Patent Application
20120276367
The present disclosure relates to the field of microporous membranes obtained by means of cast films precursors. More particularly, the disclosure relates to a method of controlling the morphology of cast films.
Among a wide range of resins, polypropylene (PP) is a well-known semicrystalline polymer and, in comparison with polyethylene (PE), PP has a higher melting point, lower density, higher chemical resistance, and better mechanical properties, which make it useful for many industrial applications.
The crystalline phase orientation in semicrystalline polymers such as polypropylene enhances many of their properties particularly mechanical, impact, barrier, and optical properties [1]. Obtaining an oriented structure in PP is of great interest to many processes such as film blowing, fiber spinning, film casting, and etc. In these processes the polymer melt is subjected to shear (in the die) and elongational (at the die exit) flow and crystallize during or subsequent to the imposition of the flow.
It is well-known that strain under flow strongly enhances the crystallization kinetics and allows the formation of a lamellar structure instead of the spherulitic one. The effect of flow on crystallization is called flow-induced crystallization (FIC) while the flow can be shear, extensional or both [2]. FIC molecular models show that flow induces orientation of polymer chains, resulting in enhancement of the nucleation rate [2-4]. Under flow, two major types of crystallization can occur, depending on the magnitude of the stress [1]: low stress results in twisted lamellae, while high stress produces a shish-kebab structure in which the lamellae grow radially on the shish without twisting [1].
Similar to shear flow, it has also been reported that extensional flow promotes fibrillar like structure oriented in the flow direction that serves as nucleation for radial growth of chain-folded lamellae perpendicular to the stress direction [5].
The effects of material parameters on shear induced crystallization process for PP have been investigated using in-situ small angle X-ray scattering (SAXS) and/or wide angle X-ray diffraction (WAXD) analyses [6-8]. Agarwal et al. [6] examined the influence of long chain branches on the stress induced crystallization. Adding a certain level of branches improved the orientation of the crystal blocks and the crystallization kinetics due to the longer relaxation time and the molecular structure. Somani et al. [7] followed the orientation development upon applying different shear rates. They found that, at a certain shear rate, only molecules with a chain length (molecular weight) above a critical value (critical orientation molecular weight, Mc) can form stable oriented row nuclei (shish). The shorter chains create lamellae over these nuclei sites. In another study, Somani et al. [8] compared the oriented microstructure under shear flow of isotactic polypropylene melts (PP-A and PP-B) with the same number average molecular weight but different molecular weight distribution (MWD). The amount of the high molecular weight species was larger in PP-B than in PP-A. Their results showed that the shish structure evolved much earlier for PP-B, which had more pronounced crystal orientation and faster crystallization kinetics. They concluded that even a small increase in the concentration of the high molecular weight chains led to a significant increase in the shish or nuclei site formation. In our recent study [9], addition of up to 10 wt % of a high molecular weight component to a low molecular weight one enhanced the formation of the row-nucleated structure probably due to an increase in the nucleating sites.
The crystallization behavior of semicrystalline polymers is significantly influenced by the process conditions. Under quiescent isothermal crystallization, the size of the spherulites, the degree of crystallinity, and the kinetics depend on temperature, while in quiescent non-isothermal conditions, both temperature and cooling rate are influencing factors [2].
Numerous studies have focused on the structure of PE and PP blown films using various materials under different processing conditions. However, as far as Applicants know no experimental study has been conducted on the cast film process with emphasis on the various parameters that can influence the morphology of the films.
Microporous membranes are commonly used in separation processes such as battery separators and medical applications to control the permeation rate of chemical components. Due to the wide range of chemical structures, optimum physical properties, and low cost of polymers and polymer blends, these materials are known as the best candidates for the fabrication of microporous membranes.
The two main techniques to develop polymeric membranes are: solution casting and extrusion followed by stretching. High cost and solvent contamination are the main drawbacks of the solution technique. Techniques to make porous membranes from polymers without using any solvent were developed in the seventies of the last century for some applications, but most of the information on these processes remains proprietary to the companies' and are not available to the scientific community. One of the techniques is based on the stretching of a polymer film containing a row-nucleated lamellar structure [29]. Then, three consecutive stages are carried out to obtain porous membranes: (1) creating a precursor film having a row-nucleated lamellar structure by mechanism of shear and elongation-induced crystallization, (2) annealing the precursor film at temperatures near the melting point of the resin to remove imperfections in the crystalline phase and to increase lamellae thickness, and (3) stretching at low and high temperatures to create and enlarge pores, respectively [29, 30]. In fact, in this process the material variables as well as the applied processing conditions are parameters that control the structure and the final properties of the fabricated microporous membranes [29]. The material variables include molecular weight, molecular weight distribution, and chain structure of the polymer. These factors mainly influence the row-nucleated structure in the precursor films at the first step of the formation of microporous membranes.
A few studies have investigated the fabrication of porous membranes by stretching of lamellar morphology using polypropylene [35-37]. Sadeghi et al. [35, 36] considered the influence of molecular weight on orientation of the row-nucleated lamellar structure. They found that molecular weight was the main material parameter that controlled the orientation of the crystalline phase. It was demonstrated that the resin with high molecular weight developed larger orientation and thicker lamellae than the resin with low molecular weight. Sadeghi et al. [37] realized that an initial orientation was required in order to obtain a lamellar structure. The crystalline orientation in the precursor film depended on the molecular weight of the resin and the type of process (i.e. cast film or film blowing). It was shown that the cast film process was more efficient than film blowing for producing precursor films with the appropriate crystalline orientation.
Although quite a few authors have investigated the formation of porous membranes from various resins, information is still lacking on the control of morphology and performances of membranes.
According to one aspect, there is provided a method for controlling the morphology of a cast film, the method comprising extruding a cast film by controlling a cooling rate of the cast film by applying on the film a gas at a gas cooling rate of at least about 0.4 cm3/s per kg/hr.
According to one aspect, there is provided a method for controlling the morphology of a cast film, the method comprising extruding a cast film by controlling a cooling rate of the cast film by applying on the film a gas at a gas cooling rate of at least about 0.4 cm3/s per kg/hr in accordance with the extrudate flow rate.
According to another aspect, there is provided a method for preparing a microporous membrane comprising preparing a cast film by controlling the morphology of the cast film as indicated in the method previously described, annealing the film, and stretching the film.
According to another aspect, there is provided a multilayer microporous membrane comprising at least two cast films prepared by controlling the morphology of the cast films as indicated in the method previously described.
According to another aspect, there is provided a method of preparing a microporous membrane comprising preparing a multilayer cast film, annealing the film, and stretching the film.
According to another aspect, there is provided a method of preparing a microporous membrane comprising preparing a multilayer cast film, annealing the film, and stretching the film, wherein the multilayer cast film comprises, in the following order, a first polypropylene layer, a polyethylene layer, and a second polypropylene layer.
According to another aspect, there is provided a method of preparing a microporous membrane comprising preparing a multilayer cast film, annealing the film, and stretching the film, wherein the multilayer cast film comprises, in the following order, a first linear polypropylene layer, a high density polyethylene layer, and a second linear polypropylene layer.
In the appended drawings which represent various examples of the present disclosure:
FIGS. 33A1, 33A2, 33B1, 33B2, 33C1 and 33C2 shows, for examples according to the present disclosure, SEM micrographs of the surface (top images) and cross-section (bottom images) of the microporous membranes. made with: PP28 (FIGS. 33A1 and 33A2), 5 wt % PP08 blend (FIGS. 33B1 and 33B2), and 10 wt % PP08 blend (FIGS. 33C1 and 33C2); DR=70, cold stretching of 35%, followed by hot stretching of 55;
The following embodiments are presented as non-limiting examples.
In the method previously mentioned, the gas used for cooling the film can be air. It can also be various other gases commercially vailable such as nitrogen, argon, helium etc.
For example, the cast film can be prepared by extruding the film at a draw ratio (DR) of at least 50, 55, 60, 65, 70, 75, or 80. For example, the draw ration can be about 50 to about 100 or about 60 to about 90.
For example, the film can have a thickness of about 20 μm to about 60 μm, about 30 μm to about 50 μm, or about 32 μm to about 45 μm.
According to one embodiment, the gas can be blown on the film by means of at least one air knife.
For example the cast film can be a monolayer film or a multilayer film (such as having from 2 to 10 layers, 2 to 7 layers, 2 to 5 layers, 2 to 4 layers, 2 layers or 3 layers).
For example, the gas cooling rate can be of at least 0.5, 0.6, 0.7, 0.8, 0.9, 1.0, 1.2, 1.5, 2.0, 3.0, 3.5, 4.0, 4.5, 5.0, 5.5, 6.0, 6.5, 7.0, 7.5, 8.0, 8.5, or 10 cm3/s per kg/hr in accordance with the extrudate flow rate. Alternatively, the gas cooling rate can be about 0.5 to about 9.0, about 0.6 to about 5.5, or 0.7 to about 4.5 cm3/s per kg/hr in accordance with the extrudate flow rate.
For example, the gas cooling rate can be at least proportional to square of the extrudate flow rate or it can be proportional to the reciprocal of extrudate film width.
According to one embodiment, the film can be extruded by means of a die and rolled up on a at least one cooling drum.
For example, the least one cooling drum can be at a temperature of about 20° C. to about 150° C., about 40° C. to about 140° C., about 50° C. to about 140° C., about 75° C. to about 140° C., about 80° C. to about 130° C., about 85° C. to about 115° C., about 90° C. to about 120° C., or about 100° C. to about 110° C.
For example, the can film comprise polypropylene, polyethylene or a mixture thereof.
For example, the can film comprise linear polypropylene, high density polyethylene, or a mixture thereof.
For example, the film can have a lamellar crystal structure. For example, the film can have a crystallinity of at least 40, 50, 60, 70, 80, or 90%.
When preparing a microporous membrane by using a cast film prepared according to the method as previously discussed, the film can be annealed at temperatures below the melting temperature. For example, the film can also be annealed at a temperature of about 100° C. to about 150° C., about 110° C. to about 140° C., or about 120° C. to about 140° C. For example, the film can be stretched at a first temperature and the film can be stretched at a second temperature. For example, the first temperature can be about 10° C. to about 50° C., about 15° C. to about 40° C., or 20° C. to about 30° C. For example, the second temperature can be about 90° C. to about 150° C., about 100° C. to about 140° C., or about 110° C. to about 130° C.
For example, the film can be stretched of about 20% to about 75% at the first temperature and the film can be stretched of about 40 to about 200% at the second temperature.
For example, the film can be stretched of about 30% to about 70% at the first temperature and the film can be stretched of about 50 to about 175% at the second temperature.
For example, the film can be stretched of about 30% to about 40% at the first temperature and the film can be stretched of about 50 to about 60% at the second temperature.
For example, the film can be stretched of about 50% to about 60% at the first temperature and the film can be stretched of about 70 to about 80% at the second temperature.
When preparing a multilayer microporous membrane comprising at least two cast films prepared by controlling the morphology of the cast films as described in the method previously mentioned the at least two cast films can be annealed and stretched. For example, the at least two cast films can be annealed at temperatures below the melting temperature of each film. For example, the at least two cast films can be annealed at about 100° C. to about 130° C., about 110° C. to about 130° C., or about 120° C. to about 130° C.
For example, the at least two films can be stretched at a first temperature and then the at least two films can be stretched at a second temperature. For example, the first temperature can be about 10° C. to about 50° C., about 15° C. to about 40° C., or about 20° C. to about 30° C. For example, the second temperature can be about 90° C. to about 130° C., about 100° C. to about 130° C., or about 110° C. to about 130° C.
For example, the at least two films can be stretched of about 20% to about 75% at the first temperature and the at least two films can be stretched of about 40 to about 200% at the second temperature.
For example, the at least two films can be stretched of about 30% to about 70% at the first temperature and the at least two films can be stretched of about 50 to about 175% at the second temperature.
For example, wherein the at least two films can be stretched of about 30% to about 40% at the first temperature and the at least two films can be stretched of about 50 to about 60% at the second temperature.
For example, wherein the at least two films can be stretched of about 50% to about 60% at the first temperature and the at least two films can be stretched of about 70 to about 80% at the second temperature.
For example, the multilayer membrane can comprise three films, the multilayer membranes comprising, in the following order, a first linear polypropylene layer, a high density polyethylene layer, and a second linear polypropylene layer.
1—Effect of Processing on the Crystalline Orientation, Morphology, and Mechanical Properties of Polypropylene Cast Films and Microporous Membrane Formation
A commercial linear polypropylene (PP5341) supplied by ExxonMobil Company was selected. It had a melt flow rate (MFR) value of 0.8 g/10 min (under ASTM conditions of 230° C. and 2.16 kg) Its molecular weight was estimated from the relationship between the zero-shear viscosity and the molecular weight [10] and found to be around 772 kg/mol. The resin showed a polydispersity index (PDI) of 2.7, as measured using a GPC (Viscotek model 350) at 140° C. and 1,2,4-Trichlorobenzene (TCB) as a solvent. Its melting point, Tm, and crystallization temperature, Tc, obtained from differential scanning calorimetry at a rate of 10° C./min, were 161° C. and 118° C., respectively.
The cast films were prepared using an industrial multilayer cast film unit from Davis Standard Company (Pawcatuck, Conn.) equipped with a 2.8 mm thick and 122 cm width slit die and two cooling drums. The extrusion was carried out at 220° C. and the distance between the die exit to the nip roll was 15 cm. The die temperature was set at 220° C. and draw ratios of 60, 75, and 90 were applied. An air knife with dimensions of 3 mm opening and 130 cm width was mounted close to the die to provide air to the film surface right at the exit of the die. The variables of interest were chill roll temperature, amount of air flow, and draw ratio. The films were produced under chill roll temperatures of 120, 110, 100, 80, 50, and 25° C. For all the cast roll temperatures, the air cooling rates used were 0, 1.2, 7.0, and 12 L/s. These air cooling conditions are noted as: no air flow rate (N-AFR), low air flow rate (L-AFR), medium air flow rate (M-AFR), and high air flow rate (H-AFR), respectively.
For membrane fabrication, the precursor films with a thickness, width and length of 35 μm, 46, and 64 mm, respectively, were used. The films were first annealed at 140° C. for 30 min and then cold and hot stretched at 25° C. and 120° C., respectively. Both annealing and stretching were performed using an Instron machine equipped with an environmental chamber. A drawing speed of 50 mm/min was applied during cold and hot stretching steps. The details for the fabrication of the microporous membranes can be found elsewhere [9].
Fourier transform infrared spectroscopy (FTIR): For FTIR measurements, a Nicolet Magna 860 FTIR instrument from Thermo Electron Corp. (DTGS detector, resolution 2 cm−1, accumulation of 128 scans) was used. The beam was polarized by means of a Spectra-Tech zinc selenide wire grid polarizer from Thermo Electron Corp. The measurement is based on the absorption of infrared light at certain frequencies corresponding to the vibration modes of atomic groups present within the molecule. In addition, if a specific vibration is attributed to a specific phase, the orientation within that phase can be determined [11]. If the films are oriented, the absorption of plane-polarized radiation by a vibration in two orthogonal directions, specifically parallel and perpendicular to a reference axis (MD), should be different. The ratio of these two absorption values is defined as the dichroic ratio, D [11]:
where is the absorption parallel A⊥ is the absorption perpendicular to a specific reference axis. The Herman orientation function of this vibration is obtained according to [11]:
For polypropylene, absorption at the wavenumber of 998 cm−1 is attributed to the crystalline phase (c-axis) while absorption at the wavenumber of 972 cm−1 is due to the contribution of both crystalline and amorphous phases. From the former absorption, the orientation of the crystalline phase, fc, can be determined while from the latter, the average orientation function, ƒav, is obtained. The orientation of the amorphous phase, fam, can be calculated according to:
ƒav=Xcfc+(1−Xc)fam (3)
where Xc is the degree of the crystallinity. Using FTIR, the global, crystalline and amorphous orientations can be determined.
X-Ray Diffraction
XRD measurement was carried out using a Bruker AXS X-ray goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper CuKα radiation (λ=1.542 A°) was selected using a graphite crystal monochromator. The sample to detector distance was fixed at 9.2 cm for wide angle diffraction and 28.2 cm for small angle X-ray scattering analysis. To get the maximum diffraction intensity several film layers were stacked together to obtain the total thickness of about 2 mm.
Wide angle X-ray diffraction (WAXD) is based on the diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the polymer crystalline phase. Using a pole figure accessory, the intensity of the diffracted radiation for a given hkl plane is measured as the sample is rotated through all possible spherical angles with respect to the beam. This gives the probability distribution of the orientation of the normal to hkl plane with respect to the directions of the sample.
The Herman orientation function of a crystalline axis is given by [12]:
where φ is the angle between the unit cell axes (a, b, and c) and reference axes. Details about the calculations can be found elsewhere [12].
The orientation factors from WAXD are mainly due to the crystalline part, therefore no information about the orientation of the amorphous phase can be obtained. Small angle X-ray scattering (SAXS) was used to compare the level of the lamellae formation for the different samples and to estimate the long period between lamellae.
Thermal Analysis
Thermal properties of specimens were analyzed using a TA instrument differential scanning calorimeter (DSC) Q 1000. The thermal behavior of films was obtained by heating from 50 to 220° C. at a heating rate of 10° C./min. The reported crystallinity results were obtained using a heat of fusion of 209 J/g for fully crystalline polypropylene (PP) [13].
Mechanical and Tear Analysis
Tensile tests were performed using an Instron 5500R machine equipped with an environmental chamber for tests at high temperature. The procedure used was based on the D638-02a ASTM standard. A standard test method for the tear resistance of plastic films based on ASTM D1922 was used to obtain the MD and TD tear resistances. According to this standard, the work required in tearing is measured by the loss of energy of the encoder, which measures the angular position of the pendulum during the tearing operation.
Morphology
To clearly observe the crystal arrangement of the PP cast films, an etching method was employed to remove the amorphous part. The PP films were dissolved in a 0.7% solution of potassium permanganate in a mixture of 35 volume percentage of orthophosphoric and 65 volume percentage of sulfuric acid. The potassium permanganate was slowly added to the sulfuric acid under rapid agitation. At the end of the reaction time, the samples were washed as described in Olley and Bassett [14].
A field emission scanning electron microscope (FESEM—Hitachi S4700) was employed for the observation of the etched films surfaces as well as microporous membranes. This microscope provides high resolution of 2.5 nm at a low accelerating voltage of 1 kV and high resolution of 1.5 nm at 15 kV with magnification from 20× to 500 kx.
Water Vapor Transmission
The permeability to water vapor was measured via a MOCON PERMATRAN-W Model 101K at room temperature. It is composed of three chambers: an upper chamber containing liquid water and separated from the center chamber by two porous films. Water vapor diffuses from the first film to fill the space between the films to reach 100% relative humidity (RH). The center chamber is separated from the lower one by the test film. The diffused vapor is swept away by N2 gas to a relative humidity (RH) sensor.
Rheological Characterization
Dynamic rheological measurements were carried out using a Rheometric Scientific SR5000 stress controlled rheometer with a parallel plate geometry of 25 mm diameter and a gap equal to 1.5 mm at the temperatures of 180, 195, 210, and 225° C. under nitrogen atmosphere. Molded discs of 2 mm thick and 25 mm in diameter were prepared using a hydraulic press at 190° C. Prior to frequency sweep tests, time sweep tests at a frequency of 0.628 rad/s and different temperatures were performed for two hours to check the thermal stability of the specimens. No degradation (less than 3% changes) was observed at test temperatures for the duration of the frequency sweep measurements. The dynamic data were obtained in the linear regime and used to evaluate the weighted relaxation spectra of the samples.
Experimental data that clearly demonstrate the effects of process conditions are firstly shown. More particularly, air cooling and drum temperature, on crystallization, orientation of the amorphous and crystalline phases, and also tear and mechanical properties. Subsequently, two morphological pictograms are proposed to describe the observed experimental data and the reasons for these observations are discussed. Finally, the structure and properties of the microporous membranes obtained from the PP cast films having different microstructures are presented.
The effects of cast roll temperature (Tcast) and air cooling on thermal behavior of the films were examined using differential scanning calorimetry (DSC) and the results are shown in
The orientation and arrangement of the crystal lamellae in the cast films are influencing factors in controlling the final properties of manufactured films.
The effect of the amount of applied strain, either in shear or elongational flow, on the lamellar structure of the various resins has been investigated recently [7, 15-17]. The authors reported that as the level of strain increased, more lamellae accompanied with better orientation were generated. The influence of draw ratio on the orientation of the crystalline and amorphous phases has also been considered [9, 17, 18]. In the cast film process of PP, an almost linear relationship between the draw ratio and the orientation factor was reported [9, 18]. At low draw ratios, the lamellae were not well aligned perpendicular to the flow direction, but at high draw ratios, the lamellae aligned themselves perpendicular to the machine direction, resulting in higher orientation. In this study, the influence of the draw ratio on the orientation function with and without the use of air flow is illustrated in
The effect of air cooling on orientation of the crystalline phase was also considered using WAXD, as shown in
The crystalline orientation can also be analyzed quantitatively from the pole figures of the 110 and 040 planes, as illustrated in
The orientation features, in terms of cos2 (φ) of the crystalline axes (i.e. a, b, and c (see the sketch in
The degree of crystallinity (Xc) of the samples determined using WAXD and DSC is presented in Table 1. In WAXD, the contributions arising from the crystalline and amorphous parts were extracted via peak fitting of the 28 diffraction pattern. Similarly to DSC results, it was observed that cooling improves crystallinity. However, the crystallinity obtained from WAXD was slightly higher than that from DSC. In addition, the average crystal width in the directions of the 110 and 040 crystallographic planes were determined from the full width at half maximum Δ(2θ) of the deconvoluted diffraction profiles according to the following equation [20]:
where K is a crystallite form coefficient that is taken equal to 1 and λ is the X-ray wavelength. Although it is known that this equation is not accurate because it neglects the broadening due to the lattice distortions, it allows a useful comparison of the crystalline structures for the various films. Table 1 also presents the variations of the D110 and D040 with air cooling for the films casted at 120° C. Both D110 and D040 are enhanced by the use of low air cooling and do not vary by further increases of air flow. The D040 crystallite size corresponds to an average size of the crystallites that are oriented parallel to the film plane. Therefore, the increase of D040 suggests that the crystallite size is increased in a direction parallel to the b crystallographic axis. The effect of the cast roll temperature on the D040 was also evaluated (not shown here) and no noticeable impacts were found.
trates the SAXS patterns as well as the azimuthal intensity profile for the films obtained under different air cooling rates. The equatorial streak in the SAXS patterns is attributed to the formation of the shish, while the meridian maxima are attributed to the lateral lamellae or kebabs [6]. Looking at the meridian intensity (either pattern or azimuthal profile), the formation of more lamellae for the air cooled samples is obvious. In addition, for all conditions, it is clear that the contribution of the shish to the crystalline phase is much less than that of lamellae, confirming the results of Somani et al. [21] for PE and PP.
The long period distance, Lp, was estimated from the position of the Lorentz corrected intensity maxima, as demonstrated in
A first order peak arising from stacks of parallel lamellae and a second order peak indicating that the periodicity of the lamellae is high [22] is observed. Air cooling slightly shifts the peaks to higher values, indicating a decrease of the long period spacing. Long period spacing results for the no air cooled as well as the air cooled specimens are also reported in
The previous results imply that an oriented shish-kebab crystal structure is obtained by the use of air cooling in addition to the chill roll, compared to a much less ordered crystal structure for the no air cooled films. These differences can be clearly visualized from SEM surface images of the etched films (etching removes the amorphous region), as demonstrated in
It is well established that the structure of the crystalline and amorphous phases strongly influence the mechanical and tear properties of films. In other words, the mechanical and tear behaviors are closely related to structure changes. Zhang et al. [24] studied the microstructure of LLDPE, LDPE, and HDPE blown films and showed that the type of oriented structure was greatly dependent on the type of polyethylene as well as on the processing conditions. These structure differences were shown to translate into different ratios for MD and TD tear and tensile strengths [24].
The enhancement of the mechanical response along MD, due to the high orientation of the cooled specimens along MD, was, however, accompanied by significant reductions of the elongation at break along TD (
Table 2 reports the mechanical properties along MD and TD for the films produced at Tcast of 120, 110, 100° C. under N-AFR as well as L-AFR conditions. Compared to the no air cooled films, a significant effect of low air blowing on the mechanical properties, at all drum temperatures, is observed. It is clear that the Young modulus, yield stress, tensile toughness, and tensile strength along MD decrease as Tcast decreases. This is due to the formation of more lamellae at higher Tcast for the no air subjected films (see
It is well understood that tear measurements are very sensitive to the type and alignment of crystalline morphology [24]. Along MD, tear resistance values of 0.178, 0.154, 0.146, and 0.121 g/μm were measured for the films obtained at Tcast=120° C. and N-AFR, L-AFR, M-AFR, and H-AFR, respectively: the higher the orientation of the crystalline and amorphous phases, the lower the tear resistance along MD. It was observed that measurements of the tear resistance along TD for samples subjected to air cooling were impossible, because the tearing direction deviated most of the time to MD. In fact, there is a high resistance in TD when compared to MD, which causes a crack in MD and created non-reproducible data that are not reported here. This implies that for air cooled films, a shish-kebab lamellar crystal structure with shishs aligned in MD is formed.
Based on observations from thermal analysis, FTIR results, WAXD and SAXS patterns, microscopy, mechanical and tear properties, two microstructural pictograms (one for the no air cooled cast film and another for the cast films produced with the use of air cooling) are proposed as depicted in
For the films produced at high chill roll temperatures and without air cooling, FTIR data, WAXD and SAXS patterns suggest the presence of a lamellar crystalline structure (rows of lamellae and/or cross-hatched), which is not preferentially oriented in MD (see
In contrast, for the films produced under air cooling, FTIR data, SAXS and WAXD patterns suggest the presence of a stacked lamellar crystal structure, which is preferentially oriented into MD (see
Based on FTIR results for the orientation of the amorphous phase (see
In the following section, justifications regarding the roles of air cooling and drum temperature on the final crystal microstructures are presented. The orientation and morphological differences are believed to originate from the rheological characteristics and crystallization kinetics. It is well known that temperature has an effect on the relaxation time of polymer chains as well as on the crystallization rate. In order to consider the effects of temperature on the applied stress and relaxation time of the molecules, linear dynamic rheological measurements were carried out.
In general, the rate of crystallization is first controlled by nucleation and then by the growth and packing of the crystals [28]. In our case, the air cooling causes a large decrease in the extruded film temperature such that crystallization temperature of the resin is reached before the frost line is formed. This increases the number of nuclei sites resulting in a much faster crystallization rate. This fact, together with the intrinsic temperature effects on the relaxation time, as discussed above, determines a significant coupling between temperature and flow, yielding to a novel highly oriented lamellar structure. In other words, the use of air cooling in addition to chill rolls in the cast film process helps flow induced crystallization to occur at lower temperatures. This will noticeably increase the number of shish or nuclei sites, and consequently the crystallization kinetics is promoted resulting in a well oriented shish-kebab structure.
To produce microporous membranes by the stretching technique, precursor films with an adequate orientation and alignment of the crystal lamellae are needed [9, 18]. In this study, the effects of microstructure differences of the PP cast films on the microporous membranes morphology and water vapor transmission rate were investigated. Three consecutive stages were carried out to obtain porous membranes: cast or precursor film formation, annealing, and stretching in two steps (cold and hot). During cold stretching, the pores were created whereas in the subsequent hot stretching they were enlarged. WAXD and FTIR measurements clearly showed that cooling drastically enhanced orientation of crystal lamellae in the precursor films; hence, a microporous membrane with more pore density and better tortuosity is expected as air cooling is utilized.
In view of the above, it can thus be said that:
Two commercial linear polypropylenes (PP28, PP08) were selected. Both PPs were supplied by ExxonMobil Company and had MFR values of 2.8 g/10 min (under ASTM conditions of 230° C. and 2.16 kg) and 0.8 g/10 min, respectively. The main characteristics of the resins are shown in Table 3. The molecular weight of the linear PPs was evaluated using the relation between the zero-shear viscosity and the molecular weight [39]. The melting point, Tm, and the crystallization temperature, Tc, of the resins were obtained using differential scanning calorimetry. For the rheological characterization, blends containing 2, 5, 10, 30, 50, and 70 wt % PP08 were prepared using a twin screw extruder (Leistritz Model ZSE 18HP co-rotating twin screw extruder) followed by water cooling and pelletizing. The temperature profile along the barrel was set at 160/180/190/200/200/200/200° C. The extrusion was carried out at 80 rpm. During blending, 3000 ppm of a stabilizer, Irganox B225, was added to avoid thermal degradation of the polymers. To make sure that all samples have the same thermal and mechanical history, unblended components were extruded under the same conditions.
Dynamic rheological measurements were carried out using a Rheometric Scientific SR5000 stress controlled rheometer with a parallel plate geometry of 25 mm diameter and a gap equal to 1.5 mm at the temperature of 190° C. under nitrogen atmosphere. Molded discs of 2 mm thick and 0.25 mm in diameter were prepared using a hydraulic press at 190° C. Time sweep test was first performed at a frequency of 0.628 rad/s for two hours. Material functions such as complex viscosity, elastic modulus, and weighted relaxation spectrum in the linear viscoelastic regime were determined in the frequency range from 0.01 to 500 rad/s. In order to obtain more accurate data, the frequency sweep test was carried out in four sequences while the amount of applied stress in each sequence was determined by a stress sweep test.
The precursor films from PP28 and blends containing 2, 5, 10, and 20 wt % PP08 were prepared by extrusion using a slit die of 1.9 mm thick and 200 mm width. An air knife was mounted on the die to supply air to the film surface right at the exit of the die. The main parameters were die temperature, cooling rate, and draw ratio (ratio of the roll speed to the die exit velocity) [7]. In this study the die temperature was set at 220° C. and the maximum speed of the fan was applied, thus the only variable was the draw ratio. The film samples were prepared under draw ratios of 70, 80, and 90.
For membrane preparation, precursor films having a thickness, width and length of 35 μm, 46 mm and 64 mm, respectively, were used. Both annealing and stretching were performed using an Instron machine equipped with an environmental chamber. A drawing speed of 50 mm/min was applied during the cold and hot stretching steps.
Fourier transform infrared spectroscopy (FTIR): For FTIR measurements, a Nicolet Magna 860 FTIR instrument from Thermo Electron Corp. (DTGS detector, resolution 4 cm−1, accumulation of 128 scans) was used. The measurement is based on the absorption of infrared light at certain frequencies corresponding to the vibration modes of atomic groups present within the molecule. In addition, if a specific vibration is attributed to a specific phase, the orientation within that phase can be determined [36]. If the films are oriented, the absorption of plane-polarized radiation by a vibration in two orthogonal directions, specifically parallel and perpendicular to a reference axis (MD), should be unequal. The ratio of these two absorption values is defined as
the dichroic ratio, D [40]:
where is the absorption parallel and A⊥ is the absorption perpendicular to a specific reference axis. The Herman orientation function of this vibration is obtained according to [40]:
For polypropylenes, absorption at the wavelength of 998 cm−1 is attributed to the crystalline phase (c-axis) while absorption at the wavelength of 972 cm−1 is due to the contribution of both crystalline and amorphous phases. From the former absorption, the orientation of the crystalline phase, fc, can be determined while from the latter, the average orientation function, ƒav, is obtained. The orientation of the amorphous phase, fam, can be calculated according to:
f
av
=X
c
f
c+(1−Xc)fam (8)
where Xc is the degree of the crystallinity. Using FTIR, the global, crystalline and amorphous orientations can be determined.
X-Ray Diffraction
XRD measurement was carried out using a Bruker AXS X-ray goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper Cu Kα radiation (λ=1.542 A°) was selected using a graphite crystal monochromator. The sample to detector distance was fixed at 9.2 cm for wide angle diffraction and 28.2 cm for small angle x-ray scattering analysis. To get the maximum diffraction intensity several film layers were stacked together to obtain the total thickness of about 2 mm.
Wide angle X-ray diffraction (WAXD) is based on the diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the polymer crystalline phase. Using a pole figure accessory, the intensity of the diffracted radiation for a given hkl plane is measured as the sample is rotated through all possible spherical angles with respect to the beam. This gives the probability distribution of the orientation of the normal to hkl plane with respect to the directions of the sample.
The Herman orientation function is given by [35]:
where φ is the angle between the unit cell axes (a, b, and c) and reference axes. Details about the calculations can be found elsewhere [35].
The orientation factors from WAXD are mainly due to the crystalline part, therefore no information about the orientation of the amorphous phase can be obtained. Small angle X-ray scattering (SAXS) was utilized to estimate the long period distance between the lamellae.
Thermal Analysis
Thermal properties of specimens were analyzed using a TA instrument differential scanning calorimeter (DSC) Q 1000. Samples were heated from 50 to 220° C. at a heating rate of 10° C./min.
BET Measurement
To obtain the surface area and pore diameter of the membranes, a flowsorb Quantachrome instrument BET ASI-MP-9 was used. A nitrogen and helium gas mixture was continuously fed through the sample cell, which was kept at liquid nitrogen temperature. At different pressures, the total volume of nitrogen gas adsorbed on the surface was measured. The volume of gas needed to create an adsorbed monomolecular layer was calculated as follows [41]:
where P is the experimental pressure, P° is the saturation pressure, v is the volume of the adsorbate, vm is the volume of gas required to form an absorbed monomolecular layer, and c is a constant. The procedure for estimating the surface area from Eq. 5 can be found elsewhere [42].
Mercury Porosimetry
The average pore size, pore size distribution, and porosity of the membranes were also evaluated using a mercury porosimeter (PoreMaster PM33). After evacuation of the cell, it is filled by mercury and then pressure is applied to force mercury into the porous sample. The amount of intruded mercury is related to the pore size and porosity.
Water Vapor Transmission
The permeability to water vapor was measured via a MOCON PERMATRAN-W Model 101 K at room temperature. It is composed of three chambers: an upper chamber containing the liquid water and separated from the center chamber by two porous films. Water vapor diffuses from the first film to fill the space between the films to reach 100% relative humidity (RH). The center chamber is separated from the lower one by the test film. The diffused vapor is swept away by N2 gas to a RH sensor.
Mechanical Analysis
Tensile tests were performed using an Instron 5500R machine equipped with a chamber for running tests at high temperature. The procedure used was based on the D638-02a ASTM standard.
Puncture Resistance
Puncture tests were performed using a 10 N load cell of the Instron machine used for the tensile tests. A needle with 0.5 mm radius was used to pierce the samples. The film was held tight in the camping device with a central hole of 11.3 mm. The displacement of the film was recorded against the force (Newton) and the maximum force was reported as the puncture strength.
The complex shear viscosities as a function of frequency for the neat PPs as well as for the blends are shown in
log η*(ω)=φβ log(η*(ω))1+(1−φβ)log(η*(ω))2 (11)
where φβ is the PP08 content. Adding a high molecular weight component (PP08) causes monotonic increases of the complex shear viscosity, which is due to the presence of the larger molecules of PP08. Good agreement with the logarithmic mixing rule can be observed for all samples, suggesting miscibility of both PP components.
In order to quantitatively analyze the role of adding large chains on the melt relaxation of the blends, the weighted relaxation spectra evaluated from dynamic moduli (G′, G″, ω) using the NLREG (non-linear regularization) software [44] are plotted in
The main mechanism of shear and/or elongation-induced crystallization is the propagation of the lamellae based on the fibrils or nuclei sites [35, 36]. As fibrils are mostly created from the long chains [32-34] and long chains have larger relaxation time (
The relaxation behavior can also be shown in Cole-Cole plots, which are plot of η″ versus η′ as illustrated in
Draw ratios of 70, 80, and 90 were applied to the extruded films to investigate the role of the extension ratio on the orientation of the precursor films, as illustrated in
To determine the optimum annealing conditions that will lead to the largest amount of crystallinity, annealing at 140° C. without extension, at 140° C. under 5% extension, and at 120° C. without extension was carried out and the measured crystallinity values are plotted in
The effects of annealing and stretching on crystallinity were examined using differential scanning calorimetry (DSC) and the results are shown in
SAXS measurements were performed to examine the role of annealing and stretching on the lamellae spacing. The long period distance, Lp, was estimated from the position of the intensity maxima, as demonstrated in
It has been shown that annealing significantly influences the tensile response of films [35]. According to Sadeghi et al. [35], due to the planar morphology of the annealed sample, the rupture along the MD for the annealed film occurs at much smaller strain than for the non-annealed film. Puncture tests were performed to investigate the effects of annealing on the mechanical properties of the samples along the ND, and the results are presented in
The effect of blending on the mechanical properties of precursor films along MD and TD is illustrated in
To control the final structure of the produced membranes, obtaining a precursor film with the adequate orientation and alignment of the crystal lamellae is needed. The WAXD measurements were performed to consider the influence of blending on the level of orientation, as illustrated in
The neat PP28 membrane contains thicker lamellae in the precursor films and has lower orientation yielding in a difficult lamellae separation and less interconnectivity (
Properties for membranes made from blends and their neat component are reported in Table 4. The membranes containing 5 wt % PP08 and 10 wt % PP08 exhibit pore densities about twice and four times that of the PP28 membrane, respectively. The PP08 microporous membrane shows pore density much lower than the 10 wt % PP08 membrane. The table also compares the results on specific surface area and average pore size of the membranes determined by BET and mercury porosimetry. The pore diameters obtained from BET and mercury porosimetry are almost identical. The specific surface area varies from 5.9 to 26.2 m2/g depending on the PP08 content. As the 10 wt % PP08 blend microporous membranes has smaller diameter pores the larger value for the specific surface area is due to its larger pore density. An average pore diameter of 0.12 μm was determined for the PP28, 5 wt % PP28 and 10 wt % PP28 membranes. It should be noted that the neat PP08 microporous membrane shows a much lower surface area but larger pores compared to the 10 wt % PP08 microporous membrane and the reason for these is discussed in the following paragraph.
Table 4 also presents the water vapor transmission rates for the obtained microporous membranes. The permeability increased by a factor of 3 when 10 wt % PP08 was added to PP28. The addition of a high molecular weight component enhances the permeability, which is attributed to more pores, higher porosity, and better interconnection between the pores for the blend samples containing up to 10 wt % PP08. No significant increase of permeability was observed by further addition of PP08 except for the neat PP08 microporous membrane. In the blends, adding of more than 10 wt % PP08 possibly destroys the lamellar morphology, resulting in no changes or even lower permeability. Microporous membranes made of the neat PP08 showed a fibrillar structure with smaller number of lamellae (not presented here) than the 10 wt % PP08 microporous blend. This was due to the presence of larger number of long chains in PP08. Although the pore density of the PP08 porous membrane is much smaller than for the 10 wt % PP08 membrane, its pores are much larger leading to a better pore interconnection and larger WVTR. Although the permeability of the neat. PP08 membrane was larger than that of all blend membranes, the objective of this work as mentioned earlier is to control the performances of microporous membranes using polymer blends.
Also as shown by Table 4, the Young modulus of the membranes slightly increases as the amount of the high Mw PP increases (Table 4). This can be explained by the better orientation, of the lamellae for the blend films compared to the neat PP28.
In the preparation of porous membranes using the stretching technique, voids are nucleated by cold stretching and enlarged by subsequent hot stretching [1,2]. According to Johnson [2], the micro void morphologies produced via this method are a consequence of inter lamellar separation, which takes place at temperatures above Tg of the specific semi crystalline polymer. This is in contrast to amorphous polymers which reportedly form voids (i.e. crazes, this process is termed crazing) upon deformation at temperatures below their respective Tg. Sadeghi et al. [8] found that the pores size of the cold stretched films obtained from the PP resins with distinct Mw did not significantly vary. However, a difference in the lamellae thickness was observed. To find the optimum cold stretching conditions, cold drawing was carried out at 25° C. and 45° C. and under predetermined amounts of extension while the amount of hot stretching was kept constant.
Similar experiments were performed to investigate the influence of hot stretching. In contrast to cold stretching, no maximum was observed when the films were stretched to different levels (
In this work, the structure and performances of microporous membranes made from blends of linear low and high molecular weight PPs have been investigated. In addition, the influence of annealing conditions on the crystallinity and stretching variables on the water vapor transmission rate (WVTR) were examined. Our findings can be summarized as follows:
Using the puncture test, it was shown that addition of high molecular weight species does not have dramatic influence on the mechanical properties of precursor films along ND. However, tensile tests revealed a slight reduction of the mechanical properties along MD and TD.
3—Microporous Membranes Obtained from PP/HDPE Multilayer Films by Stretching
Commercially available lithium battery separators are made from polyolefins such as polypropylene (PP) and polyethylene (PE). These materials are compatible with the cell chemistry and can be used for many cycles without significant degradation in properties [54]. Lithium (Li) batteries will generate heat if accidentally overcharged. Separator shutdown is a useful safety feature for restricting thermal reactions in Li-batteries [54,55]. Shutdown occurs close to the melting temperature of the polymer, leading to pores collapse and restricting passage of current through the cell. PP separators melt around 160° C. whereas PE separators have shutdown temperature between 120 and 130° C. If in a battery, the heat dissipation is slow, even after shutdown, the cell temperature may continue to increase before starting to cool [54]. Recently, manufacturers have started producing trilayer separators where a porous PE layer is sandwiched between two porous PP layers. In such a case, the PE layer has lower shutdown temperature while PP provides the mechanical stability at and above the shutdown temperature [54].
Three commercially available processes are used for making microporous membranes: solution casting (also known as extraction process), particle stretching, and dry-stretching [56]. In the extraction process, the polymeric raw material is mixed with a processing oil or plasticizer, this mixture is extruded and the plasticizer is removed through an extraction process [57]. In the particle stretch process, the polymeric material is mixed with particles, this mixture is extruded, and pores are formed during stretching at the interface of the polymer and solid particles [58]. Costly processes and difficulties in dealing with solvent and particle contaminations are main drawbacks of such methods. However, the dry-stretch process is based on the stretching of a polymer film containing a row-nucleated lamellar structure [59]. Three consecutive stages are carried out to obtain porous membranes by this technique: 1) creating a precursor film having a row-nucleated lamellar structure through shear and elongation-induced crystallization of the polymer having proper molecular weight and molecular weight distribution, 2) annealing the precursor film at temperatures near the melting point of the resin to remove imperfections in the crystalline phase and to increase lamellae thickness, and 3) stretching at low and high temperatures to create and enlarge pores, respectively [59,60].
In fact, in this process, the material variables as well as the applied processing conditions are parameters that control the structure and the final properties of the fabricated microporous membranes [59]. The material variables include molecular weight, molecular weight distribution, and chain structure of the polymer. These factors mainly influence the row-nucleated structure in the precursor films in the first step of the formation of microporous membranes. According to Sadeghi et al. [61,62], molecular weight was the main material parameter that controlled the orientation of the row-nucleated lamellar structure. The resins with high molecular weight developed larger orientation and thicker lamellae than those with low molecular weight. In our recent study [63], the addition of up to 10 wt % of a high molecular weight component to a low molecular weight resin enhanced the formation of the row-nucleated structure due to an increase in the nucleating sites. In Sadeghi et al. [64], a superior permeability was obtained by adding a small amount of a long-chain branched polypropylene (LCB-PP) to a linear polypropylene (L-PP). The effects of process conditions such as draw ratio (DR), air flow rate (AFR), and cast roll temperature on the structure of PP cast films and microporous membranes have been investigated by the Applicants [65]. A significant enhancement in orientation was observed by applying air cooling and increasing DR. An ordered stacked lamellar structure was seen for the films subjected to low air cooling whereas the films produced without air cooling showed a spherulitic structure.
There are two main industrial processes for the production of films: film blowing and cast film extrusion. It is well known that the thickness variation in blown films are considerably greater compared to cast films. For the preparation of porous membranes, obtaining a precursor film with good thickness uniformity is strongly recommended since any non uniformity causes irregularities in the stress distribution in the following stretching process. In addition, compared to film blowing, cast film process has more flexibility in the supply of air cooling from both sides, leading to a more uniform lamellar structure in both surfaces.
A commercial linear polypropylene (PP) and a commercial high density polyethylene (HDPE) were selected. PP5341E1 was supplied by ExxonMobil and had a melt flow rate (MFR) value of 0.8 g/10 min (under ASTM D1238 conditions of 230° C. and 2.16 kg). HDPE 19A was provided by NOVA Chemicals and had an MFR value of 0.72 g/10 min (under ASTM D1238 conditions of 190° C. and 2.16 kg). The main characteristics of the resins are shown in Table 5. The molecular weight (Mw) and polydispersity index (PDI) of the HDPE was supplied by company and that of the PP was measured using a GPC (Viscotek model 350) with 1,2,4-Trichlorobenzene (TCB) as a solvent at a column temperature of 140° C. The melting point, Tm, and crystallization temperature, Tc, of the resins obtained from differential scanning calorimetry at a rate of 10° C./min are also reported in Table 5.
Dynamic rheological measurements were carried out using a Rheometric Scientific SR5000 stress controlled rheometer with a parallel plate geometry of 25 mm diameter and a gap equal to 1.5 mm at the temperature of 190° C. under nitrogen atmosphere. Molded discs of 2 mm thick and 25 mm in diameter were prepared using a hydraulic press at 190° C. Prior to frequency sweep tests, time sweep tests at a frequency of 0.628 rad/s and 190° C. were performed for two hours to check the thermal stability of the specimens. No degradation (less than 3% changes) was observed for the duration of the frequency sweep measurements. Complex viscosity and weighted relaxation spectrum in the linear viscoelastic regime were determined in the frequency range from 0.01 to 500 rad/s. In order to obtain more accurate data, the frequency sweep test was carried out in four sequences while the amount of applied stress in each sequence was determined by a stress sweep test.
The cast films were prepared using an industrial multilayer cast film unit from Davis Standard Company (Pawcatuck, Conn.) equipped with a 2.8 mm opening and 122 cm width slit die and two cooling drums. The extrusion was carried out at 220° C. and the distance between the die exit to the nip roll was 15 cm. The die temperature was set at 220° C. and draw ratios of 60, 75, and 90 were applied. An air knife with dimensions of 3 mm opening and 130 cm width was mounted close to the die to provide air to the film surface right at the exit of the die. The variables of interest were draw ratio and amount of air flow. The films were produced under chill roll temperature of 50° C. The air cooling rates used were 1.2 and 12 L/s. These air cooling conditions are noted as: low air flow rate (L-AFR) and high air flow rate (H-AFR), respectively.
For the membrane fabrication, the precursor films with thickness, width, and length of 32 μm, 46 mm, and 64 mm, respectively, were used. The films were first annealed at 120° C. for 30 min and then cold and hot stretched at 25° C. and 120° C., respectively. Both annealing and stretching were performed in an Instron tensile machine equipped with an environmental chamber. Drawing speeds of 500 mm/min and 25 mm/min were applied during cold and hot stretching steps, respectively.
Thermal Analysis
Thermal properties of specimens were analyzed using a TA instrument differential scanning calorimeter (DSC) Q 1000. The thermal behavior of films was obtained by heating from 50 to 220° C. at a heating rate of 10° C./min. The reported crystallinity results were obtained Using a heat of fusion of 209 and 280 J/g for fully crystalline PP and HDPE, respectively [66,67].
Fourier Transform Infrared Spectroscopy (FTIR)
For FTIR measurements, a Nicolet Magna 860 FTIR instrument from Thermo Electron Corp. (DIGS detector, resolution 2 cm−1, accumulation of 128 scans) was used. The beam was polarized by means of a Spectra-Tech zinc selenide wire grid polarizer from Thermo Electron Corp. The measurement is based on the absorption of infrared light at certain frequencies corresponding to the vibration modes of atomic groups present within the molecule. In addition, if a specific vibration is attributed to a specific phase, the orientation within that phase can be determined [61]. If the films are oriented, the absorption of plane-polarized radiation by a vibration in two orthogonal directions, specifically parallel and perpendicular to a reference axis (MD), should be different. The ratio of these two absorption values is defined as the dichroic ratio, D [61]:
where is the absorption parallel and A⊥ is the absorption perpendicular to a specific reference axis. The Herman orientation function of this vibration is obtained according to [61]:
For polypropylene, absorption at the wavenumber of 998 cm−1 is attributed to the crystalline phase (c-axis) while that at 972 cm−1 is due to the contribution of both crystalline and amorphous phases. From the former absorption, the orientation of the crystalline phase, Fc, can be determined while from the latter, the average orientation function, Favg, is obtained. The orientation of the amorphous phase, F8, can then be calculated according to:
F
avg
=X
c
F
c+(1−Xc)Fa (14)
where Xc is the degree of the crystallinity.
For polyethylene, absorption at the wavenumber of 730 cm−1 is attributed to the a-axis of the unit crystal cell while absorption at the wavenumber of 720 cm−1 is due to the b-axis. The similarity of the normal (N) and transverse (T) spectra confirmed that the orientation is mostly uniaxial [68]. In such a case, it is not necessary to use the tilted film technique. The orientation function of the a- and b-axes could be obtained from Eq. 13 while that of the c-axis orientation is calculated according to the orthogonality equation:
F
a
+F
b
+F
c=0 (15)
X-Ray Diffraction
XRD measurement was carried out using a Bruker AXS X-ray goniometer equipped with a Hi-STAR two-dimensional area detector. The generator was set up at 40 kV and 40 mA and the copper CuKα radiation (λ=1.542 A°) was selected using a graphite crystal monochromator. The sample to detector distance was fixed at 9.2 cm for wide angle diffraction and 28.2 cm for small angle X-ray scattering analysis. To get the maximum diffraction intensity several film layers were stacked together to obtain the total thickness of about 2 mm.
Wide angle X-ray diffraction (WAXD) is based on the diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the polymer crystalline phase. Using a pole figure accessory, the intensity of the diffracted radiation for a given hkl plane is measured as the sample is rotated through all possible spherical angles with respect to the beam. This gives the probability distribution of the orientation of the normal to hkl plane with respect to the directions of the sample.
The Herman orientation function Fij of a crystalline axis i with respect to a reference axis j is given by [69]:
where φij is the angle between the unit cell axes i (a, b, or c) and the reference axis j.
The Herman orientation functions were derived from the 110 and 040 pole figures for the PP and the 110 and 200 pole figures for the HDPE. Details about the calculations for PP can be found in Sadeghi et al. [61]. For the HDPE, since the a-axis of the unit cells is perpendicular to the 200 plane, its orientation relative to the machine direction can be measured directly as follow:
On the other hand, Fc (orientation of the c-axis) with respect to MD is determined by the combination of data of two planes for the HDPE, which are 110 and 200 [69]:
cos2φc=1−1.435 cos2φ110−0.565 cos2φ200 (18)
The orientation parameter for the b-axis can be calculated from the orthogonality relation:
cos2φb=1−cos2φa−cos2φc (19)
The orientation factors from WAXD are mainly due to the crystalline part, therefore no information about the orientation of the amorphous phase can be obtained. Small angle X-ray scattering (SAXS) was used to compare the level of the lamellae formation for the different samples and to estimate the long period between lamellae.
Mechanical Analysis and Puncture Resistance
Tensile tests were performed using an Instron 5500R machine equipped with an environmental chamber for tests at high temperature. The procedure used was based on the D638-02a ASTM standard. Puncture tests were performed using a 10 N load cell of the Instron machine used for the tensile tests. A needle with 0.5 mm radius was used to pierce the samples. The film was held tight in the camping device with a central hole of 11.3 mm. The displacement of the film was recorded against the force and the maximum force was reported as the puncture strength. Strain rates of 50 mm/min and 25 mm/min were utilized during tensile and puncture tests, respectively.
Morphology
To clearly observe the crystal arrangement of the precursor films, an etching method was employed to remove the amorphous part. The films were dissolved in a 0.7% solution of potassium permanganate in a mixture of 35 vol % of orthophosphoric and 65 vol % of sulphuric acid. The potassium permanganate was slowly added to the sulphuric acid under rapid agitation. At the end of the reaction time, the samples were washed as described in Olley and Bassett [70].
A field emission scanning electron microscope (FE-SEM—Hitachi S4700) was employed for the observation of the etched precursor films and microporous membranes surfaces as well as cross-sections. This microscope provides high resolution of 2.5 nm at a low accelerating voltage of 1 kV and high resolution of 1.5 nm at 15 kV with magnification from 20× to 500 kx.
Water Vapor Transmission
The permeability to water vapor was measured via a MOCON PERMATRAN-W Model 101 K at room temperature. It is composed of three chambers: an upper chamber containing liquid water and separated from the center chamber by two porous films. Water vapor diffuses from the first film to fill the space between the films to reach 100% relative humidity (RH). The center chamber is separated from the lower one by the test film. The diffused vapor is swept away by N2 gas to a relative humidity (RH) sensor.
BET Measurement
To obtain the surface area of the membranes, a Micromeritics, BET Tristar 3000 was used. A nitrogen and helium gas mixture was continuously fed through the sample cell, which was kept at liquid nitrogen temperature. At different pressures, the total volume of nitrogen gas adsorbed on the surface was measured. The volume of gas needed to create an adsorbed monomolecular layer was calculated as follows [71]:
where P is the experimental pressure, P° is the saturation pressure, V is the volume of the adsorbate, Vm is the volume of gas required to form an absorbed monomolecular layer, and c is a constant. The procedure for estimating the surface area from Eq. 20 can be found elsewhere [72].
Rheological and Film Characterization
The complex shear viscosities as a function of frequency for the resins are shown in
To produce microporous membranes by the stretching technique, precursor films with an adequate orientation and alignment of the crystal lamellae are needed [62,63]. The higher the crystalline alignment in the precursor, the better is expected the lamellae separation and, as a consequence, the larger the porosity and permeability of the microporous membranes. In this study, the roles of draw ratio (DR), cooling air flow rate (AFR), and annealing on the crystalline alignments of single layer films as well as the components in a multilayer film are probed using WAXD and FTIR.
From literature, in PE, two major types of crystallization can occur depending on the magnitude of stress in flow [74]: low stress produces kebabs in the form of twisted ribbons resulting in off-axis 110 and meridian 200 diffractions. In contrast, high stress produces flat kebabs (planar crystal structure) leading to the appearance of equatorial 110 and 200 diffractions. When the magnitude of flow is in-between, an intermediate arrangement is formed, resulting in off-axis 200 and 110 diffractions [74]. However, PP under flow usually generates planar lamellar morphology with less dependence on the flow magnitude [65].
The WAXD patterns as well as the diffraction intensity profiles for the PP and HDPE shown in
The orientation features, in terms of cos2 (φ) of the crystalline axes (i.e. a, b, and c) along MD, TD, and ND obtained from the Herman orientation function for the PP and HDPE single layers as well as the components in the multilayer film are presented in
Differences in the crystal structure and arrangement of the precursor films can be clearly visualized from SEM surface images of the etched films (etching removes the amorphous region), as demonstrated in
It is well established that the structure of the crystalline phase strongly influences the mechanical properties of films. Zhang et al. [79] studied the microstructure of LLDPE, LDPE, and HDPE blown films and showed that the type of oriented structure was greatly dependent on the type of polyethylene as well as on the processing conditions. In our previous study [65], significant increases in the Young modulus, yield stress, tensile strength, tensile toughness along MD and a drastic decrease in the elongation at break along TD were observed for polypropylene cast films subjected to air cooling. Table 6 reports the results on the mechanical properties of the films along MD and TD for DR=60 and 90. All the properties improve along. MD and elongation at break along TD reduces with increasing DR, due to better crystal alignment. Additionally, it should be noticed that the mechanical properties of the trilayer films are between those of the monolayer films.
In general, blends of PP and HDPE are known to be immiscible systems. The interfacial morphology for the etched multilayer film is illustrated in
The produced precursor films should be annealed at a proper temperature before to be cold and hot stretched. As annealing is performed at a temperature that is above the onset of mobility in the crystalline structure (Tα), it is postulated that during annealing, the lamellae twist and orient perpendicular to the machine direction. Also, melting of small lamellae and their recrystallization with better orientation can occur [63]. Our previous study [63] showed that annealing at 140° C. without extension was the optimum annealing condition for PP. However, because Tm of the HDPE is around 129° C., the annealing temperature of the trilayer film should be lower than the HDPE melting point and above the alpha transition temperature, Tα, of the PP (Tα,PP=110° C. obtained from the dynamic mechanical thermal analysis). Therefore, we selected 120° C. for annealing of the trilayer film. In order to be able to compare the results, the single layer precursor films were annealed at the same temperature.
Table 7 presents the mechanical properties along MD and TD as well as the puncture resistance along ND for the microporous membranes. Obviously, the porous membranes have nearby similar tensile responses in MD and, as expected, the tensile properties along TD are considerably smaller than in MD. However, the PP microporous membranes show a much lower strain at break along TD than that the HDPE and multilayer membranes. In addition, due to the presence of the elongated interlamellar microfibrils in the microporous membranes, a pronounced increase in the tensile strength and a drastic decrease in elongation at break are observed for the membranes compared to the precursor films (see Tables 6 and 7). The decrease in the modulus for the membranes compared to the precursor films is possibly due to the lower interconnection between the lamellae in the membranes as a result of the tie chains pulled out during the pores formation. Also as shown by Table 7, the maximum piercing force is considerably larger for the PP membrane than the HDPE membrane, which can be explained by the smaller pores and lower porosity for the former. Therefore, it could be concluded that in the PP/HDPE/PP membranes, the side layers (i.e. PP) improve the puncture resistance drastically.
In the preparation of porous membranes using the stretching technique, voids are formed by cold stretching and enlarged by subsequent hot stretching [59,60]. According to Johnson [60], the micro void morphologies produced via this method are a consequence of inter lamellar separation, which takes place at temperatures above Tg of the specific semicrystalline polymers. Sadeghi et al. [62] found that the pores size of the cold stretched films obtained from the PP resins with distinct Mw did not vary significantly. However, a difference in the lamellae thickness was observed. Our previous study [63] showed that the water vapor transmission rate (WVTR) of the cold stretched PP films increased as the stretch ratio increased up to 30% while further stretching resulted in a reduction in WVTR. To find the optimum cold stretching extension for the present PP as well as the HDPE, cold stretching was carried out under predetermined levels of extension while the amount of hot stretching was kept constant.
To clearly understand the reasons for the opposite extension dependence of the PP and HDPE in the first stretching step, we show in
Similar experiments (data not shown) were performed to investigate the influence of hot stretching level. In contrast to the cold stretching behavior of the PP, no maximum was observed when the cold drawn films were stretched to different hot stretch levels. The pores created in cold stretching are enlarged during the hot stretching step and consequently enhance WVTR. More flexibility of the lamellae at high temperatures can be a reason for the increase of pore size with increasing extension ratio.
Finally, the trilayer microporous membranes obtained at 55% cold extension followed by 75% hot extension showed WVTR values about 30% lower than the monolayer PP and HDPE obtained under the same conditions. This could be due to the presence of the interface and lower orientation of the PP and HDPE components in the multilayer film than the monolayer ones (see
In this work, the structure and performances of microporous membranes made from monolayer and trilayer films of the PP and HDPE have been investigated. Applicants findings can be summarized as follows:
The present disclosure has been described with regard to specific examples. The description was intended to help the understanding of the disclosure, rather than to limit its scope. It will be apparent to one skilled in the art that various modifications may be made to the disclosure without departing from the scope of the disclosure as described herein, and such modifications are intended to be covered by the present document.
| Filing Document | Filing Date | Country | Kind | 371c Date |
|---|---|---|---|---|
| PCT/CA2010/000952 | 6/18/2010 | WO | 00 | 3/19/2012 |
| Number | Date | Country | |
|---|---|---|---|
| 61219581 | Jun 2009 | US | |
| 61288042 | Dec 2009 | US |
