregular crystalline polypropylene was first synthesized in 1955 by Natta and his co-workers (1, 2). It was the isotactic form (iPP), which has asymmetric carbon atoms with all the same d or l configurations throughout the polymer. iPP has since become an important commercial polymer. Around 1960. Natta and his co-workers (3, 4) subsequently reported the synthesis of another stereoregular polypropylene, i.e. syndiotactic polypropylene (sPP), which has asymmetric carbon atoms with alternative d and l configurations. These early-stage sPPs prepared using vanadium-based Ziegler-Natta catalysts had low syndiotacticities (typically 0.6-0.8), which led to poor mechanical and thermal properties. sPP received little attention until highly stereoregular sPPs became commercially available in recent years through the application of the improved catalysts (metallocene).

Both iPP and sPP exhibit polymorphism, having several different crystalline forms. iPP exhibits monoclinic [alpha]-, hexagonal [beta]- and orthorhombic [gamma]-crystalline forms and a mesomorphic (smectic) form (5, 6). All these structures consist of the same 3/1 helices of (tgtg) chain conformations. sPP has four different crystal structures (7-11), i.e. Form I. Form II. Form III and Form IV. Form I (7) is the most stable form and consist of S(2/1)2 helices of [([t.sub.2][g.sub.2]).sub.2] chain conformations. Form II (8) has the same [([t.sub.2][g.sub.2]).sub.2] chain conformations and a slightly different orthorhombic unit cell. Form III (9, 10) has zigzag all trans (tttt) chain conformations. Form IV (11) has ([t.sub.6][g.sub.2])[t.sub.2][g.sub.2]) helical chain conformations.

Melt spinning is an important fabrication method for iPP. The development of structure in melt-spun iPP fibers has long been studied by many investigators (12-24). Activities on sPP are still in the early stages. A key investigation of crystallization during melt spinning of iPP and other polyolefins was made in a 1968 paper by Katayama. Amano and Nakamura (13) through online measurements of spinline temperature profiles, birefringence and X-ray scattering patterns. Similar online studies on melt spinning of iPP were made later by Ishizuka and Koyama (15) and Nadella et al. (16). iPP filaments spun through ambient air (13, 15-24) are usually found to possess the monoclinic [alpha]-crystalline form. Filaments quenched into ice water (12, 14, 16, 22-24) possess usually the mesomorphic structure. This structure has been also found in air-quenched iPP fibers (20, 21). Orientation development in melt-spun iPP and other polyolefin fibers has been investigated by various researchers beginning with Kitao et al. (17). The most striking feature in melt-spun iPP fibers is the existence of bimodal orientation (16, 18, 19), which originates from its peculiar crystallization behavior, i.e. epitaxial lamellar branching (25, 26). The levels of crystalline orientation factors and birefringence in melt-spun iPP fibers have been found to be correlated with spinline stress by Nadella et al. (16) and others (20, 22, 24).

Film is also an important application area for iPP. Again, activities on sPP are just beginning. Tubular blown film extrusion is an important fabrication method for polyethylene, but not common for iPP. However, tubular blown film extrusion can provide a useful tool for studying the development of structure in biaxially stressed polymer melts. Studies of structure development in tubular blown film extrusion of iPP have been reported by Shimomura et al. (27) and D. Choi and White (28). They found through WAXD pole-figure investigations that the (040) plane normal, or crystallographic b-axis, tends to orient in the film normal direction. Shimomura et al. (27) correlated the birefringences [DELTA][n.sub.ij] of iPP blown films with stress difference ([[sigma].sub.i]-[[sigma].sub.j]).

Most of the previous studies (12-24, 27, 28) of structure development in melt processing of iPP were limited to high tacticity iPPs prepared by using Ziegler-Natta catalysts. There have been few studies on structure development in melt processing of lower tacticity iPPs, which have been commercially available in recent years through the application of metallocene catalysts. In this paper, we study structure development in melt spinning of these lower tacticity iPPs and also compare the results with conventional high tacticity iPPs. Crystallization in tubular film extrusion of these low tacticity iPPs is also briefly discussed.

Structure development in melt processing of sPP has been studied by few researchers. Melt spinning of sPP has been studied by Gownder (29), Sura et al. (30, 31) and D. Choi and White (24). All these papers describe slow crystallization behavior during melt spinning of sPP. i.e. the spun filaments did not crystallize completely in the spinline. Sura et al. (30, 31) found through WAXD investigations that with increasing spinning speed, the fiber structure changed from ([t.sub.2][g.sub.2])[.sub.2] helical chain conformations (Form I) to zigzag all trans (tttt) chain conformations (Form III). They also noted that the fiber density decreased with increasing spinning speed. In this paper, we describe in more detail the structure and orientation development in melt spinning and tubular blown film extrusion of sPP.

EXPERIMENTAL

Materials

Six different commercial grades of polypropylenes, supplied by AtoFina Petrochemicals and Basell, were studied in this paper. These include two metallocene catalyst sPPs (sPP-1 and sPP-2) having syndiotacticities ~78 %[rrrr], two conventional Ziegler-Natta catalyst iPPs (H-1 and H-2) having isotacticities 97-98 %[mm] and two metallocene catalyst iPPs (L-1 and L-2) having isotacticities ~89 %[mm]. In each set of materials, we chose two different grades that had similar structural characteristics but different molecular weights. The characteristics of the materials are summarized in Table 1. The effects of regio-irregularity (i.e., head-to-head and tail-to-tail) were not considered in this paper.

Sample Preparation

Melt-spun filaments were prepared using a capillary rheometer with a circular die (diameter 1.6 mm. length-to-diameter ratio 19.3) and a take-up device. The filaments were normally spun through ambient air and sometimes into ice water. Several series of fiber samples were prepared by varying draw-down ratio (20-4000) and melt temperature (220[degrees]C and 260[degrees]C for iPPs, 180[degrees]C, 200[degrees]C and 220[degrees]C for sPP-1, and 220[degrees]C, 240[degrees]C and 260[degrees]C for sPP-2). Spinline tensions were measured during melt spinning using an electronic tensiometer.

Tubular blown films were prepared using tubular film extrusion machines with an annular die (diameter ~30 mm, gap ~1 mm). Using a high tacticity iPP (H-1) and sPP-1. several series of film samples were prepared by varying draw-down ratio (DDR) and blow-up ratio (BUR) at a constant extrusion temperature 230[degrees]C.

Sample Characterization

The crystal structure of the melt-spun fibers and tubular blown films was characterized by WAXD film patterns and 2[theta] equatorial scans.

Orientation factors of crystallographic axes were determined from the diffraction intensities of the (110)[.sub.[alpha]] and (040)[.sub.[alpha]] reflections for iPP and the (200)[.sub.Form I] and (002)[.sub.Form I] reflections for sPP. For fiber samples, the level of orientation was represented in terms of the Hermans-Stein uniaxial orientation factor (16, 22, 24, 32)

[f.sub.j]=[3 * [bar.[cos[.sup.2][[phi].sub.J]]] - 1]/2 (1)

where [[phi].sub.j] is the angle between the j-crystallographic axis (j = a, b or c) and the fiber axis. For film samples, the level of orientation was represented in terms of the White-Spruiell biaxial orientation factors (27, 33)

[f.sub.1.sub.j.sup.B] = 2 [bar.[cos[.sup.2][[phi].sub.j,1]]] + [bar.[cos[.sup.2][[phi].sub.j,2]]] - 1 (2a)

[f.sub.2.sub.j.sup.B] = 2 [bar.[cos[.sup.2][[phi].sub.j,2]]] + [bar.[cos[.sup.2][[phi].sub.j,1]]] - 1 (2b)

where the 1- and 2-axes are chosen as the machine direction (MD) and the transverse direction (TD) of the films.

The birefringence of the melt-spun fibers was measured using a cross-polarized optical microscope with a Berek tilting compensator. For film samples, two out-of-plane birefringences [DELTA][n.sub.13] and [DELTA][n.sub.23] were determined from the three principal refractive indices ([n.sub.1], [n.sub.2] and [n.sub.3]) measured using an Abbe refractometer.

[FIGURE 1 OMITTED]

The crystallinities of the fibers and films were measured by DSC at a heating rate of 20[degrees]C/min. We assume the heat of fusion ([DELTA][H.sub.m.sup.0]) for 100% crystallinity to be 177 J/g for iPP and 130 J/g for sPP (34).

RESULTS

Quiescent Crystallization Rates

The rates of crystallization of the materials were measured under both non-isothermal and isothermal conditions using a DSC. The non-isothermal crystallization experiments were carried out by cooling 200[degrees]C melts at various rates (10[degrees]C/min ~ 70[degrees]C/min). Figure 1a shows continuous cooling transformation (CCT) curves (14, 16, 23) constructed using the crystallization onset temperatures during cooling and the corresponding elapsed times. For each material set, the differences in crystallization rate between the two different grades were not so significant. The order of non-isothermal crystallization rates of the materials was found to be

H-1, H-2 > L-1. L-2 [much greater than] sPP(1), sPP(2)

The isothermal crystallization experiments were carried out by quenching 200[degrees]C melts to various crystallization temperatures and measuring the crystallization half-times ([t.sub.1/2]), i.e. the time required for crystallization to reach 50% of final crystallinity. Figure 1b shows the results. The order of isothermal crystallization rates, i.e. ([t.sup.1/2])[.sup.-1], of the materials was found to be the same as that found in the non-isothermal experiments, but the differences in crystallization rate between the two different grades of each materials set were more significant in this case. Both low and high tacticity iPPs had very high crystallization rates at low crystallization temperatures (below about 105[degrees]C). For sPP, the melts crystallized quite rapidly at crystallization temperatures below about 75[degrees]C, but the crystallization rates were not as high as in iPP.

PART I. MELT-SPUN FIBERS

Crystallinity

The DSC-determined crystallinities of the melt-spun iPP fibers are plotted in Figs. 2a and 2b as a function of draw-down ratio. For the high tacticity iPPs (H-1 and H-2), with increasing draw-down ratio, the fiber crystallinity generally slightly increased from ~0.47 to ~0.51 with some exceptions at high spinning temperature (e.g. 260[degrees]C). For the low tacticity iPPs (L-1 and L-2), with increasing draw-down ratio, the fiber crystallinity increased more significantly from ~0.45 to ~0.52 for L-1 and from ~0.42 to ~0.48 for L-2. However, at high draw-down ratios, the crystallinity of the low tacticity fibers dramatically decreased again.

[FIGURE 2 OMITTED]

In Fig. 2c, we plot the crystallinities of the melt-spun sPP fibers as a function of spinline stress. As the spinline stress increased, the fiber crystallinity first increased but slightly decreased again at high spinline stresses.

Crystal Structure

In WAXD investigations, most of the melt-spun iPP fibers prepared in this study exhibited the monoclinic crystal structure ([alpha]-form). However, the low tacticity L-2 fibers spun at 260[degrees]C exhibited the mesomorphic structure at low and modest draw-down ratios. The WAXD patterns for these fibers are shown in Fig. 3a.

Another notable observation made on the crystal structure of the melt-spun iPP fibers is the formation of hexagonal [beta]-crystalline form. Significant amounts of hexagonal [beta]-crystalline form were found, as a mixture with [alpha]-form, in the low tacticity fibers (L-1 and L-2) spun at high draw-down ratios. These fibers showed a small peak or a shoulder at 2[theta] angle about 16[degrees], corresponding to the (300)[.sub.[beta]] reflection, in the WAXD scans. This peak appeared more clearly in high resolution WAXD patterns. An example of the WAXD patterns for these fibers is illustrated in Fig. 3b. This (300)[.sub.[beta]] peak was observed in the other fibers we prepared but showed negligibly weak diffraction intensities. From the peak intensities of the (300)[.sub.[beta]] and (040)[.sub.[alpha]] reflections in the WAXD 2[theta] equatorial scans, we determined the relative [beta]-form contents. B, defined as

[FIGURE 3 OMITTED]

B = [I.sub.[beta](300)]/[[I.sub.[beta](300)] + [I.sub.[alpha](040)]] (3)

The results are shown in Fig. 4. The high tacticity iPPs (Fig. 4a) exhibited relatively small B values (about 0.08-0.15) in all spinning conditions we used. With increasing draw-down ratio, these high tacticity fibers showed a slight decrease in B value at low draw-down ratios (less than ~1000), then a small increase at intermediate draw-down ratios (near 1000), and then decreased again at high draw-down ratios. However, for the low tacticity fibers (L-1 and L-2), with increasing draw-down ratio, B first slightly decreased but sharply increased again (to about 0.25) at high draw-down ratios.

sPP also exhibited polymorphic crystallization in melt spinning. Typical WAXD patterns for the sPP fibers are illustrated in Fig. 3c for the sPP(1) fibers spun at 220[degrees]C. It exhibited Form I helical structure (35) at low draw-down ratios and Form III zigzag all trans structure (10) at high draw-down ratios. In all series of the fiber samples we prepared, this structural change was observed at the same spinline stress level (about 7-8 MPa).

[FIGURE 4 OMITTED]

Ice-water quenching did not produce any significant effects on the crystalline structure of the melt-spun sPP fibers.

Crystalline Orientation Factors

The crystalline orientation factors of the melt-spun iPP fibers are plotted in Figs. 5a and 5b as a function of spinline stress. The high tacticity fibers (H-1 and H-2) showed a good correlation of crystalline orientation factors with spinline stress. However, the low tacticity fibers (L-1 and L-2) showed much different trends in the crystalline orientation factors with spinline stress. At low spinline stresses, the L-2 fibers showed much lower orientation levels than the others. At higher spinline stresses, both L-1 and L-2 fibers had the same orientation levels, and the chain-axis (c-axis) orientation factors ([f.sub.c]) increased up to about 0.80, which is higher than in the high tacticity fibers at the same spinline stress. However, as the spinline stress further increased above a certain level (~5 MPa), [f.sub.c] sharply decreased again, and [f.sub.a*] increased. The b-axis orientation factors ([f.sub.b]) remained almost constant ([f.sub.b]~-0.45).

The crystalline orientation factors of the sPP melt-spun fibers are shown in Fig. 5c. The sPP fibers also showed a good correlation of crystalline orientation factors with spinline stress. The chain-axis (c-axis) orientation factors ([f.sub.c]) increased monotonically up to about 0.9 with spinline stress. This is much higher than in iPP. The melt-spun sPP fibers exhibited essentially the same a- and b-axes orientation factors, i.e. [f.sub.a]~[f.sub.b], which had relatively large negative values at modest spinline stress levels and decreased to about -0.45 at higher spinline stresses.

Birefringence

The birefringence sof the melt-spun iPP fibers are plotted in Fig. 6a as a function of spinline stress. For both high and low tacticity fibers, the birefringence changed in trends very similar to the chain-axis (c-axis) orientation factor ([f.sub.c]) shown in Fig. 5a and 5b. The high tacticity fibers (H-1 and H-2) showed a good correlation of birefringence with spinline stress. As the spinline stress increased, the birefringence rapidly increased to a certain level (~0.017), then showed a plateau, and then gently increased again to ~0.020. The low tacticity fibers (L-1 and L-2) showed much different trends in the birefringence with spinline stress. At low spinline stresses, the L-2 fibers showed lower birefringence levels than the others. At higher spinline stresses, both L-1 and L-2 fibers exhibited the same levels of birefringence, which increased up to ~0.022 (This is almost 30% higher than the value of the high tacticity fibers at the same spinline stress). As the spinline stress further increased, the birefringence sharply decreased again to about 0.015.

The melt-spun sPP fibers showed a good correlation of birefringence with spinline stress as shown in Fig. 6b. As the spinline stress increased, the birefringence increased up to about 0.11 and then slightly decreased. The levels of birefringence of the sPP fibers were much lower than in the iPP fibers at the same spinline stresses.

PART II. TUBULAR BLOWN FILMS

Crystal Structure

In tubular blown film extrusion, both high and low tacticity iPPs exhibited the monoclinic [alpha]-crystalline form under our processing conditions, and sPP Form I helical structure (35).

[FIGURE 5 OMITTED]

[FIGURE 6 OMITTED]

Pole Figures

The pole figures for the iPP (H-1) tubular blown films are shown in Fig. 7. For the films prepared at high DDR/BUR ratios, the (040) reflection showed orientation patterns close to uniaxial symmetry about the film MD. With decreasing DDR/BUR ratio, the (040) pole figures showed more increasingly the maximum intensity in the film normal direction (ND), indicating that the crystallographic b-axis is preferentially oriented in the ND. The film sample prepared at DDR/BUR ~1 (e.g. 3/2.8) showed almost equal-biaxial orientation with respect to the MD and TD.

The pole figures for the sPP tubular blown films are presented in Fig. 8. The (200) pole figures always showed the maximum intensity in the ND, indicating that the a-axis tends to orient in the film ND. The (002) reflection representing the crystalline chain-axis orientation showed the maximum intensity in the film MD at low BUR/DDR ratios and in the TD at high BUR/DDR ratios.

Crystalline Orientation Factors

Figure 9a shows the crystalline biaxial orientation factors for the iPP blown films. The b-axis orientation factors [f.sub.1b.sup.B] and [f.sub.2b.sup.B] had negative values, indicating that the b-axis is preferentially oriented in the film ND. The c-axis orientation factors [f.sub.1c.sup.B] and [f.sub.2c.sup.B] had positive values, indicating that the polymer chains are preferentially oriented parallel to the film surface. For all the iPP films we prepared. [f.sub.1c.sup.B] was larger than [f.sub.2c.sup.B], indicating higher chain orientation in the MD than in the TD. With increasing DDR at a constant BUR, the chain-axis (c-axis) orientation factors showed a more significant increase in [f.sub.1c.sup.B] than in [f.sub.2c.sup.B]. With increasing BUR at a constant DDR, it showed a significant increase in [f.sub.2c.sup.B] while [f.sub.1c.sup.B] changed little.

[FIGURE 7 OMITTED]

Figure 9b shows the crystalline biaxial orientation factors for the sPP blown films. It shows more significant variations with processing conditions. The a-axis orientation factors [f.sub.1a.sup.B] and [f.sub.2a.sup.B] had negative values, indicating that the a-axis is preferentially oriented in the film ND. With increasing draw-down ratio, both [f.sub.1a.sup.B] and [f.sub.2a.sup.B] decreased significantly. At high DDR/BUR ratios, [f.sub.1c.sup.B] was larger than [f.sub.2c.sup.B], i.e. the polymer chains tend to orient more preferentially in the MD. At low DDR/BUR ratios, [f.sub.1c.sup.B] was smaller than [f.sub.2c.sup.B], i.e. the polymer chains tend to orient more preferentially in the TD rather than in the MD.

Birefringence

In Fig. 10, we plot [DELTA][n.sub.13] versus [DELTA][n.sub.23] for the tubular blown films prepared at various DDRs and BURs. For both iPP and sPP films, the birefringences varied with processing conditions in trends quite similar to the crystalline chain-axis orientation factors, i.e. [f.sub.1c.sup.B] versus [f.sub.2c.sup.B]. They were in the ranges of 0.001-0.006 for [DELTA][n.sub.13] and 0.001-0.004 [DELTA][n.sub.23]. For both iPP and sPP films, with increasing DDR at a constant BUR, [DELTA][n.sub.13] significantly increased while there were little changes in [DELTA][n.sub.23]. With increasing BUR at a constant DDR, [DELTA][n.sub.23] significantly increased while [DELTA][n.sub.13] changed little.

DISCUSSION

Crystallization

In the melt spinning of iPP, some of the low tacticity iPP fibers prepared at low and modest draw-down ratios or spinline stresses had a mesomorphic structure. Formation of the mesomorphic structure in airquenched iPP fibers has been reported in the literature (20, 21). Lu and Spruiell (20) argued that the structure in melt-spun iPP filaments is governed by two competitive factors, i.e. spinline stress and cooling rate. We estimated the cooling rates in melt spinning at the position of solidification in spinline by solving the energy balance equation suggested by Kase and Matsuo (36). From the cooling rates estimated and spinline stresses, we constructed structural phase diagrams for the melt-spun iPP fibers as done previously by C. H. Choi and White (23). The results are shown in Fig. 11. It is seen that as the tacticity in iPP decreases, the boundary line between the monoclinic [alpha]-form and the mesomorphic form shifts toward lower cooling rates and higher spinline stresses region. This indicates that the mesomorphic structure is formed more readily in a lower tacticity iPP.

[FIGURE 8 OMITTED]

[FIGURE 9 OMITTED]

[FIGURE 10 OMITTED]

[FIGURE 11 OMITTED]

We found the hexagonal [beta]-crystalline form, coexisting with [alpha]-form, in melt-spun iPP fibers. Formation of the [beta]-crystalline form in melt-spun iPP fibers has not been reported in earlier studies (12-24). The amounts of [beta]-form crystals were negligibly small in most fiber samples except for a few low tacticity fibers spun at high spinning speeds. The relatively high [beta]-form contents and the low crystallinity levels found in these high-speed spun low tacticity fibers seem to be associated with its slower crystallization rates and the associated strong supercooling during melt spinning.

In the melt spinning of sPP, it also exhibited polymorphic crystallization behavior. Form I helical structure was formed at low draw-down ratios, and the zigzag all trans Form III structure at high draw-down ratios. This observation is in agreement with earlier studies by Sura et al. (30, 31). We found that the governing factor for this structural change is the level of spinline stress. The critical stress level for this structural change was found to be about 7-8 MPa. It is considered that when sPP molecules are in an extended form under high spinline tensions, the crystallization of the extended molecules into a helical structure is not favorable, but they prefer to crystallize into the zigzag structure (Form III). Notably, the crystallinity in sPP fibers decreased as the spinline stress increased. This is unlike most other linear polymers, where the crystallization is more favorable under high chain orientation. For sPP, the effect of cooling rate on crystallization in melt spinning seems negligible. This might be associated with its slow crystallization rate.

The crystal structures developed in both iPP and sPP tubular blown films were the same as those in the fibers spun at spinline stresses. This is because the processing conditions in tubular blown film extrusion are usually much milder than in melt spinning, i.e. the levels of applied stress and cooling rate are not so high.

Crystalline Orientation and Birefringence

For the melt-spun high tacticity iPP fibers (H-1 and H-2), the correlations of birefringence and crystalline orientation factors with spinline stress are generally in agreement with the earlier correlations found by other researchers (16, 22). However, the low tacticity fibers (L-1 and L-2) showed significant discrepancies with these correlations at both low and high spinline stress levels. The lower birefringence and crystalline chainaxis orientation factors of the low tacticity L-2 fibers spun at low spinline stresses are considered to be attributed to severe crystalline imperfection and the formation of mesomorphic structure.

In the low tacticity iPP fibers spun at high draw-down ratios or spinline stresses, the crystalline chainaxis orientation factor and birefringence were remarkably decreased. This can be explained by considering the bimodal orientation behavior of iPP (16, 18, 19). Two different orientation groups, i.e. the c- and a-axis oriented groups, coexist. By separating these two populations in the WAXD azimuthal scans of the (110) reflection, as done previously by Nadella et al. (16), we determined their relative populations. The results are shown in Fig. 12. It is seen that the relative population of the a-axis oriented group (branched lamellae) significantly varied with processing conditions, and the trends were almost opposite to the birefringence and the crystalline chain-axis orientation factors. For the high tacticity fibers, with increasing draw-down ratio, the percentage of the a-axis oriented group decreased from ~40% to ~20% at low draw-down ratios, then showed a small increase at intermediate draw-down ratios (around 1000), and then decreased again at high draw-down ratios. For the low tacticity fibers, with increasing draw-down ratio, the a-axis oriented group more significantly decreased (up to ~10%) at low draw-down ratios but remarkably increased again at high draw-down ratios. We consider that this increase of branched lamellae, in which the polymer chains are preferentially oriented perpendicular to the fiber axis, is directly responsible for the decreases of crystalline chain-axis orientation factor and birefringence in the low tacticity fibers spun at high spinline stresses.

[FIGURE 12 OMITTED]

Binsbergen and De Lange (25) found that the extent of epitaxial lamellar branching strongly depends on degree of supercooling or crystallization temperature, and low crystallization temperature or high degree of supercooling leads to intensive lamellar branching. When a low tacticity iPP is spun at high spinning speeds, one may expect high degrees of supercooling in the spinline due to its slower crystallization rate and rapid cooling. This should lead to severe lamellar branching.

Melt-spun sPP fibers showed much higher chainaxis (c-axis) orientation factors ([f.sub.c]) than iPP at the same spinline stresses. However, the levels of birefringence in sPP fibers were found to be much lower than in iPP fibers. By considering the levels of crystallinity (0.45-0.50 for iPP and 0.25-0.30 for sPP) of the melt-spun fibers, these experimental results may suggest that melt-spun sPP fibers possess lower orientation levels in amorphous phase than in iPP fibers.

In tubular blown film extrusion, iPP and sPP exhibited different characteristics in the development of orientation, iPP exhibited almost equal-biaxial orientation when DDR/BUR ~1, and [f.sub.1c.sup.B] > [f.sub.2c.sup.B] when DDR/BUR > 1. The sPP blown films we prepared exhibited much higher chain orientation in the TD than in the MD (i.e. [f.sub.1c.sup.B] < [f.sub.2c.sup.B]) when DDR/BUR ~1 and sometimes even when DDR/BUR [much greater than] 1. These differences seem to be associated with bubble shapes. For iPP bubbles, a distinct freeze line was formed very close to the die surface. Then, the melt extensions in both MD and TD occurred at almost the same positions on the bubble. This will lead almost the same orientation levels in the MD and TD when DDR/BUR ~1. However, for sPP bubbles, no distinct freeze line was found, but the bubble diameter increased gradually over a wide region from the die exit to much above it. Meanwhile, the melt extension in the MD occurred mostly near the hot die surface. This should lead to much higher chain orientation in the TD than in the MD at DDR/BUR ~1. These differences in the bubble shapes of iPP and sPP might be associated with the different crystallization rates.

CONCLUSIONS

In the melt spinning of iPP, we found significant differences in structure between low and high tacticity melt-spun fibers. The level of tacticity in iPP affects not only the crystal structure but also the level of chain orientation in fibers. The melt-spun fibers possess normally monoclinic [alpha]-crystalline form and sometimes a mesomorphic form under low spinline stress and high cooling rate conditions. The mesomorphic structure is more readily formed in lower tacticity fibers. In the low tacticity fibers spun at high draw-down ratios, significant amounts of hexagonal [beta]-form crystals were also found. These low tacticity fibers showed significant decreases in both crystallinity and crystalline chain-axis orientation. We consider that these unusual phenomena found in high-speed spun low tacticity fibers are attributed to strong supercooling during melt spinning, caused by its slower crystallization rates and rapid cooling, and the subsequent increase of epitaxially branched lamellae.

Melt-spun sPP fibers exhibited Form I helical structure at low spinning speeds and Form III zigzag all trans structure at high spinning speeds. We found that the level of spinline stress is the governing factor for this structural change. The critical stress level is found to be about 7-8 MPa. Melt-spun sPP fibers exhibit much higher chain-axis (c-axis) orientation factors ([f.sub.c]) and lower birefringence than in iPP fibers at the same spinline stresses. Considering the different levels of fiber crystallinity, this may suggest that melt-spun sPP fibers possess lower orientation levels in amorphous phase than in iPP fibers.

In tubular blown film extrusion, sPP exhibits the limit-disordered Form I crystal structure, and its a-axis tends to orient perpendicular to the film surface, while the b-axis does so in iPP blown films (all the monoclinic [alpha] structure). For sPP, high blow-up ratios were readily achievable owing to its slow crystallization characteristics, which led to high chain orientation in the film transverse direction.

ACKNOWLEDGMENT

We would like to thank Dr. Joe Schardl at AtoFina Petrochemicals Inc. for his kind provision of iPP and sPP materials.
Table 1. Characteristics of the Materials.

 iPP (high tacticity) iPP (low
 tacticity)
Sample Code H-1 H-2 L-1
Tacticity 0.97 (a)) 0.98 (a)) 0.89 (a))
[T.sub.m] ([degrees]C) (c)) 162 162 150
Crystallinity (c)) 0.56 0.56 0.50
Melt index (g/10 min) 4 9 3
[M.sub.w](X [10.sup.5] (d)) 2.73 1.67 2.19
[M.sub.w]/[M.sub.n] (d)) 6.8 7.8 5.8

 iPP (low
 tacticity) sPP
Sample Code L-2 sPP-1 sPP-2

Tacticity 0.89 (a)) 0.78 (b)) 0.78 (b))
[T.sub.m] ([degrees]C) (c)) 149 124 124
Crystallinity (c)) 0.49 0.26 0.26
Melt index (g/10 min) 8 4 10
[M.sub.w](X [10.sup.5]) (d)) 1.37 1.80 1.24
[M.sub.w]/[M.sub.n] (d)) 4.0 2.2 2.5

(a)) Meso-triad contents, [mm], estimated from IR absorbance ratio
([A.sub.841 cm-1]/[A.sub.973cm-1]) (37).
(b)) Racemic-pentad contents, [rrrr], provided by the material supplier.
(c)) Measured by DSC at 10[degrees]C/min.
(d)) Estimated from the shear viscosity data using Yamane and White
correlations (38).


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DONGMAN CHOI and JAMES L. WHITE

Institute of Polymer Engineering

The University of Akron

Akron, Ohio 44325
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