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Effect of a compatibilizer on the structural development of a thermotropic liquid crystalline polymer/polystyrene blend.

INTRODUCTION

Thermotropic liquid crystalline polymers (TLCPs) have been one of the most interesting developments in the chemistry and technology of polymeric materials during the past two decades (1-3). Upon melting, TLCPs give rise to highly organized liquid phases (mesophases) that tend spontaneously to pack parallel to one another to form highly oriented domains. Under elongational processing conditions, these oriented domains can develop a fibril morphology with a high degree of orientation, leading to enhanced mechanical properties (2). These properties enable TLCPs to be used as a reinforcing filler which is not present as a solid phase during processing of the composite, but instead forms when the material is cooled to a solid state. These blends have been called "in-situ" composites because of their in-situ shaping during processing (2). In-situ composites have attracted a great deal of interest because they can solve some problems that arise during the processing of conventional fiber-reinforced composites (3-5).

There have been innumerable studies of in-situ composites and TLCP blends with commercial thermoplastics (2, 4). Fibrils of the TLCP phase are usually formed near the die exit as a result of extensional forces in the flow direction (6). If a self-reinforcing fibril structure is to be obtained, the rheological properties of the TLCP blends should be carefully controlled (2). Not all TLCP blends with commercial thermoplastics can produce in-situ composites. It is generally accepted that in blends that include a minor phase of lower viscosity, the minor phase becomes elongated and the blend viscosity is reduced (7, 8). It is also accepted that in a capillary rheometer, the minor phase of lower viscosity becomes elongated into fibrils, provided that there is a sufficient quantity of the minor phase (for TLCP, greater than 10%) and that there are extensional flows involved, such as high rates of flow through the entrance to the capillary rheometer or drawing after die exit (9). Taylor (10) and others (11-13) studi ed two-phase systems of Newtonian fluids and found that the viscosity ratio, the interfacial tension, and the applied shear stress were crucial to the resultant morphology of the drop when there was no elongational force. A number of workers have applied these ideas, including the importance of both the viscosity ratio and elasticity ratio between the dispersed phase and the matrix phase, to polymer blend systems (9, 14, 15). However, the case in which the viscosity of TLCP is higher than that of the matrix has not been fully investigated yet.

All previous findings agree that some form of elongational deformation, as well as a matrix polymer with a higher viscosity than the TLCP phase, is necessary for deformation of the TLCP phase (12, 13, 16-19). However, most thermoplastics studied thus far are incompatible with TLCP. The reinforcing effect is less than that obtained from miscible systems (20-22). As in a blend of two flexible chain polymers, the complete lack of compatibility in TLCP/thermoplastic blends leads to a deterioration of the mechanical properties, even when micro fibrils are sometimes formed (23-25). This is more serious in the transverse direction. A compatibilized system can improve some of the properties of immiscible TLCP/thermoplastic blends (21, 25, 26), and one way to provide such compatibility to an immiscible binary blend is to add a compatibilizer (27). Very recently, compatibilization of thermoplastic blends with TLCP has been attempted (4, 22, 24, 26, 28). Seo and coworkers have intensively investigated the possibility of the deformation of a TLCP phase dispersed in a matrix whose viscosity is lower than that of the TLCP phase (29-33). Their conclusions can be summarized as follows: Deformation of the dispersed TLCP phase in a matrix whose viscosity is lower than that of TLCP phase is possible if the interfacial tension becomes low enough.

Though there have been innumerable studies on TLCP/thermoplastic blends, only a few studies on the blends of polystyrene (PS) and TLCP have been done previously (24, 34-38). Weiss and coworkers (34) investigated such a blend. The TLCP used was immiscible with the PS, but when an extensional component of flow was present during processing, the TLCP formed an elongated fibrous phase oriented in the flow direction. On the other hand, only spherical domains were observed without the extensional component. According to Zhuang et al.'s experiments (35), high extrusion rates and melt spinning result in fibrillar structures of TLCP in a PS matrix. Crevecoeur and Groeninckx (37) observed that the TLCP phase in injection-molded samples of PS/TLCP blends was only moderately elongated into fibrils and that the mechanical properties of the blends were below the values predicted by using the rule of mixtures. Ogata et al. (38) used three different grades of PS to study the PS/TLCP blends and found that the domain size of t he TLCP dispersed phase decreased with increasing PS viscosity (higher molecular weight). Kobayashi et al. (39) reported that 5% of a compatibilizer (thermotropic liquid crystalline blockgraft copolymer) added into PS/TLCP blends resulted in the best mechanical properties. Using a styreneglycidyl methacrylate copolymer that reacted with the TLCP functional groups to form a graft copolymer, Chiou et al. (24) observed that the compatibilized PS/TLCP blends had a reduced number of TLCP fibrils and had a tendency to form droplet domains.

In this study, our goal is to gain an understanding of the structure development of the TLCP phase in ternary blends of PS, TLCP, and a compatibilizer, poly-(styrene-co-maleic anhydride), based on the molecular characteristics of the macromolecules and the processing conditions, especially when the matrix has a lower viscosity than the TLCP does.

EXPERIMENTAL

The chosen TLCP was an aromatic liquid crystalline poly(ester amide), Vectra B950 (VB) (a copolymer based on 6-hydroxy-2 naphthoic acid (60%), terephthalic acid (20%), and aminophenol (20%)) produced by Hoechst Celanese Co. It has the high nematic transition temperature (283[degrees]C). Polystyrene used as a matrix polymer was purchased from Miwon Petro Chemical Co. Its [M.sub.w] was 29,000 g/mol and [M.sub.n] was 15,200 g/mol. Poly(styrene-co-maleic anhydride) (SMA) was obtained from Aldrich. Its [M.sub.w] was 224,000 g/mol and it contained ca. 7 wt% maleic anhydride.

The pellets of the PS and Vectra B were dried in a vacuum oven at 60[degrees]C and 120[degrees]C, respectively, for at least 24 hours before use. SMA was dried in a vacuum oven at 60[degrees]C for 72 hours. The TLCP content was kept at about 30 wt%. Dried pellets of VB and PS, and SMA powder were mixed in a container before blending in the extruder. Blending was carried out in a 42-mm Brabender twin-screw extruder (AEV651) at a fixed rotation speed of 20 rpm. It was equipped with a pulling unit imparting different draw ratio (DR), which we defined as the diameter at die exit to far down stream. Many strands of low DR (less than 8) were obtained. The extrusion temperatures of the feeding zone/transporting zone/melting zone/die were set as 140/300/300/290[degrees]C, respectively.

The Fourier transform infrared (FT-i.r.) spectrum was obtained using an Alpha Centauri Spectrometer (Mathson Instruments) with an average of 32 scans at 4 [cm.sup.-1] resolution.

Differential scanning calorimetry (DSC) studies of the thermal property characteristics were performed on a DuPont 910 DSC controlled by a 9900 thermal analyzer. Every thermogram was repeated at least twice to verify the reproducibility of the measurement. A DuPont 2000 thermal gravimetric analyzer was also used to observe the degradation of samples. The heating rate was 10[degrees]C/min and samples were heated to 800[degrees]C. Dynamic mechanical thermal analysis (DMTA) of the blends was carried out with a Polymer Laboratories Dynamic Mechanical Thermal Analyzer (Model 2) at a frequency of 1 Hz. A clamping geometry for the tensile mode was used.

Scanning electron microscopy (SEM) observations of the composite samples were performed on a Hitachi S-2200C model. Fractured surfaces of the blends were prepared by cryogenic fracture in liquid nitrogen followed by coating with gold in an SPI sputter coater. The morphology was determined using an accelerating voltage of 15 keV.

The rheological properties were measured using a UDS200 (Physica, Germany) rheometer on which 25-mm diameter cone and plate were mounted. The frequency range was set at 0.1-500 rad/sec, and the applied strain was 5%. Before the measurement, the samples were prepared using a compression molder at 260[degrees]C. The measurements were done under the nitrogen atmosphere.

Testing of the mechanical properties was done using an Instron Universal Testing Machine (Model 4204) at a constant temperature. A gauge length of 30 mm and a crosshead speed of 10 mm/min were used. All the reported results are averages of at least ten measurements.

RESULTS AND DISCUSSION

Transitions of the binary blends and ternary blends were resolved by using DSC and DMTA. In the DSC thermograms, a shift in the glass transition temperatures ([T.sub.g]) of PS in the ternary blends was obvious, but not for that of VB. This result was more obvious in the DMTA thermogram. The variation in the peak of tan[delta], which corresponds to the glass transition temperature, with increasing temperature is shown in Fig. 1. In the binary blend, two tan[delta] peaks are observed; the one at at 103[degrees]C is associated with the [T.sub.g] of PS and the other small one at 143[degrees]C is associated with the [T.sub.g] of VB. These two distinct values of the [T.sub.g] indicate that PS and VB are immiscible. However, in ternary blends, the higher and lower tan[delta] peaks, corresponding to [T.sub.g] of PS and VB, shifted toward each other reflecting improved compatibility although that of VB appears less vividly as a small shoulder. This is consistent with previous studies that reported changes in the glass transition temperatures of TLCP and polyamides (nylon 46 (29), nylon 66 (32), and nylon 6 (31)) with the addition of a compatibilizer.

The flow curves of the pure components and blends are presented in Fig. 2. The viscosity of VB is much higher than that of PS. At the processing shear rate (the apparent shear rate was about 40 [sec.sub.-1]), VB had a viscosity almost three orders of magnitude higher than that of PS. The viscosity values of the binary and the ternary blends in the range of investigated frequencies were higher than those of PS because of the blending with VB and its much higher viscosity. When a small amount of SMA (less than 2 wt%) was added, the viscosity of the ternary blend was close to that of the binary blend whereas the addition of more SMA increased the viscosity. In this case, the effect of TLCP acting as a processing aid to lower the blend viscosity was not observable. This is different from the results for nylon 6 and nylon 46 blends (29, 31) where blend viscosities lower than those of pure polymers were observed. For two different reasons, the reduction in the blend viscosities of those nylon blends with the additi on of TLCP has been attributed to the interlayer slip of the phases: a) incompatibility between the two phases and b) the formation of elongated fibrils of the TLCP phase, which tend to lubricate the melt (1). The viscosity increase in the ternary blend can be partly attributed to the molecular weight increase due to the formation of graft-block copolymers. The moduli of all of the blends were higher than those of PS because of the high modulus value of VB.

SEM micrographs can manifest the microstructure of the fractured surfaces of the composite extrudates prepared under various extrusion conditions. Figure 3 shows the fractured surfaces of a VB/PS binary blend and VB/PS/SMA ternary blends for a draw ratio of 1, which means that almost no drawing was applied. TLCP exists as spherical domains in the binary blend while TLCP has deformed shapes in the ternary blends with 1 and 2 wt% SMA. The ternary blends having 5 wt% SMA show a more irregular morphology. The morphologies of the binary and the ternary blends show the effect of the SMA. In the binary blend, most of the dispersed TLCP phases have spherical shapes, indicating almost no deformation, about 10-50 [micro]m in diameter, and are very nonuniformly dispersed in the PS matrix. The open craters and holes, as well as the gaps, between the dispersed TLCP phase and the PS matrix, indicate relatively poor adhesion at the interface. Interfacial adhesion is definitely improved in the ternary blends due to the inter action provided by SMA, but the fibrillarity (i.e., the number and the fineness of fibrils) deteriorates when an excessive amount of SMA (more than 3 wt%) is added. It looks as if coagulation and coalescence of the dispersed phase occur when an excessive amount of SMA is added.

In Fig. 3. a low draw ratio of 1 was applied to the samples so that the morphology of the final extrudates would predominantly reflect what had developed within the die. The effect of the draw ratio on the blend microstructure is demonstrated in Fig. 4. The textures of the samples reveal the presence of uniformly deformed TLCP fibrils. Figures 3 and 4 clearly show that textured droplets of TLCP are more abundant in more highly drawn samples. The micrographs in Figs. 3 and 4 exhibit the influence of elongational flow on the development of TLCP fibrils with SMA addition. The development of the fibrils reflects improved interfacial adhesion in the melt state. However, the fractured surfaces show that TLCP exists as fine and coarse fibrils, as well as ellipsoids and droplets, in the PS matrix when the draw ratio is low (Fig 3). This is partly attributable to the non-uniform distribution of the SMA. Although the possibility of thermal degradation of the PS matrix existed, that was disregarded when thermo-gravimetr ic analysis proved that the degradation of PS at the processing temperature (290[degrees]C) was very slight.

The fibrillation and droplet deformation due to the addition of SMA can be ascribed to two possible mechanisms: the occurrence of a chemical reaction or an interaction such as hydrogen bonding. O'Donnel and Baird (28) investigated the possibility of these two mechanisms for ternary blends of polypropylene/maleic anhydride grafted polypropylene/TLCPs (Vectra A 950, Vectra B950, LC 3000). We prepared a binary blend of VB and SMA (50:50) under the same processing conditions as were used for the ternary blends. Then, using a soxhlet extractor with boiling toluene, we extracted the soluble portion (SMA phase). After seven days of extraction, the extracted solution was poured into methanol. The precipitate was dried in a vacuum oven at 60[degrees]C for two weeks to remove any solvent residue. Then, FT-i.r. analysis was performed on the remnants. Though the results are not shown here due to space limit, the difference spectrum between the extract and the SMA clearly showed the characteristic peaks of naphthalene and amine groups, which cannot be seen in the SMA spectrum, at 1400 [cm.sup.-1], 1620 [cm.sup.-1] and 3400 [cm.sup.-1]. We emphasize that these peaks were not seen in the spectrum of the binary blend extract. This implies that the VB moiety is included in the ternary blend extract. Since VB is not soluble in toluene, we speculate that some reactions have occurred to produce a kind of graft copolymer. Recently, Seo (33) proposed reaction schemes occurring between the maleic anhydride group and the end groups of nylon 6 or VB. It is well known that maleic anhydride can react with the amine end groups of nylons (40). Thus, we speculate that a chemical reaction might have occurred to graft the VB moiety to SMA. It may be conjectured that the reacted VB moiety is from the low molecular weight VB formed by thermal depolymerization. However, a TGA thermogram of the VB showed no signs of thermal degradation until a temperature of 350[degrees]C was reached. The amine group of VB may react with the anhydride group of SMA to form a kind of (comb shape) block copolymer which has different branches and acts as a compatibilizer at the interface (41). Such copolymer molecules obviously exists In the interfacial region, and contribute significantly to a reduction in the interfacial tension and to the adhesion between the two phases (Figs. 3 and 4). Graft copolymers increase the mechanical strength at the interface by increasing the number of molecular entanglements between the two phases. However, there are two arguments against t he existence of the compatibilizer at the interface: diffusion and micelle formation. These factors prevent the migration of the graft copolymers to the interface. When there are not enough graft copolymers, some of the dispersed phase may have the same shape as that in the binary blend, then, the dispersed phase can not be deformed due to poor adhesion. On the other hand, when there are more than enough graft copolymers, coagulation and coalescence of the dispersed phase can take place (Fig. 3).

Blending of 30 wt% TLCP increases the blend viscosity values higher than that of PS (Fig. 2). The overall shape of the blend curve follows that of PS. However, ternary blends show complicated viscosity behavior because of the mixed morphology. In the viscous flow field, under the assumption that elongational fields are absent and the primary normal stress difference is negligible, which is not clearly proven yet (42), the deformation of the dispersed droplets is a balancing process between the shear force imposed by the flow field and the interfacial tension (interfacial stress divided by the curvature of the droplet). In the shear flow field, therefore, the final shape of the dispersed phase is a dynamic equilibrium between the shear stress and the compatibilizing action of SMA. Thus, efficient stress transfer by improved adhesion at the interface is the only practical way to achieve the deformation of TLCP droplets in a matrix whose melt viscosity at the processing conditions is lower than that of TLCP. Whe n a proper amount of the compatibilizer is added, some fine fibrils are formed in the ternary blend system even without a strong elongational drawing. This result is consistent with our previous works (29, 30, 32, 41) The presence of fibrils in the microstructure correlates with an increased tensile modulus. This is also In accordance with the fractured surface morphology. These results clearly show that fibril formation Is favored when a sufficient level of drawing is applied (Fig. 4).

The tensile modulus and strength for the composites are shown in Figs. 5 and 6. Though the sensitivity of the tensile properties to SMA addition is not obvious when the drawing is low (around 1), the tensile strengths for the ternary blends are higher than that of the binary blend, which is consistent with the fractured surface morphology. We speculate that the difference is due to the good adhesion and deformation of the TLCP phase caused by the addition of SMA. The influence of the draw ratio on the tensile modulus and the tensile strength is definitely attributed to the deformation of the TLCP phase. The overall tensile properties of the ternary blends are similar to those found in previous studies (26, 31), showing a definite Increase with the draw ratio. The tensile modulus and strength show interesting behaviors that depend on the amount of SMA added, i.e., they show maximum values when ca. 1 wt% of SMA is added and then decrease when more SMA is added. Though the tensile moduli of the 1 wt% SMA- and th e 2 wt% SMA-added ternary blends are almost the same, the tensile strength of the 1 wt% SMA-added system is higher. This implies that more elongation occurs in this system because of better adhesion at the TLCP/PS interface (see Fig. 7). The tensile modulus rapidly increases with draw ratio and then reaches a plateau value. Ternary blends containing 1 and 2 wt% SMA show steeper increases. The tensile strength also increases with the draw ratio. The tensile strength of the ternary blend having 1 wt% SMA shows the highest value, followed by the 2 wt%, 3 wt%, and 5 wt% SMA-added blends in that order. When more than 2 wt% SMA is added, the tensile strength decreases. However, all ternary blends show higher tensile strengths than the binary blend. The elongation at break is slightly affected by the addition of SMA, as shown in Fig. 7. Addition of TLCP to PS in the binary blend promotes the propagation of fractures before the matrix reaches its break point. Hence, the small elongation of the binary blend compared w ith that of PS indicates that the specimen is acting as a two-phase structure that will separate and fail prematurely, instead of acting as an integral specimen (29). The addition of SMA improves the elongation due to better adhesion, but excessive amounts of SMA will decrease the elongation.

Fibrous composites can fail during monotonic loading because of a number of competing micromechanisms of fracture, such as fiber breakage, matrix cracking and fiber pull-out (22). Depending on the form by which stored elastic strain energy in the fiber is released as well as on the strengths and toughnesses of the fiber, fiber-matrix interface, and the matrix, brittle fractures result from the earliest failure events (matrix cracks, for instance) and follow the distribution of flaws. A short fiber (or undeformed droplets) is pulled out if the force on the fiber is sufficient to cause some debonding. In contrast, for long fibers embedded in a matrix under a tensile stress, the fraction of fibers pulled out rather than broken approaches zero as the fiber length increases. This means that such fibers must fracture along a common crack plane or fracture into smaller segments before pulling out of the matrix. Examinations of the fracture surfaces of the composites (Figs. 3 and 4) reveal broken fibers of various le ngths protruding above the fracture surfaces of the matrix. These breaks occur because of a variation in the strength of the fiber at weak points beneath the surface of the cracked matrix. When a fiber fractures, in order for the new surfaces to move apart, the matrix must crack or plastically deform and the next fiber must fracture in the crack plane or beyond the crack plane. Then, it is axially pulled out of the matrix. Matrix plastic deformation and fiber pull-out require additional energy input to the material and provide a means of the energy dissipation necessarily associated with fracture. Since the frictional shear force opposes any force applied to extract the fiber, work must be done in overcoming this frictional force. A compatibilized composite has a larger frictional shear force due to the strong adhesion at the interface between the matrix and the fiber, thus requiring more energy to pull out the strands (26). As a result, the tensile strength of the system increases.

If the strand still maintains contact with the sheath of the matrix surrounding it, work must be done in pulling the fiber fragments against the restraining frictional force at the fiber-matrix interface (26). If the fiber does not maintain contact with the sheath of the matrix, the fibers can be easily pulled-out, so the elongation can not be increased. However, additional energy must be expended to break the strong adhesion at the interface in a compatibilized system. Fibers will not be simply debonded; they will sustain their fibril shapes over the gap between crack surfaces until additional energy is supplied (30). This allows the blends containing a small amount of SMA to attain a greater elongation at break than the binary system (Fig. 7). Excess SMA, however, brings about the coalescence of dispersed phase (Fig. 3). Owing to poor dispersion, the tensile strength and the tensile modulus are decreased as does the elongation because both total contacting area and the energy restraining the crack are decre ased. The elongation at break of all blends is decreased remarkably by blending with VB. The hard and strong, but inextensible strands of VB restrain the deformation of the polymer composite, thus leading to low elongation. The low elongation for large amounts of SMA added system is partly ascribable to aggregations of SMA. which act as the defects in the blend. They fracture first under a tensile force. Because of this low elongation, the tensile strength of systems containing large amounts of SMA is decreased. By the same token, the tensile strength for these ternary blends is decreased at high draw ratios.

CONCLUSIONS

We have investigated the possible deformation and fibril shape formation of TLCP (VB) in a PS matrix, whose viscosity is lower than that of VB, when small amounts of SMA, whose maleic anhydride moiety can interact with amine groups of VB to produce graft copolymers that can act as a compatibilizer at the interface, are added. We demonstrate that, even in a TLCP/PS blend characterized by an expectedly unfavorable viscosity ratio, fibrillation of a TLCP minor phase can be obtained under shear flow conditions, if an optimum amount of a compatiblizer capable of appropriately increasing the interfacial adhesion is employed. Good adhesion at the interface is believed to promote efficient stress transfer to deform TLCP droplets in the melt state at the processing temperature used. This result corroborates that the interfacial adhesion and the size of the dispersed droplets are as much decisive factors as the viscosity ratio of the TLCP to that of the isotropic matrix polymer in determining the deformation and the st ructure development of the TLCP phase. Scanning electron micrographs prove that if the proper amount of compatibilizer is added, the elongation and the orientation of the TLCP phase to a strand structure can occur even when the viscosity of the matrix polymer is much lower than that of the dispersed TLOP phase.

The mechanical properties (elongation, as well as the tensile strength and the tensile modulus) are improved when the proper amount of SMA is added, which enables improved adhesion at the interface. The loss of ductility of the matrix due to the addition of TLCP is ascribed to the incompatibility between PS and VB, which can be lessened by adding SMA. When excess SMA is added, however, nonuniform structures appear which cause the mechanical properties to deteriorate. As already noted in previous studies, an optimum amount of the compatibilizing agent is the key to obtaining the best morphology; hence, it is also the key to obtaining the best mechanical properties.

[FIGURE 1 OMITTED]

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[FIGURE 7 OMITTED]

ACKNOWLEDGMENT

This work was supported by the Korea Ministry of Commerce, Industry and Energy (Grant No. 2M11320).

REFERENCES

(1.) P. J. Collings and M. Hird, Introduction to Liquid Crystals, Taylor and Francis. London (1997).

(2.) F. P. La Mantia ed., Thermotropic Liquid Crystal Polymer Blends Technomic Publishing. Lancaster, Pa. (1993).

(3.) D. J. Williams Adv. Polym. Tech, 10, 173 (1990).

(4.) V. Handlos and D. G. Baird, J. Macromol. Sci Rev., C35 (2), 183 (1995).

(5.) C. Carfagna, E. Amendola, and M. R. Nobile, in International Encyclopedia of Composites. Vol. II, S. M. Lee, ed., VCH Publishers, New York (1990).

(6.) G. Crevecoeur and G. Groeninckx, Polym. Compos., 13, 244 (1992).

(7.) F. P. La Mantia and A. Valenza MakroinoL Chem. MacromoL Symp., 56. 151 (1992).

(8.) D. G. Baird and R. Ramanathan, in Multiphase Macromolecular Systems, B. M. Culbertson, ed., Plenum Press, New York (1989).

(9.) F. P. La Mantia, A. Valentia, M. Paci, and P. L. Magagnini, Polym. # Sci., 30, 7(1990).

(10.) R G. Taylor, Proc. R. Soc. A, 146, 501 (1934).

(11.) F. D. Rumscheidt and S. G. Mason, J. Colloid Sci. 16, 238 (1961).

(12.) H. P. Grace, Chem. Eng. Commun., 14, 225 (1982).

(13.) I. Delaby, D. Froelich, and R Muller, Macromol. Symp., 100, 131 (1995).

(14.) D. Graebling, R. Muller, and J. F. Palierne, Macromolecules. 26, 320 (1993).

(15.) P. G. Ghodgaonkar and U. Sundararaj, Polym. Eng. Sci, 36, 1656 (1996).

(16.) R. A. Weiss, W. Huh, and L. Nicolais, in High Modulus Polymers, A. E. Zacharides and R. S. Porter. Eds., Marcel Dekker, New York (1988).

(17.) F. P La Mantia, M. Saiu, A. Valenza, M. Pad, and P. L. Magagnini, Eur. Polym. J., 26, 323 (1990).

(18.) F. P La Mantia, M. Saiu, A. Valenza, M. Pad, and P. L Maganini, J. AppL Polym. Sci., 38, 583 (1989).

(19.) D. Beriy, S. Kenig, and A. Siegmann, Polym. Eng. Sci, 31, 451 (1991).

(20.) S. S. Bafna T. Sun. and D. G. Baird, Polymer, 34, 708 (1993).

(21.) J. Lee, J. Jang, Y. Seo, S. M Hong, S. S. Hwang, and K. U. Kim, Int. Polym. Proc., 12, 19 (1997).

(22.) Y. Seo, S. M. Hong, S. S. Hwang, T. S. Park, K. U. Kim, S. Lee, and J. W. Lee, Polymer, 36, 515 (1995).

(23.) W. C. Lee and T. DiBenedetto, Polymer, 34, 683 (1993).

(24.) Y. Chiou, D. Chang, and F. Chang, Polymer, 37, 5653 (1996).

(25.) A. Datta and D. G. Baird. Polymer, 36, 505 (1995).

(26.) Y. Seo, S. M. Hong, S. S. Hwang, T. S. Park, K. U. Kim, S. Lee, and J. W. Lee, Polymer 36, 525 (1995).

(27.) D. R. Paul, J. W. Barlow, and H. Keskkula, in Encyclopedia of Polymer Science and Engineering, Vol. 12, 2nd ed., p. 399. H. F. Mark et al., Wiley, New York (1989).

(28.) H. J. O'Donnel and D. G. Baird, Polymer, 36, 3113 (1995).

(29.) Y. Seo, S. M. Hong, and K. U. Kim, Macromolecules, 30, 278 (1997).

(30.) S. Lee, S. M. Hong, Y. Seo, T. S. Park, S. S. Hwang, K. U. Kim, and J. W. Lee, Polymer, 35, 519 (1994).

(31.) Y. Seo and K. U. Kim, Polym. Eng. Sci., 38, 596 (1998).

(32.) Y. Seo, B. Kim, and K. U. Kim, Polymer, 40, 4483 (1999).

(33.) Y. Seo J. AppL Polym. Sci., 67, 359 (1997).

(34.) R. A. Weiss, W. Huh, and L Nicolais, Polym. Eng. Sci., 27, 684 (1987).

(35.) P. Zhuang T. Kyu, and J. L. White, Polym. Eng. Sci., 28, 1095 (1988).

(36.) B. R. Bassett and A. F. Yee, Polym. Eng. Sci., 30, 10 (1990).

(37.) G. Crevecoeur and G. Groeninckx, Polym. Eng. Sci., 30, 532 (1990).

(38.) N. Ogata, T. Tanaka, T. Ogihara, K. Yoshida, Y. Kondo, K. Hayashi, and N. Yoshida, J. Appl. Polym. Sci., 48, 383 (1993).

(39.) T. Kobayasbi, M. Sato, N. Takeno, and K. Mukaida, Eur. Polym. J., 29, 1625 (1993).

(40.) F. Ide and A. Hasegawa, J. Appl. Polym. Sci., 18, 963 (1974).

(41.) Y. Seo, H. Kim, B. Kim, S. M. Hong, S. S. Hwang, and K. U. Kim, Korea Polym. J., 9, 238 (2001).

(42.) G. Marrucci, in Thermotropic Liquid Crystal Polymer Blends, F. P. La Mantia, ed.. Technomic Publishing, Lancaster, Pa. (1993).

YONGSOK SEO (*)

(*) To whom correspondence should be addressed: ysseo@idst.re.kr.
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Author:Seo, Yongsok; Kim, Hyong-Jun; Kim, Youngjun; Rhee, Hee Woo
Publication:Polymer Engineering and Science
Article Type:Abstract
Date:May 1, 2002
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