Formation, stability, and properties of in-situ composites based on blends of a thermotropic liquid crystalline polymer and a thermoplastic elastomer.
In recent years, numerous studies on two-phase blends of thermoplastic polymers and thermotropic liquid crystalline polymers (TLCPs) have been reported (1-3). Interest in blends containing TLCPs as the minor phase is mainly focused on the fact that under appropriate melt processing conditions, the dispersed TLCP phase can be elongated into fine, oriented fibers which, if preserved by solidification, can reinforce the matrix. Since in this case an oriented fiber/matrix morphology is generated during processing, these blends are often referred to as in-situ composites (4).
In-situ composites usually are generated with conventional processing equipment like single or twin screw extruders. In these processing techniques, the TLCP is often first dispersed as fine (micron sized) droplets, subsequently elongated into fibers in a stretching step and finally frozen in by solidification. The stretching step, typically realized by flow through a conical die and/or by post melt-drawing, is generally considered necessary, not only to form high aspect ratio fibers, but also to induce a high degree of molecular orientation in the fibers. Only in TLCP fibers with a significant level of molecular orientation, the reinforcing effect becomes manifest.
The generation of in-situ composites by conventional, mainly dispersive, processing techniques tends to be highly system and equipment dependent. Some disadvantages of using dispersive mixing equipment for the generation of in-situ composites are:
* Control of the TLCP phase size is relatively difficult.
* Frequent occurrence of a skin-core morphology (5-7) in extruded strands or injection molded parts.
* Not suitable for matrix polymers with no overlap with the relatively high processing temperatures of TLCPs.
For the generation of fiber/matrix morphologies static mixers are much more appropriate than conventional blending equipment. A static mixer operates on the distributive mixing principle of repeated splitting, stretching, and recombining the melt streams of both polymers. In this paper, a coextrusion technique is described where TLCP and matrix are fed separately by two extruders to a static mixer with an adjustable number of mixing elements. The TLCP phase dimensions can be controlled very accurately by varying the number of mixing elements, and a fiber/matrix morphology can be formed directly, without an intermediate droplet/matrix morphology. Additionally, TLCP/thermoplast systems that normally have no overlap in processing temperatures, can be processed with this coextrusion technique. TLCP and matrix polymer can be fed at different temperatures to the static mixer, which itself can, if necessary, be set at a temperature below the actual melting point of the TLCP, exploiting its supercooling tendency (8). The risk of thermal degradation of matrix polymers can thus greatly be reduced. This concept of blending polymers has been used in the past to generate, e.g., multi-layered structures (9-12). Recently Sukhadia et al. (13) reported a similar concept ("dual-extruder mixing method") to generate in-situ composites.
The process of in-situ fiber formation from droplet/matrix dispersions can qualitatively be described using the theory of deformation and break-up of isolated droplets, originally developed by Taylor as early as the 1930s (14, 15). The main parameters governing the process of deformation and break-up of an isolated droplet of viscosity [[Eta].sub.d] suspended in a matrix of viscosity [[Eta].sub.m] are:
* the capillary number Ca = [Tau]R/[Sigma], being the ratio of the deforming (shear-) stress [Tau] imposed by the flow field and the stabilizing interfacial stress [Sigma]/R (where [Sigma] is the interfacial tension and R is the droplet radius).
* the viscosity ratio p = [[Eta].sub.d]/[[Eta].sub.m].
Upon slightly exceeding a critical capillary number [Ca.sub.crit] (where the shear stress overrules the interfacial stress), no stable equilibrium drop shape exists and the droplet will usually break into two daughter droplets and possibly a few much smaller satellite droplets. Grace (16) and others (17), determined [Ca.sub.crit] experimentally under more or less equilibrium conditions using Newtonian model liquids and found a strong dependence on the flow type (shear or elongational) and the viscosity ratio p. In shear flow, for example, [Ca.sub.crit] has a minimum ([Ca.sub.crit] [similar to] 0.5) for 0.1 [less than] p [less than] 1, whereas for p [greater than] 4, [Ca.sub.crit] [approaches] [infinity], indicating that a droplet under these conditions can hardly be deformed or broken. However, in elongational flow deformation and break-up always seem possible.
Upon a sudden gross exceeding of the critical capillary number (Ca [much greater] [Ca.sub.crit], supercritical flow) a droplet will deform affinely with the matrix into a long slender body or thread. Elemans (11) has shown that affine deformation in simple shear flow of Newtonian droplets in a Newtonian matrix occurs at Ca [greater than] 2[Ca.sub.crit]. Recently, Janssen (18) determined the condition for affine deformation in 2D-elongational flow to be Ca [greater than] 5[Ca.sub.crit].
The main break-up mode in case of a long cylindrical thread is the Rayleigh distortion mechanism (19), in which the thread exhibits interfacial tension driven wavelike distortions causing break-up into a line of droplets. The break-up of viscous threads embedded in a quiescent or flowing viscous matrix is described theoretically by various authors (20-23). In general, break-up of threads during flow is retarded compared to break-up under quiescent conditions, because the flow of the matrix fluid extinguishes developing distortions on the thread/matrix interface.
Generating in-situ composites in conventional blending equipment takes place in two steps. In the initial stage of the melt-blending process, Ca [much greater than] [Ca.sub.crit], the TLCP phase domains (size of the order of mm) will be extended affinely into long slender bodies. During this distributive mixing stage (passive interfaces, no break-up) the slender bodies continuously decrease in diameter. Finally, the interfacial stress will become of the same order as the applied stress and dispersive mixing (active interfaces, with break-up) starts to play a role. The slender bodies break up into smaller droplets, according to the Rayleigh distortion mechanism. Depending on the local capillary number these droplets may be deformed and broken up again. Apart from deformation and break-up, which lead to finer morphologies, there is a coarsening effect due to coalescence. At the end of the first step the capillary number is usually of the order of [Ca.sub.crit].
In the subsequent stretching step, the capillary number must significantly exceed [Ca.sub.cit] to provide affine deformation of the TLCP droplets into high aspect ratio fibers. In particular, in the last stage of morphology development, the time effects of several microtheological processes (deformation, break-up, and coalescence) are important, since a non-equilibrium morphology has to be frozen in by solidification. The stability of the non-equilibrium fiber/matrix morphology of in-situ composites above the melting point of the TLCP has received little attention. Earlier work in our laboratory, however, has demonstrated that TLCP fibers (diameter [approximately equal to] 5-10 [[micro]meter]) embedded in various thermoplastic matrices (e.g., PES, PP (24), SEBS (25)], are highly unstable in the molten state and rapidly break into droplets. To obtain a fiber/matrix morphology the fiber break-up times, evidently, must be larger than the residence and solidification times in the final stages of processing.
The fiber formation step, which needs Ca [greater than] [Ca.sub.crit], is thus favored by large phase domains, high stress levels, and a small interfacial tension. The influence of the phase size is reflected in the experimental observation that a certain minimum TLCP concentration (and thus TLCP phase size) is necessary for fiber formation (26-29). The average droplet size, as determined by the dynamic equilibrium between dispersive mixing and coalescence, increases with increasing concentration dispersed phase and consequently fiber formation is easier.
An often claimed condition for fiber formation is that the viscosity of the dispersed (TLCP) phase must be smaller than that of the matrix (p [less than] 1) (30-33). The requirement that p [less than] 1 for fiber formation, however, seems only valid under shear flow conditions. In elongational flow, deformation is still possible. An explanation for the reported impossibility to form fibers at p [greater than] 1 may be that Ca (which in the final dispersive mixing stage is close to [Ca.sub.crit]) during subsequent stretching, does not exceed [Ca.sub.crit] sufficiently to obtain affine deformation into fibers.
However, dispersive mixing, involving the formation of droplets, is redundant if the desired morphology is a fiber/matrix morphology. A fiber/matrix morphology can be realized by a single distributive mixing process. In conventional equipment it is difficult, if not impossible, to separate distributive and dispersive mixing.
This paper reports the results of blending a high melting TLCP (Vectra A900) with a polyether-ester thermoplastic elastomer (Arnitel em630), using a co-extrusion technique involving separately feeding the components to a Ross static mixer. The study was aimed at: a) forming in-situ composites by a single distributive mixing process, in particular, the influence of the number of mixing elements and extrudate drawing on morphology and tensile properties of the blends, b) feeding TLCP and matrix polymer at different temperatures, thus enlarging the processing window for in-situ composites, and c) Investigating the break-up process of TLCP fibers in the matrix at temperatures above their melting point, in particular break-up times in comparison with residence and solidification times.
The Vectra/Arnitel system was chosen for the following two reasons. First, the processing temperatures of Arnitel em630 and Vectra A900 do not overlap: blending on a single screw extruder was not possible due to the limited thermal stability and the low viscosity of Arnitel at the processing temperature of the TLCP. Second, the tensile properties of Arnitel and Vectra differ greatly, enabling easy detection of the reinforcing effect of the TLCP, particularly at low concentration.
The TLCP used is an aromatic random copolyester consisting of 73% 4-hydroxy-benzoic acid (HBA) and 27% 2-hydroxy-6-naphtoic acid (HNA) supplied by Hoechst Celanese as Vectra A900. It was dried for 4 h at 180 [degrees] C under nitrogen atmosphere and stored in a vacuum oven at 80 [degrees] C before further use. The thermoplastic elastomer used is Arnitel em630 supplied by DSM. Arnitel is a semicrystalline block copolymer consisting of 25% poly-oxytetramethylene (POTM) acting as the soft amorphous segment and 75% poly-butylene terephtalate (PBT) acting as phase-separated semicrystalline crosslinks (34). Arnitel was oven dried at 80 [degrees] C and stored before further use.
The coextrusion die is shown schematically in Fig. 1. A Collin laboratory extruder equipped with a transport screw (D = 20 mm, L/D = 20) was used for feeding the thermoplastic elastomer to the mixing section. The TLCP melt is injected centrally in the matrix stream just in front of the mixing section, by a Handle laboratory extruder (D = 17 mm, L/D = 20). The mixing section consists of a sequence of 15 mm diameter Ross mixing elements, the number of which can be varied. The four-way Ross static mixer was chosen because of the low retained volume (17% of the equivalent open pipe), relative short residence time and high shear rates (35). After mixing, the melt enters a conical die with a 60 [degrees] entrance angle, a diameter of 3 mm, and an L/D ratio of 10. The conical die is mounted directly behind the sequence of Ross mixers, in order to avoid creating a cavity, which could act as a "dead zone" (low shear rate zone) in the processing equipment. Subsequently, the extruded strand is stretched and quenched in a water bath. The distance between die and water bath was held constant at 30 cm in these experiments. Different draw ratios were obtained by varying the speed of a take-up device. The draw ratio (DR) can be obtained from the ratio of cross sectional areas of the die and the extrudate: DR = ([D.sub.die]/[D.sub.ex]). The draw ratio of the extruded strands ranged from 1 to 15.
Maximum barrel temperature settings during processing were as follows: matrix extruder: 260 [degrees] C, TLCP extruder: 300 [degrees] C. The screw rotational speed of the matrix extruder was 10 rpm. Blends with different compositions were obtained by varying the screw speed of the TLCP extruder.
The blend composition of the extrudates was determined by an extraction method. Extrudates of known weight were placed in a closed sample holder made of porous metal. The Arnitel matrix was removed by Soxhlet extraction in 1,1,2,2-tetrachloroethane for 24 h. After extraction the undissolved material was vacuum dried at 80 [degrees] C for several days and weighed again. All values are the result of averaging eight measurements. The blend composition will be given in terms of volume fraction Vectra in this paper. The ambient density values of Vectra A900 and Arnitel em630 used for the calculations were 1.37 and 1.23 kg/[dm.sup.3], respectively.
The optimum temperature of the mixing section was 270 [degrees] C for our system. Maintaining a mixing at 280 [degrees] C and higher, resulted in a melt strength that was too low to obtain smooth extrudates. Mixing at 260 [degrees] C and lower resulted in irregularly shaped extrudates, possibly because the Vectra phase was partly solidified.
Three series of coextrusion experiments were performed, using 8, 11, and 14 mixing elements. The total mass flow rate ([Q.sub.m]) during the coextrusion experiments was [approximately]10 g/min. Assuming a melt-density of 1 kg/[dm.sup.3], a rough estimate of the shear rate at the wall in the channels of the Ross mixing elements (a Ross element contains 4 cylindrical channels, with a diameter of 2.73 mm) at this mass flow rate is: [Mathematical Expression Omitted].
The total residence time in the mixing section (with 11 elements) and the conical die at the above mentioned flow rate was about 60 s. This value was determined experimentally by injecting a small amount of Vectra (with the TLCP extruder) into the Arnitel stream. As the residence time was taken the time between the moment of injection and the first appearance of Vectra at the die exit, which is in fact the minimum residence time. These tests also demonstrated that the residence time distribution was quite narrow, indicating that no "dead spaces" were present. The experimental residence time of 60 s, agrees well with the calculated residence time of 61 s, which follows from the total retained volume (= retained volume of the mixing elements + volume between injection point and mixing elements + die volume), which is 10.2 [cm.sup.3]. The calculated residence time in the 11 mixing elements alone (retained volume 4.2 [cm.sup.3]) is 25 s.
Two types of experiments were performed to determine the stability of the morphology of the blends above the melting point of the TLCP:
1. Annealing experiments, which were carried out by placing small pieces of coextruded strands with a DR of 4 on a glass slide in a Mettler FP82 hot stage. Then, the sample was heated (20 [degrees] C/min) to a temperature above the melting point of the TLCP and held at this temperature for periods of time. After annealing the sample was rapidly removed from the hot stage and cooled to room temperature in ambient air.
2. Capillary instability experiments, carried out on a Linkham THM600 hot stage under an optical microscope (Jenapol). Samples were prepared by positioning a TLCP thread (diameter 5-15 [[micro]meter], drawn from a molten granule) between two sheets (thickness 100-200 [[micro]meter], compression molded) of Arnitel. Subsequently, the sample was heated in the hot stage to the desired temperature and capillary instabilities developing on the thread/matrix interface were monitored. Details of the procedure will be given elsewhere (24).
Thermal characterization of the polymers was carried out with a Perkin Elmer DSC-7 calorimeter. All heating scans were performed with a heating rate of 20 [degrees] C/min. For Vectra A900 additional cooling scans were made from 300 [degrees] C to 50 [degrees] C with various cooling rates, to determine the effect of cooling rate on the liquid crystalline to crystalline transition temperature.
Rheological measurements on the pure polymers were performed using a single-barrel and a twin-barrel capillary rheometer. The twin-barrel capillary rheometer (Rosand RH7-8/2) was equipped with two 0.5 mm diameter capillary dies (entrance angle: 90 [degrees]) with L/D ratios of 4 and 40. On the viscosity data, Bagley and Rabinowitsch corrections were applied. The single-barrel capillary rheometer was equipped with a 1.5 mm diameter capillary die (entrance angle: 90 [degrees]) with L/D ratio of 50. The Bagley correction was found to be negligible in this case.
In addition, oscillatory shear measurements on Vectra A900 were conducted during cooling from the melt, using a Rheometrics RMS-800 mechanical spectrometer. The samples for these experiments were prepared by compression molding Vectra A900 granules into discs at 310 [degrees] C for 5 min, followed by immediate cooling in ambient air. The dynamic measurements were performed in the parallel plate mode with a plate radius of 12.5 mm and a gap of 1.2 mm, using a strain of 5%. During the experiments the dynamic properties ([[Eta].sup.*], G[prime], and G[double prime]) at a frequency of 10 rad/s were monitored as a function of temperature upon cooling from 300 [degrees] C. The cooling rate was 4 [degrees] C/min.
The morphology of the blends was studied by scanning electron microscopy (Philips XL 20). SEM-samples were prepared by cryogenic fracture of extrudates in liquid nitrogen both perpendicular and parallel to the extrusion direction. Gold-coating was applied to improve conductivity of the samples. The microscope was equipped with an on-line size-measurement facility, which was used to estimate the diameter of the dispersed TLCP phase.
Tensile modulus and strength of the extrudates as a function of draw ratio were measured on a Zwick 1445 tensile tester with a constant crosshead speed of 10 mm/min using an initial gauge length of 46 mm. Nearly all specimens, except those of low Vectra content, broke at the clamps as a result of stress concentrations. Tensile strength measurements were therefore carried out in a separate run, where the extrudates were secured between cork-filled rubber pads in the clamps. By adjusting the clamp pressure (higher pressures were necessary for samples with higher Vectra content), breaking of the samples at the clamps could be avoided. For extrudates exhibiting a yield behavior (pure Arnitel and strands with a low fraction TLCP), the yield strength is taken as the tensile strength. For all other blends the strength at break was taken as the tensile strength.
Dynamic mechanical properties of extrudates with a DR of 9 as a function of temperature were determined by a dynamic mechanical analyzer (Perkin Elmer DMA 7). Measurements were performed in the tension mode with a frequency of 1 Hz and a heating rate of 2 [degrees] C/min in the temperature range of -125 [degrees] C to 150 [degrees] C.
Information on the molecular orientation of the TLCP fibers in some extrudates of different draw ratios, was obtained by Wide Angle X-ray Diffraction (WAXS). X-ray diffraction patterns were made with a flat film Kiessig pinhole camera, using point collimated Ni-filtered CuK[Alpha] radiation. The distance between sample and film was 100 mm.
Thermal Properties of the Polymer
DSC analysis of Vectra A900 revealed a Tg [similar to] 100 [degrees] C and a melting peak at 280 [degrees] C, which corresponds to the crystalline to nematic phase transition (36). For Arnitel em630, the only transition that could be determined with DSC was a melting endotherm at 210 [degrees] C, which corresponds to the melting point of the PBT segments. With DMA, a single broad transition is found in the temperature region -55 [degrees] C to 20 [degrees] C and a small transition around 80 to 100 [degrees] C [ILLUSTRATION FOR FIGURE 9 OMITTED]. The first transition is probably caused by an overlap of the glass and melting transitions of the POTM segments (37), while the second may correspond to the glass transition of the PBT segments.
Some characteristic thermograms of Vectra A900, upon cooling from 300 [degrees] C with different cooling rates, are shown in Fig. 2. As can be seen, under quiescent conditions the TLCP can be supercooled significantly ([approximately]40 [degrees] C with a cooling rate of 5 [degrees] C/min) before actual solidification takes place.
Rheology of the Polymers
The shear viscosity vs. shear rate behavior of Arnitel at 240, 250, and 260 [degrees] C, as determined with the capillary rheometer, is shown in Fig. 3. In the shear rate range investigated, this polymer shows Newtonian flow behavior at 250 and 260 [degrees] C. At 240 [degrees] C, a slight shear thinning is observed. Because of thermal degradation from long residence times (10 to 15 min) in the barrel of the capillary rheometer, no reliable results could be obtained for processing at 270 [degrees] C. The viscosities obtained at the lower temperatures were therefore extrapolated, by means of an Arrhenius-type equation, to estimate the viscosity at 270 [degrees] C. This results in: [[Eta].sub.270 [degrees] C] [approximately equal to] 190 Pas.
The flow curves of Vectra A900 at 290 and 300 [degrees] C are presented in Fig. 4. Shear thinning is observed over the entire range of shear rates investigated, which agrees with results of earlier investigators, e.g. Wissbrun et al. (38). It is known that the TLCP melt at temperatures [less than]300 [degrees] C is liable to recrystallization (36), which affects the flow. In fact, only at temperatures [greater than]300 [degrees] C the flow of the TLCP is stable. Relatively stable nematic melts at [less than]300 [degrees] C can be obtained by preheating at higher temperatures, as described by Lin et al. (39). Nevertheless, the results at 290 [degrees] C are used to have at least an indication of the flow of the TLCP at temperatures close to the mixing temperature.
A typical example of a cooling experiment performed with the plate-plate rheometer is shown in Fig. 5. Upon cooling from 300 [degrees] C, the complex viscosity [[Eta].sup.*] shows an abrupt increase [less than]270 [degrees] C. At this point the supercooled melt starts to solidify. Again these results indicate, as do the DSC measurements, the possibility of processing Vectra below its actual melting poInt once it is molten. The supercooling of the TLCP melt under dynamic conditions is much smaller than under quiescent conditions in the DSC.
Figure 6 shows the morphology of the blends prepared usIng 8 mixing elements. A stratified morphology of continuous Vectra layers (or ribbons), dispersed in the Arnitel matrix, is observed. In a blend containing 11 vol% Vectra [ILLUSTRATION FOR FIGURE 6A OMITTED], the thickness of these layers is [approximately]4 [[micro]meter]. Between the Vectra layers numerous micron-sized fibers can be observed [ILLUSTRATION FOR FIGURE 6D OMITTED]. Stratified morphologies are quite common when static mixers are used for mixing two or more separate polymer melt streams (9, 11, 12). SInce TLCPs in general display extremely low permeabilities for gasses and liquids (40, 41), the resultant morphology of dispersed TLCP layers in a thermoplastic matrix may be interestIng for gas barrier films.
The morphologies of the blends made by 11 mixing element coextrusion are shown in Figs. 7a through 7h. Already at Vectra concentrations of 6 vol% [ILLUSTRATION FOR FIGURES 7A AND 7B OMITTED] high aspect-ratio fibers are formed in the extrusion direction. At DR = 15 the average fiber diameter is [approximately] 1 [[micro]meter]. The fibers appear to be uniformly distributed over the extrudate cross-section. Higher Vectra concentrations (38 vol%, [ILLUSTRATION FOR FIGURES 7C THROUGH 7H OMITTED]), also lead to a distinct fiber/matrix morphology in which the fibers appear to be continuous over the length of the extrudate.
SEM micrographs of blends made with 14 mixing elements are shown in Fig. 8a through 8i. In this case, a distinct skin-core morphology has been generated [ILLUSTRATION FOR FIGURE 8A THROUGH 8C OMITTED]. In the core of the extrudate, only moderately extended Vectra droplets are observed, while in the skin high aspect-ratio fibers are present. The skin-core morphologies are more pronounced at relatively thick extrudates and low Vectra content ([ILLUSTRATION FOR FIGURES 8A THOUGH 8C OMITTED], 17 vol% Vectra, DR = 3), but even at higher Vectra content and at high DR ([ILLUSTRATION FOR FIGURES 8G THROUGH 8I OMITTED], 33 vol% Vectra, DR = 15), droplets are still present in the core region.
Table 1 presents the average, minimum, and maximum TLCP fiber diameters in the extrudates of draw ratios 3 and 15, made with 11 mixing elements, derived from images depicted in Fig. 7e.
This Table shows some trends regarding the morphology and the influence of the final drawing step, outside the die.
* The minimum fiber diameter appears to be independent of both the Vectra concentration and the draw ratio.
* The average and maximum fiber diameters increase with higher Vectra concentrations.
* The average and maximum fiber diameter decrease with higher draw ratios.
* The ratio between the average TLCP fiber diameters at DR = 3 and DR = 15 (for a given Vectra content), increases with higher Vectra content, indicating that the TLCP fibers are easier deformed when present at higher concentration. However, this ratio is still significantly smaller than the ratio between the macroscopic diameters of the extrudates with DR = 3 and DR = 15, which is 2.2. Hence it appears that the deformation of the TLCP fibers in the blend, during the final drawing step, is less than affine. This is not surprising, considering the differences in viscosities and solidification temperatures of the blend components during the non-isothermal drawing process.
Dynamic Mechanical Properties
The effect of the TLCP on the dynamic mechanical behavior of the extrudates (11 mixing elements, DR = 9) is illustrated in Fig. 9, where the storage modulus (E[prime]), measured in the tension mode, in the temperature range between -125 [degrees] C and 150 [degrees] C is shown. The loss modulus (E[double prime]) could not be determined accurately in this deformation mode. As is clear from Fig. 9, the storage modulus increases significantly with increasing Vectra content at all temperatures. At 25 [degrees] C, the addition of 15 vol% Vectra leads to a value 9 times higher than that of pure Arnitel. Also, the rubbery plateau of Arnitel disappears with the addition of (even fairly small amounts) Vectra. No significant change with the addition of Vectra can be observed in the small transition at 80 to 100 [degrees] C of Arnitel, which can possibly be attributed to the Tg of the PBT segments.
The addition of Vectra to Arnitel as a reinforcing agent also becomes manifest in the increase in tensile modulus and strength, measured by tensile testing of extrudates. Figures 10 and 11 are typical for the tensile modulus and strength of the extrudates, respectively, as a function of DR. Both modulus and strength increase with the DR, as a result of increased molecular orientation of the TLCP fibers, induced by the final drawing process. At low Vectra content, modulus and strength appear to level off to a constant value. For blends containing [greater than]9 vol% Vectra, no plateau level is reached at the highest DR of 15.
Table 2 shows the tensile modulus/strength of Vectra/Arnitel blends of, more or less, comparable compositions and DR = 15, generated with 8, 11, and 14 mixing elements. The tensile modulus and strength of pure Arnitel are 0.3 GPa and 16 MPa, respectively. This Table shows that the use of 8 and 11 mixing elements results in more or less the same increase in tensile properties. However, the reinforcing effect is considerably less using 14 mixing elements as concluded from a comparison of the values at 15 and 17 vol% Vectra, respectively, for 11 and 14 elements. Apparently, the continuity of the TLCP phase is important for the mechanical properties.
The level of molecular orientation in the TLCP fibers of some extrudates of different DR was studied qualitatively with wide-angle X-ray scattering. The diffraction [TABULAR DATA FOR TABLE 1 OMITTED] pattern of pure Arnitel (DR = 10), shows concentrical circles, indicative of a crystalline sample with no preferential orientation [ILLUSTRATION FOR FIGURE 12A OMITTED]. The diffraction patterns of the blends (15 vol% Vectra/Arnitel, 11 mixing elements, DR = 3 and 15) show an additional pair of equatorial reflections, resulting from the oriented TLCP fiber phase [ILLUSTRATION FOR FIGURE 12B AND C OMITTED]. More quantitative information on the molecular orientation in the TLCP fibers can be obtained by calculating the Hermans orientation factor from an azimuthal densitometer scan over the equatorial (110) TLCP reflection. The Hermans orientation factor S is defined as:
S = 1 - 3/2<[sin.sup.2] [Beta]> (1)
where [Beta] is the angle between the director of the liquid crystalline domains and the fiber axis. A value of S = 1 means perfect orientation in the fiber direction, S = 0 means random orientation of the liquid crystalline domains. The term <[sin.sup.2] [Beta]> was calculated from a Pearson VII fit of the measured azimuthal intensity distribution (42). The calculated values of the orientation factor for the 15 vol% Vectra/Arnitel blends of DR = 3 and 15, are 0.61 and 0.72, respectively, indicating that the molecular orientation in the TLCP fibers has increased with extrudate DR. This agrees with the observed increase in mechanical properties.
A typical example of a capillary instability experiment is presented in Fig. 13. An isolated Vectra thread (diameter 9.4 [[micro]meter]) embedded in an Arnitel matrix, at 300 [degrees] C, breaks up into droplets in [approximately]70 to 80 sec. The fragmentation process proceeds by two mechanisms. First, break-up takes place through the growth of rather irregular Rayleigh distortions. Subsequently, the remaining thread fragments exhibit, depending on the shape, either Rayleigh type distortion growth or retraction to a sphere. In a forthcoming paper (24) capillary instability experiments with TLCP threads embedded in various thermoplastic matrices will be discussed in more detail.
Table 2. Tensile Modulus and Strength of Vectra/Arnitel Blends of Draw Ratio 15 of Comparable Compositions, Generated With 8, 11, and 14 Mixing Elements. number of Modulus Strength mixing vol% (GPa) (MPa) elements morphology Vectra DR = 15 DR = 15 8 'stratified' 23 5.1 90 34 9.5 185 11 'fiber/matrix' 15 2.4 73 38 10.2 205 14 'skin-core' 17 1.7 51 33 4.6 67
The effect of annealing above the melting point of the TLCP on the morphology of a coextruded 15 vol% Vectra/Arnitel strand (DR [approximately equal to] 4, 11 mixing elements) is shown in Fig. 14. Mere heating of the blend to 290 [degrees] C with 20 [degrees] C/min, followed by rapid cooling in air, results in a complete rearrangement from a fiber/matrix [ILLUSTRATION FOR FIGURE 14A OMITTED] into a droplet/matrix morphology [ILLUSTRATION FOR FIGURE 14B OMITTED].
In several coextrusion experiments, the influence of the cooling/quenching conditions on the morphology of the extruded strand was investigated, by reducing the distance between the die and the water bath. Figure 15 shows SEM micrographs of the core region of 28 vol% Vectra/Arnitel extrudates, prepared by 14 mixing element coextrusion with a die-water bath distance of 30 and 2 cm, respectively. In extrudates frozen-in after 30 cm, significantly more droplets and fragmented fibers are observed [ILLUSTRATION FOR FIGURE 15A OMITTED] than in those frozen-in after 2 cm [ILLUSTRATION FOR FIGURE 15B OMITTED], indicating that the break-up process can also take place outside the die.
Formation and Stability of In-situ Composites
As is demonstrated here, the static mixer with separate feeding provides a possibility to generate in-situ composites from TLCP/thermoplast systems, which normally have little or no overlap in processing temperatures. In our specific system, blending the components in one extruder is not even possible. The coextrusion method turns out to have some distinct advantages over conventional processing techniques for generating in-situ composites. First, it is possible to control the mixing intensity by the number of mixing elements. This is clearly shown in the strong dependence on the number of mixing elements of the morphology of the coextruded Vectra/Arnitel blends, ranging from a stratified morphology (8 elements, [ILLUSTRATION FOR FIGURE 6 OMITTED]), via a continuous fiber/matrix morphology (11 elements, [ILLUSTRATION FOR FIGURE 7 OMITTED]) to a skin-core morphology (14 elements, [ILLUSTRATION FOR FIGURE 8 OMITTED]).
Second, the viscosity ratio can be set more favorably (p [less than] 1) for in-situ fiber formation, as mentioned also by Sukhadia et al. (13), by feeding the TLCP melt at high and the matrix at low temperature to the mixing section. Based on a shear rate of 21 [s.sup.-1], calculated above, an estimate of the viscosity ratio, using Figs. 2 and 3, is: p = [[Eta].sub.Vectra](300 [degrees] C)/[[Eta].sub.Arnitel](270 [degrees] C) [approximately] 1.5. This estimate of p represents a lower limit, because the actual mixing temperature will be [less than] 300 [degrees] C (the temperature at which the TLCP melt is injected into the Arnitel matrix of 270 [degrees] C). The temperature of the Vectra phase during mixing will decrease, while the temperature of the Arnitel phase will increase somewhat as a result of temperature leveling and thus the viscosity ratio will be higher. Hence, it appears possible with the present coextrusion technique, to generate fiber/matrix morphologies under conditions where the viscosity of the matrix is lower than that of the dispersed TLCP phase component (p [greater than] 1). We suppose that this is due to the mainly distributive mixing action of the static mixer, where the viscosity ratio is of less importance in determining whether fibers are formed. This is in contrast to the mainly dispersive mixing techniques (which are more common so far) and where the necessity of p [less than] 1 for fiber formation is reported.
As far as the in-situ fiber formation in this specific system is concerned, under the chosen processing conditions an optimum number of 11 mixing elements is found. Sukhadia et al. (13) report that 18 Kenics mixing elements are necessary for their "dual-extruder mixing method" to obtain a fiber/matrix morphology. The difference is caused by the fact that the Ross ISG mixer generates twice as many striations (viz. n [multiplied by] [4.sup.e], where n is the number of input streams, e is the number of elements) per element than the Kenics static mixer (n [multiplied by] [2.sup.e]). The mixing effectivity of the Ross static mixer, however, is still less than would be expected on the basis of the assumption of perfect splitting and recombining the melt streams. When perfect splitting and recombination is assumed, a number of 7 or 8 mixing elements would be expected to reduce the striation thickness to 0.1-1 [[micro]meter] in our case. A possible explanation for the difference between experimental and theoretical thicknesses might be that stream splitting of the dispersed TLCP phase is less perfect when its dimension becomes too small compared to the channel diameter and therefore flow through the channels occurs without any splitting at all. Also, coalescence effects in recombining the melt streams may lead to the larger phase dimensions.
Further characteristics of the Ross ISG mixer are its relative short residence time and high shear rates, which is an advantage, particularly when using low melting or thermally unstable thermoplastics as a matrix polymer. Sukhadia (43) reports typical residence times of 65 and 139 s at flow rates of 88 and 37 g/min, respectively, for the dual-extruder mixing method, while we find [approximately]60 s at a flow rate of 10 g/min. The corresponding shear rates are 4.7 [s.sup.-1] at 88 g/min (43) for the dual-extruder mixing method compared to 21 [s.sup.-1] at 10 g/min in the Ross mixing section. The pressure drop is higher in our coextrusion process: 35 bar compared to 17.6 bar in case of the dual-extruder mixing process (43).
After passing the static mixer section the blend is mainly subjected to elongational flow in the conical die (part E in Fig. 1) and in the strand drawing process outside the die (except in the capillary section of the die). This results in further deformation of the TLCP fibers formed in the static mixer (e.g. in case of 11 mixing elements) and, moreover, induces the necessary molecular orientation within the fibers. The orientation induced in the entrance of the die, however, may relax in the capillary section of the die. Turek et al. (44), investigated the orientation development in Vectra A950 (having the same composition as Vectra A900), extruded through capillary dies of different L/D ratios. The flow in the capillary section of the die permits relaxation of the orientation developed in the die entrance and is, therefore, dependent on the L/D ratio. The decrease in orientation (determined by WAXS) as a function of capillary L/D ratio, was modeled with a simple exponential decay function, resulting in a characteristic average relaxation time of 0.5 s. The residence time of the melt in the capillary section of the die is [approximately]1.4 s. Hence, it is likely that the main part of the orientation development within the fibers and thus reinforcement of the blend, takes place in the drawing process outside the die, which will be discussed in the Mechanical Properties Section.
The fiber/matrix morphology in the present system is highly unstable in the molten state. The fibers fragment into droplets usually by a combination of Rayleigh distortions and retraction [ILLUSTRATION FOR FIGURE 13 OMITTED]. For Vectra threads of [approximately]10 [[micro]meter] diameter the time scale for Rayleigh distortions to cause break-up is less than a minute. The time for break-up, [t.sub.b], by growth of Rayleigh distortions, as derived for viscous threads embedded in a viscous matrix under quiescent conditions, is (45):
[t.sub.b] = 1/q ln (0.82 [R.sub.0])/[[Alpha].sub.0]) (2)
in which the distortion growth rate q [[s.sup.-1]], is given by:
q = [Sigma] / 2[[Eta].sub.m][R.sub.0] [Omega]([Lambda], p) (3)
[t.sub.b] = Time to the point of break-up [s].
[R.sub.0] = initial thread radius [m].
[[Alpha].sub.0] = Initial distortion amplitude [m].
[Sigma] = Interfacial tension [N/m].
p = Viscosity ratio [[Eta].sub.d]/[[Eta].sub.m], [[Eta].sub.d], and [[Eta].sub.m] are the dispersed and matrix phase viscosity, respectively.
[Lambda] = Wavelength of a sinusoidal distortion [m].
[Omega]([Lambda], p) = Dimensionless growth rate, a function of [Lambda] and p (20).
Taking break-up in the viscous case as a first approximation, the ratio of break-up times of a 10 and 1 [[micro]meter] diameter thread with an initial distortion, [[Alpha].sub.0], of 0.001, 0.01, and 0.1 [[micro]meter] is 14, 16, and 27 respectively. This implies that the break-up times of the TLCP fibers in the extruded blends (with a diameter of say 1 [[micro]meter]) probably lay in the order of a few seconds (not strongly dependent on the value of the initial distortion [[Alpha].sub.0]). This is more or less confirmed in the annealing experiments with the extruded strands [ILLUSTRATION FOR FIGURE 14 OMITTED].
The occurrence of the skin-core morphology as observed in the experiments using 14 mixing elements may thus be the result of an unfavorable balance between the residence time/solidification time and break-up times in the final stage of the blending process. The residence time in the capillary section of the die is [approximately]1.4 s, which is probably enough for the submicron fibers in the low shear rate core region of the capillary to get seriously distorted or even broken. Fibers near the capillary wall are stabilized against break-up by the relatively high shear rates. As is demonstrated in the work of e.g. Khakhar et al. (23), break-up of threads during flow is retarded compared to break-up under quiescent conditions.
Outside the die break-up can also take place as a result of a slow cooling process. The residence time between the die and the water bath and the time needed for solidification, are larger than the estimated break-up time of the TLCP fibers in the core of the extrudate. The residence time ([t.sub.r]) of the extruded strand in ambient air as a function of the die-to-water bath distance, l, and volumetric flow rate, [Q.sub.v], is given by:
[t.sub.r] = [Pi] / [Q.sub.v] [integral of] [r.sup.2](z)dz between limits t and 0 (4)
where r(z) is the extrudate radius at axial distance z from the die-exit. Assuming, for example, a hyperbolic decrease of r(z), the extrudate radius between the die and the water bath can be expressed as (7, 46):
r(z)/[r.sub.0] = 1 - a[z/(1 + z)] (5)
where [r.sub.0] is the initial extrudate radius (=die exit radius) and a is a constant determined by the final extrudate radius. Expressing this constant in terms of the DR and the distance (l) between the die and the water bath, gives:
a = (1 - 1 / [square root of DR]) (1 + l / l) (6)
In Table 3, the calculated residence times as a function of the draw ratio are shown under the processing conditions presented in the Experimental Section. In particular at low draw ratios the residence time between the die and water bath is of the same order of magnitude as the fiber break-up times. By similar considerations, e.g. by following the procedure presented by Blizard et al. (46), it can be shown that at least the core of the cooling extrudate is not solidified (assuming that the Vectra phase solidifies at [approximately]250 [degrees] C) within the residence times indicated in Table 3, so fiber break-up may go on outside the die. The experiments where the residence time before quenching was decreased by reducing the distance between die and water bath Indeed produce an Indication that part of the break-up can take place while the strand is cooling in ambient air.
The difference in morphology with the strands using 11 mixing elements, where skin-core morphologies are absent, is probably due to the fact that the fibers generated are somewhat larger in diameter and therefore more stable and able to "survive" the processes mentioned above. In fact, by using 14 mixing elements the generated fibers are too thin to remain stable. The optimum number of mixing elements for fiber formation is that, which reduces the mixing scale to a level that break-up into droplets is just avoided. It is probably because of this fact that there appears to be a certain minimum achievable TLCP phase domain diameter in the blends more or less independent of TLCP content and draw ratio (see Table 1). Assuming that these phase domains are indeed TLCP fibers, then fibers with a smaller diameter would break up under the chosen processing conditions into droplets with a larger diameter. We thus conclude from our coextrusion and break-up experiments that skin-core morphologies, as frequently observed in-situ composite generation, originate from break-up of fibers in the core region of the extruded and cooling melt. Some authors (6, 47), have suggested that skin-core morphologies are caused by a migration process during blending, in which the lower viscous TLCP phase migrates to high shear rate regions (e.g. extruder barrel wall or the capillary wall of extrusion dies). A higher TLCP concentration in the skin region results in easier fiber formation there. It is unlikely, however, that this occurs within the blending time in the Ross static mixer section, which causes radial mixing as well.
Table 3. Residence Time of the Extruded Strand Between the Die and Water Bath. Calculations Performed With: I = 30 cm, [r.sub.0] = 0.15 cm and [Q.sub.v] = 0.165 c[m.sup.3]/s ([approximately equal to]10 g/min). DR [t.sub.r] (s) 3 4.9 9 2.1 15 1.5
In general, the instability of in-situ composites in the molten state and the short times needed for break-up presents a serious limitation for use of TLCPs as in-situ generated reinforcing agents in blends. Particularly, problems can be expected in processing/shaping techniques requiring long cooling/solidification times such as Injection molding or extrusion of thick parts. The stability of in-situ composites can in principle be enhanced by the addition of suitable compatibilizers that lower the interfacial tension and retard the break-up process. A problem that may arise then, which has also been pointed out by Baird et al. (3), is that in systems with a low interfacial tension dispersive mixing easily leads to TLCP phase domains that are too small and therefore too stable to be elongated into high aspect ratio fibers. A flexible distributive mixing technique, however, such as the coextrusion method described here, in which the TLCP phase dimensions can be controlled flexibly, should be able to cope with this problem.
Properties of In-situ Composites
Dynamic Mechanical Properties
One of the aims of characterizing the extrudates with dynamic mechanical measurements was to investigate whether there was any interaction between Vectra and Arnitel. In particular, the polyester segments of both polymers might show some interaction, which could be detected by a shift of the Tg for these segments. However, no clear change in the small transition [approximately]90 [degrees] C, which we attributed to the Tg of the PBT segments of Arnitel, was observed with the addition of Vectra [ILLUSTRATION FOR FIGURE 9 OMITTED]. From the dynamic mechanical properties alone, however, it appears impossible to draw any reliable conclusion regarding the incompatibility in this system, because the Tgs of Vectra and the PBT-segments (partly) overlap.
The addition of Vectra to Arnitel results in an increase of the tensile modulus and strength. The level of reinforcement is highly dependent on the morphology and the draw ratio of the extrudates. The effect of the morphology becomes apparent when the tensile properties of the blends prepared with 8 or 11 mixing elements are compared with those of 14 elements (Table 2). Using 14 mixing elements, the observed increase in tensile properties (at more or less similar Vectra content and draw ratio) is significantly less than in case of 8 and 11 mixing elements. This is probably due to the fact that in case of 8 and 11 elements the Vectra phase is present as continuous layers or fibers in the blends. In the blends prepared with 14 mixing elements (skin-core morphology), a large fraction of the dispersed Vectra phase is present as droplets, in particular at lower draw ratios.
Higher draw ratios have two effects, which both lead to increased tensile properties of the in-situ composite strands. First, the aspect ratio of the TLCP fibers in the blends increases as was demonstrated by the measurements of the diameter of the dispersed fibers (see Table 1). Second, higher draw ratios result in an increased molecular orientation in the TLCP fibers, as was shown by WAXS, leading to higher fiber modulus and strength.
The longitudinal modulus of in-situ composites is often modeled with the Halpin-Tsai Equation (48):
E / [E.sub.m] = 1 + AB[[Phi].sub.f] / 1 - B[[Phi].sub.f] (7)
where: A = 2[L.sub.f] / [D.sub.f] and B = [E.sub.f]/[E.sub.m-1] / [E.sub.f]/[E.sub.m] + A (7)
and E, [E.sub.f], and [E.sub.m] are the elastic moduli of the composite, the reinforcing fiber phase and the matrix phase, respectively; [[Phi].sub.f] is the volume fraction and [L.sub.f]/[D.sub.f] is the aspect ratio of the fiber phase. The equation is based on the assumption of continuity of stress and strain along the fiber/matrix interface. The Halpin-Tsai Equation can be approximated by the simple rule of mixtures for [L.sub.f]/[D.sub.f] [greater than] 100 (26). The simple rule of mixtures is given by:
E = (1 - [[Phi].sub.f]) [E.sub.m] + [[Phi].sub.f][E.sub.f] (8)
In case of 11 mixing elements a high aspect ratio fiber/matrix morphology is already present at low DR. Therefore, we assume that the rule of mixtures is applicable here, but with a fiber modulus ([E.sub.f]), which is dependent on DR. In Table 4 the calculated fiber moduli of extrudates of DR 3 and 15 are listed. The calculated fiber moduli seem quite low compared to values in literature for Vectra A-type fibers in blends at similar DR. Based on the rule of mixtures Crevecoeur et al., (49) calculates a value of 65 GPA for Vectra A950 fibers dispersed in PS at a draw ratio of 17. Of course, one has to be careful in comparing these values, because of the differences in thermal and flow history [TABULAR DATA FOR TABLE 4 OMITTED] prior to the final drawing step. On the other hand, the largest increase in the mechanical properties of TLCPs or TLCP containing blends is typically achieved in the final drawing step. The reason why we find such low values for the TLCP fiber modulus is probably twofold. First, we have shown that the deformation of the TLCP fibers during drawing is less than arline (Table 1), probably caused by slippage effects between matrix and TLCP, as described by Kenig (50). This implies that the TLCP fibers in the blend experience a smaller effective DR than actually applied. Second, the melt-drawing process takes place at a temperature where the TLCP phase may be in a supercooled state. As pointed out by DiBenedetto et al. (51), at low draw temperatures, the LC domains in the TLCP phase may be too rigid to be deformed in the extensional stress field. Higher melt-drawing temperatures might thus lead to better composite properties (46). For the same reason, feeding the TLCP at higher temperatures to the static mixer, as for example Sukhadia et al. (13) did with the PET/Vectra and PP/Vectra system (Vectra fed to static mixer at 330 [degrees] C) might give better results. In case of our blend system, however, this was not possible because of possible thermal degradation of the matrix.
Another striking feature that can be deduced from Table 4 is that the fiber modulus is higher in the blends with a higher TLCP content. Crevecoeur et al. (49) and also others (52, 53) report an increase in molecular orientation (and thus a higher fiber modulus) with higher TLCP content. They attributed this to the difference in TLCP phase size in the blends. The larger dispersed TLCP particles, which are present at higher TLCP concentration, are easier to deform in the melt-drawing or spinning step. Indeed, in our case, the deformation of the TLCP phase is higher at higher TLCP content (see Table 1), which may thus account for the higher fiber moduli. In addition, the higher melt-drawing temperature at higher TLCP content (since more TLCP at 300 [degrees] C is injected in the matrix stream of 270 [degrees] C) may also lead to higher fiber moduli.
It has been shown that in-situ composites can be generated via a single distributive mixing process, where TLCP and matrix are fed separately by two extruders to a static mixer. Moreover, it is possible to blend TLCP/thermoplast systems, which normally do not have an overlap in the temperature-processing window. In-situ composites have been generated with this processing method, from blends consisting of a Vectra A900 TLCP and an Arnitel em630 thermoplastic elastomer, which could not be processed in one extruder.
By varying the number of mixing elements of the static mixer, strands with various morphologies can be obtained. Using 8 mixing elements leads to a strat-flied morphology of continuous Vectra layers in Arnitel, whereas 11 mixing elements results in the desired (continuous) fiber/matrix morphology. By using 14 mixing elements, a skin-core morphology is obtained, with fibers in the skin and droplets or low aspect ratio fibers situated in the core region of the extrudate.
All strands show an increase in tensile modulus/strength with higher extrudate draw ratio, as a result of the increased molecular orientation of the TLCP phase, which is induced in the final drawing step. In case of the strands made with 8 and 11 mixing elements, a more or less similar increase in tensile properties is achieved, whereas in the strands made using 14 elements, this increase is significantly less.
The total level of reinforcement, however, is lower than expected, which is thought to be due to the low temperature of the post melt-drawing process, where the TLCP is in fact in a supercooled state.
Above the TLCP melting point, the fiber/matrix morphology of this system is unstable, resulting in break-up into a droplet/matrix morphology through Rayleigh distortions and retraction. The time scale for break-up of TLCP fibers with a diameter of [approximately]1 [[micro]meter], is of the order of a few seconds. This short break-up time compared to the residence/solidification time in the final parts of the blending process, appears to be the cause of the skin-core morphology in the strands made using 14 mixing elements. In general, the instability of the fiber/matrix morphology in the molten state presents a serious limitation regarding use of TLCPs as reinforcing agents in blends, in particular when using processing techniques requiring long cooling/solidification times.
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|Author:||Machiels, A.G.C.; Denys, K.F.J.; Van Dam, J.; Posthuma de Boer, A.|
|Publication:||Polymer Engineering and Science|
|Date:||Oct 15, 1996|
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