Classifications

• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C55/00—Shaping by stretching, e.g. drawing through a die; Apparatus therefor
  • B29C55/005—Shaping by stretching, e.g. drawing through a die; Apparatus therefor characterised by the choice of materials
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D67/00—Processes specially adapted for manufacturing semi-permeable membranes for separation processes or apparatus
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D67/00—Processes specially adapted for manufacturing semi-permeable membranes for separation processes or apparatus
  • B01D67/0002—Organic membrane manufacture
  • B01D67/002—Organic membrane manufacture from melts
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D67/00—Processes specially adapted for manufacturing semi-permeable membranes for separation processes or apparatus
  • B01D67/0002—Organic membrane manufacture
  • B01D67/0023—Organic membrane manufacture by inducing porosity into non porous precursor membranes
  • B01D67/0025—Organic membrane manufacture by inducing porosity into non porous precursor membranes by mechanical treatment, e.g. pore-stretching
  • B01D67/0027—Organic membrane manufacture by inducing porosity into non porous precursor membranes by mechanical treatment, e.g. pore-stretching by stretching
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D69/00—Semi-permeable membranes for separation processes or apparatus characterised by their form, structure or properties; Manufacturing processes specially adapted therefor
  • B01D69/12—Composite membranes; Ultra-thin membranes
  • B01D69/1212—Coextruded layers
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D69/00—Semi-permeable membranes for separation processes or apparatus characterised by their form, structure or properties; Manufacturing processes specially adapted therefor
  • B01D69/12—Composite membranes; Ultra-thin membranes
  • B01D69/1216—Three or more layers
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D71/00—Semi-permeable membranes for separation processes or apparatus characterised by the material; Manufacturing processes specially adapted therefor
  • B01D71/06—Organic material
  • B01D71/26—Polyalkenes
  • B01D71/261—Polyethylene
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D71/00—Semi-permeable membranes for separation processes or apparatus characterised by the material; Manufacturing processes specially adapted therefor
  • B01D71/06—Organic material
  • B01D71/26—Polyalkenes
  • B01D71/262—Polypropylene
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C48/00—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor
  • B29C48/03—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor characterised by the shape of the extruded material at extrusion
  • B29C48/07—Flat, e.g. panels
  • B29C48/08—Flat, e.g. panels flexible, e.g. films
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C48/00—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor
  • B29C48/16—Articles comprising two or more components, e.g. co-extruded layers
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C48/00—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor
  • B29C48/25—Component parts, details or accessories; Auxiliary operations
  • B29C48/88—Thermal treatment of the stream of extruded material, e.g. cooling
  • B29C48/911—Cooling
  • B29C48/9135—Cooling of flat articles, e.g. using specially adapted supporting means
  • B29C48/915—Cooling of flat articles, e.g. using specially adapted supporting means with means for improving the adhesion to the supporting means
  • B29C48/917—Cooling of flat articles, e.g. using specially adapted supporting means with means for improving the adhesion to the supporting means by applying pressurised gas to the surface of the flat article
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29D—PRODUCING PARTICULAR ARTICLES FROM PLASTICS OR FROM SUBSTANCES IN A PLASTIC STATE
  • B29D7/00—Producing flat articles, e.g. films or sheets
  • B29D7/01—Films or sheets
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D2323/00—Details relating to membrane preparation
  • B01D2323/08—Specific temperatures applied
  • B01D2323/081—Heating
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B01—PHYSICAL OR CHEMICAL PROCESSES OR APPARATUS IN GENERAL
  • B01D—SEPARATION
  • B01D2323/00—Details relating to membrane preparation
  • B01D2323/08—Specific temperatures applied
  • B01D2323/082—Cooling
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C48/00—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor
  • B29C48/03—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor characterised by the shape of the extruded material at extrusion
  • B29C48/09—Articles with cross-sections having partially or fully enclosed cavities, e.g. pipes or channels
  • B29C48/10—Articles with cross-sections having partially or fully enclosed cavities, e.g. pipes or channels flexible, e.g. blown foils
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C48/00—Extrusion moulding, i.e. expressing the moulding material through a die or nozzle which imparts the desired form; Apparatus therefor
  • B29C48/25—Component parts, details or accessories; Auxiliary operations
  • B29C48/88—Thermal treatment of the stream of extruded material, e.g. cooling
  • B29C48/911—Cooling
  • B29C48/9135—Cooling of flat articles, e.g. using specially adapted supporting means
  • B29C48/914—Cooling drums
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29C—SHAPING OR JOINING OF PLASTICS; SHAPING OF MATERIAL IN A PLASTIC STATE, NOT OTHERWISE PROVIDED FOR; AFTER-TREATMENT OF THE SHAPED PRODUCTS, e.g. REPAIRING
  • B29C55/00—Shaping by stretching, e.g. drawing through a die; Apparatus therefor
  • B29C55/02—Shaping by stretching, e.g. drawing through a die; Apparatus therefor of plates or sheets
  • B29C55/023—Shaping by stretching, e.g. drawing through a die; Apparatus therefor of plates or sheets using multilayered plates or sheets
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29K—INDEXING SCHEME ASSOCIATED WITH SUBCLASSES B29B, B29C OR B29D, RELATING TO MOULDING MATERIALS OR TO MATERIALS FOR MOULDS, REINFORCEMENTS, FILLERS OR PREFORMED PARTS, e.g. INSERTS
  • B29K2023/00—Use of polyalkenes or derivatives thereof as moulding material
  • B29K2023/04—Polymers of ethylene
  • B29K2023/06—PE, i.e. polyethylene
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29K—INDEXING SCHEME ASSOCIATED WITH SUBCLASSES B29B, B29C OR B29D, RELATING TO MOULDING MATERIALS OR TO MATERIALS FOR MOULDS, REINFORCEMENTS, FILLERS OR PREFORMED PARTS, e.g. INSERTS
  • B29K2023/00—Use of polyalkenes or derivatives thereof as moulding material
  • B29K2023/10—Polymers of propylene
  • B29K2023/12—PP, i.e. polypropylene
• • B—PERFORMING OPERATIONS; TRANSPORTING
  • B29—WORKING OF PLASTICS; WORKING OF SUBSTANCES IN A PLASTIC STATE IN GENERAL
  • B29K—INDEXING SCHEME ASSOCIATED WITH SUBCLASSES B29B, B29C OR B29D, RELATING TO MOULDING MATERIALS OR TO MATERIALS FOR MOULDS, REINFORCEMENTS, FILLERS OR PREFORMED PARTS, e.g. INSERTS
  • B29K2105/00—Condition, form or state of moulded material or of the material to be shaped
  • B29K2105/04—Condition, form or state of moulded material or of the material to be shaped cellular or porous
• • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
  • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
  • Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
  • Y10T428/00—Stock material or miscellaneous articles
  • Y10T428/249921—Web or sheet containing structurally defined element or component
  • Y10T428/249953—Composite having voids in a component [e.g., porous, cellular, etc.]
  • Y10T428/249978—Voids specified as micro

Definitions

• 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.
• polypropylene 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 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.
• 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].
• 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.
• 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].
• 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 method for controlling the morphology of a cast film 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 cm 3 /s per kg/hr.
• a method for controlling the morphology of a cast film 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 cm 3 /s per kg/hr in accordance with the extrudate flow rate.
• 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.
• 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.
• a method of preparing a microporous membrane comprising preparing a multilayer cast film, annealing the film, and stretching the film.
• 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.
• 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.
• Figures 14A and 14B show, for examples according to the present disclosure, proposed pictograms of the molecular structure for: no air cooled cast films (Figure14A) and air cooled cast films (Figure 14B) (the solid lines represent the tear path along MD and the dash lines show the tear path along TD);
• Figure 15 shows, for examples according to the present disclosure, weighted relaxation spectra for different melt temperatures (the vertical dash lines represent the range of frequencies covered during the experiments);
• WVTR water vapor transmission rate
• Figure 21 shows, for examples according to the present disclosure
• Figure 22 shows, for examples according to the present disclosure, crystalline orientation function (obtained from FTIR) as a function of draw ratio for precursor films;
• Figure 27 shows, for examples according to the present disclosure
• Figure 38 shows, for examples according to the present disclosure
• Figures 4OA, 4OB and 40C show, for examples according to the present disclosure, normalized 2D WAXD patterns and pole figures for films obtained under different DR, AFR, and annealing: PP monolayer ( Figure 40A), PP in multilayer ( Figure 40B), and HDPE monolayer ( Figure 40C), wherein annealing was performed at 120 0 C for 30 min;
• Figures 41 A, 41 B and 41 C show, for examples according to the present disclosure, orientation characteristics as cos 2 ( ⁇ ) of the crystal axes (a, b and c) along MD, TD, and ND for the films obtained under different DR, AFR, and annealing: c-axis ( Figure 41A), a-axis ( Figure 41B), and />axis ( Figure 41C), wherein annealing was performed at 120 0 C for 30 min;
• Figure 42 shows, for examples according to the present disclosure
• Figure 43 shows, for examples according to the present disclosure
• Figure 45 shows, for examples according to the present disclosure
• Figure 46 shows, for examples according to the present disclosure
• Figure 50 shows, for examples according to the present disclosure
• Figure 51 is a schematic representation of an apparatus used for carrying out an example of a method according to the present disclosure, wherein there is shown the distance between the die exit to the nip roll, and wherein delta X represents (Td-Tc) the difference of temperature between the extruder and the cast roll (cooling drum) and the Tcast, wherein Ua and Ta represents the gas cooling rate and the temperature of the gas;
• 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.
• 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.
• DR draw ratio
• the draw ration can be about 50 to about 100 or about 60 to about 90.
• the film can have a thickness of about 20 ⁇ m to about
• the gas can be blown on the film by means of at least one air knife.
• 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).
• the gas cooling rate can be of at least 0.5, 0.6, 0.7,
• 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 cm 3 /s per kg/hr in accordance with the extrudate flow rate.
• 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.
• the film can be extruded by means of a die and rolled up on a at least one cooling drum.
• the least one cooling drum can be at a temperature of about 20° C to about 150 0 C, about 40° C to about 140 0 C, about 50° C to about 14O 0 C, about 75° C to about 140 0 C, about 80° C to about 13O 0 C, about 85° C to about 115 ° C, about 90° C to about 12O 0 C, or about 100° C to about 110 ° C.
• the can film comprise polypropylene, polyethylene or a mixture thereof.
• the can film comprise linear polypropylene, high density polyethylene, or a mixture thereof.
• the film can have a lamellar crystal structure.
• the film can have a crystallinity of at least 40, 50, 60, 70, 80, or 90 %.
• the film can be annealed at temperatures below the melting temperature.
• the film can also be annealed at a temperature of about 100° C to about 15O 0 C, about 110° C to about 140 0 C, or about 120° C to about 14O 0 C.
• the film can be stretched at a first temperature and the film can be stretched at a second temperature.
• 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.
• 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.
• 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.
• 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.
• 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.
• 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.
• the at least two cast films can be annealed and stretched.
• the at least two cast films can be annealed at temperatures below the melting temperature of each film.
• the at least two cast films can be annealed at about 100° C to about 130 0 C, about 110° C to about 130 0 C, or about 120° C to about 13O 0 C.
• 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.
• 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.
• the second temperature can be about 90° C to about 13O 0 C, about 100° C to about 130 0 C, or about 110° C to about 13O 0 C.
• 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.
• 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.
• 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.
• 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.
• 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.
• ExxonMobil Company was selected. It had a melt flow rate (MFR) value of 0.8 g/10min (under ASTM conditions of 230 0 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 0 C and 1 ,2,4-Trichlorobenzene (TCB) as a solvent. Its melting point, T m , and crystallization temperature, T 0 , obtained from differential scanning calorimetry at a rate of 10 °C/min, were 161 0 C and 118 0 C, respectively.
• the cast films were prepared using an industrial multilayer cast film unit from Davis Standard Company (Pawcatuck, CT) equipped with a 2.8 mm thick and 122 cm width slit die and two cooling drums.
• the extrusion was carried out at 220 0 C and the distance between the die exit to the nip roll was 15 cm.
• the die temperature was set at 220 0 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 0 C.
• 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.
• the precursor films with a thickness, width and length of 35 ⁇ m, 46, and 64 mm, respectively, were used.
• the films were first annealed at 140 0 C for 30 min and then cold and hot stretched at 25 0 C and 120 0 C, respectively. Both annealing and stretching were performed using an lnstron 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].
• FTIR Fourier transform infrared spectroscopy
• 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]:
• Wide angle X-ray diffraction is based on the diffraction of a monochromatic X-ray beam by the crystal lographic planes (hkl) of the polymer crystalline phase.
• hkl crystal lographic planes
• 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.
• ⁇ 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.
• SAXS Small angle X-ray scattering
• 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 0 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].
• 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 2Ox to 500kx.
• 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 N 2 gas to a relative humidity (/?/-/) sensor.
• the effect of air cooling on orientation of the crystalline phase was also considered using WAXD, as shown in Figure 5.
• the first and second rings represent the patterns for the 110 and 040 crystalline planes, respectively [12].
• a diffraction ring is seen for the 110 crystallographic plane for the no air cooled film, indicating a low crystalline phase orientation.
• arcs that are sharper and more concentrated in the center are observed for the air cooled samples, implying more orientation.
• This behavior can be better shown when the intensity is plotted as a function of the azimuthal angle.
• the crystalline orientation can also be analyzed quantitatively from the pole figures of the 110 and 040 planes, as illustrated in Figure 6.
• the normal to the 110 plane is the bisector of the a and b axes and the 040 plane is along the b-axis of unit crystal cells [12].
• the film obtained without air cooling slight orientations of the 110 and 040 planes are detected in MD and TD, respectively.
• a significant orientation of the 110 plane is observed along TD and that of the 040 plane (b-axis) is in both TD and ND.
• the pole figures for the samples obtained at higher air flow rates i.e. M-AFR and H-AFR
• the schematics in Figure 6 represent the crystal alignments based on their pole figures.
• orientation functions obtained using FTIR were slightly larger than the values from the WAXD pole figure.
• the discrepancies in the values of the measured c-axis orientation may be due to different factors such as peak deconvolution, contribution of the amorphous phase etc., as discussed for PE and PP elsewhere [1 ,19].
• 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 Figure 9.
• the rectangle in Figure 10a illustrates the interface of a spherulite and rows of lamellae and its higher magnification image is shown on the right. More spherulites and lamellar branching are observed for the films prepared at r cast of 110 0 C and without the use of air cooling (Figure 10b).
• the lamellar branching from the primary lamellae is produced by epitaxial growth due to cross-hatched lamellar texture, which is a unique characteristic of PP crystalline structure [1].
• the rectangle in Figure 10b exhibits the impingement of spherulites, clearly demonstrated by its higher magnification micrograph on the right.
• the stress-strain response for the no air cooled samples along MD exhibits the typical behavior of films of a spherulitic structure with an elastic response at low deformation, yielding and plastic behavior at medium deformation, and strain hardening at high elongation.
• the stress-strain response of the air cooled specimens along MD reflects the typical behavior of films of a lamellar crystalline morphology with an initial elastic response at low deformation followed by two strain hardening zones. An extensive discussion on this behavior can be found in Samuels [25].
• Table 2 reports the mechanical properties along MD and TD for the films produced at 7 cast of 120, 110, 100 0 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 r cast decreases. This is due to the formation of more lamellae at higher 7 cast for the no air subjected films (see Figures 10a and 10b) and annealing effect at high r cast for the air subjected films.
• T Ca31 IOO 0 C L-AFR 8538( 198) 221(10) 978( 18) 24(00) 469( 08) 235(49) 005 (001)
• tear measurements are very sensitive to the type and alignment of crystalline morphology [24].
• the melt temperature right at the exit of the die decreases and, as a consequence, the applied stress (or relaxation time) rises drastically. This promotes the number of shishs or nuclei sites, resulting in a noticeable increase of the probability for the formation of lamellae by the low molecular weight chains.
• the rate of crystallization is first controlled by nucleation and then by the growth and packing of the crystals [28].
• 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.
• 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.
• Figure 16 presents SEM micrographs of the surface of the fabricated membranes. Very thick lamellae, non uniform pores and a small amount of pores for the porous membrane obtained from the no air cooled film is observed ( Figure 16a). However, for the membrane prepared from the cast film subjected to a small air cooling rate, the number of pores noticeably increase and more uniform pore sizes and a better morphology are observed as well ( Figure 16b). Note more and thinner lamellae appear for the latter compared to the former, supporting the previous results.
• FIG 17 presents the water vapor transmission rates (WVTR) of the produced microporous membranes.
• WVTR water vapor transmission rates
• Small WVTR were recorded for the no air cooled samples produced at different cast roll temperatures.
• the WVTR increased by a factor of 20 when the film surface was subjected to a low air flow, which is attributed to the formation of more pores, higher porosity, and better interconnection between the pores.
• further increases of the air cooling rate i.e. at M-AFR and H-AFR
• did not significantly increase the permeability was (see the inset of Figure 17), indicating that more air cooling does not dramatically influence the lamellar structure, in accordance with the previous results (see Figures 2 to 8).
• Microporous membranes having high pore density, large porosity, and high water vapor permeability were obtained by lamellae separation for cast films prepared using air cooling.
• 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 0 C.
• the extrusion was carried out at 80 rpm.
• 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.
• 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 0 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.
• precursor films having a thickness, width and length of 35 ⁇ m, 46 mm and 64 mm, respectively, were used. Both annealing and stretching were performed using an lnstron machine equipped with an environmental chamber. A drawing speed of 50 mm/min was applied during the cold and hot stretching steps.
• FTIR Fourier transform infrared spectroscopy
• Wide angle X-ray diffraction is based on the diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the polymer crystalline phase.
• 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.
• ⁇ is the angle between the unit cell axes (a, b, and c) and reference axes.
• SAXS Small angle X-ray scattering
• 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]:
• 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 ⁇ fe gas to a RH sensor.
• Puncture resistance Puncture tests were performed using a 10 N load cell of the lnstron 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 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 ( Figure 20), therefore, adding a high molecular weight component favors the preparation of precursor films with an adequate level of crystalline lamellae.
• the relaxation behavior can also be shown in Cole-Cole plots, which are plot of ⁇ " versus ⁇ 1 as illustrated in Figure 21.
• the semicircular shape of the Cole-Cole for the blends is another evidence of miscibility [45, 46].
• Figure 24 presents the Herman orientation function of the crystalline phase as well as of the amorphous phase for the precursor and annealed films for all the blends. It is obvious that adding up to 10 wt% PP08 enhances the orientation of the both crystalline and amorphous phases. In addition, in comparison with the non-annealed films, significant improvement of the orientation is observed for the annealed specimens in the entire range of compositions. As annealing is performed at a temperature that is close to the onset of mobility in the crystalline structure (T 0 ), it is postulated that during annealing, the lamellae twist and orient perpendicular to the machine direction.
• the arcs are sharper and more concentrated in the center for annealed samples, suggesting more orientation.
• the orientation features as cos 2 ( ⁇ ) of the crystalline axes (i.e. a, b, and c (see Figure 29)) along MD, TD, and ND obtained from the Herman orientation function for the precursor, annealed, and stretched films are plotted in the triangular graph of Figure 26d. It is obvious that annealing causes movement of the c-axis of the crystals towards MD, while the a- and b-axes take a position closer to the TD-ND plane. This obviously shows that annealing improves the orientation of the films, in accordance with the FTIR data. Clearly stretching does not dramatically influence the orientation of the unit cell.
• SAXS can detect the distance between the lamellae, and during stretching the lamellae are separated and pores are created, so dramatic influence of stretching is observed in the SAXS patterns.
• Figure 28 shows the 2D SAXS patterns for the PP28 and 10 wt% PP08 blend films.
• the equatorial streak in the SAXS patterns is attributed to the formation of the shish, while meridian maxima are attributed to the lamellae or kebabs [3]. Looking at the meridian intensity, the formation of more lamellae for the blend containing 10 wt% PP08 is evident.
• Figure 30 shows that blending reduces the elongation at break for the precursor films along the MD, except for the 5 wt% PP08 sample that shows a peak.
• the reason for the peak in the elongation at break for the 5 wt% PP08 is unclear at present.
• the following decrease in the elongation at break for the precursor films containing more PP08 is possibly due to the more orientation of the amorphous and crystalline phases for these samples (see Figure 24).
• the tie chains between lamellae are more oriented and more fibrils are expected as the level of PP08 increases. It is well known that better orientation along MD results in less deformation at break. For the transverse direction, a reduction of the maximum stress and amount of elongation at break are also observed (Figure 31). This can be explained by the larger number of fibrils and smaller lamellae for the blend films.
• the WAXD measurements were performed to consider the influence of blending on the level of orientation, as illustrated in Figure 32.
• the first and second rings of the pole figures show the patterns for the 110 and 040 crystalline planes, respectively [51].
• the normal to the 110 plane is the bisector of the a and b axes and 040 is along the b-axis of the unit crystal cells [38]. More intense arcs in the meridian and equatorial zones are apparent for the blend sample, indicating greater orientation for the crystal lamellae [51].
• Figure 33 presents SEM micrographs of the surface and cross- section of the fabricated membranes for PP28 and blends containing 5 wt% PP08 and 10 wt% PP08.
• Figure 33a1 exhibits very thick lamellae and a small amount of pores for PP28 (i.e. low molecular weight component).
• PP08 i.e. high molecular weight component
• Figures 33b1 and 33c1 This behavior can be explained as follows. As the content of the high molecular weight component increases, the number of fibrils or nuclei sites increases resulting in smaller lamellae and more pores.
• the neat PP28 membrane contains thicker lamellae in the precursor films and has lower orientation yielding in a difficult lamellae separation and less interconnectivity (Figure 33a2). More pores and thinner lamellae for the 10 wt% PP08 blend membrane lead to a better interconnectivity of the pores ( Figure 33c2).
• the membrane tortuosity which is defined as the length of the average pore to the membrane thickness [52] seems to be less for the 10 wt% PP08 blend membrane than for the PP28 membrane.
• SEM micrographs of the surface of the cold and hot stretched films were compared (SEM micrographs of the cold stretched films are not shown), the number of pores was found to be larger for the hot stretched samples. This confirms the findings of Sadeghi et al. [36] and it was explained by the melting of some lamellae and their recrystallization in the form of interconnected bridges during the hot stretching step.
• 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.
• 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.
• 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.
• 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.
• Figure 34 reports the pore size distribution for the microporous neat PP and membranes containing 5 wt% and 10 wt% PP08. It is obvious that blending does not dramatically influence the peak positions in the pore size distribution curves and all the three samples reveal peaks around 0.11 ⁇ m. However, adding PP08 significantly enhances the area under curves, indicating porosity increases. Porosity values of 30, 35, and 44 % were evaluated for the PP28, 5 wt% blend, and 10 wt% blend membranes, respectively. The lower porosity of the neat PP membrane is attributed to its thicker lamellae and, consequently, more resistant to lamellae separation.
• 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 0 C whereas PE separators have shutdown temperature between 120 and 130 0 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].
• PP polypropylene
• microporous membranes Three commercially available processes are used for making microporous membranes: solution casting (also known as extraction process), particle stretching, and dry-stretching [56].
• solution casting also known as extraction process
• particle stretching 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].
• 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.
• 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].
• 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.
• a commercial linear polypropylene (PP) and a commercial high density polyethylene (HDPE) were selected.
• PP5341 E1 was supplied by ExxonMobil and had a melt flow rate (MFR) value of 0.8 g/10min (under ASTM D1238 conditions of 230 0 C and 2.16 kg).
• HDPE 19A was provided by NOVA Chemicals and had an MFR value of 0.72 g/10min (under ASTM D1238 conditions of 190 0 C and 2.16 kg).
• MFR melt flow rate
• the molecular weight (M w ) 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 0 C.
• T m melting point
• T c crystallization temperature
• the cast films were prepared using an industrial multilayer cast film unit from Davis Standard Company (Pawcatuck, CT) equipped with a 2.8 mm opening and 122 cm width slit die and two cooling drums.
• the extrusion was carried out at 220 0 C and the distance between the die exit to the nip roll was 15 cm.
• the die temperature was set at 220 0 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 0 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.
• the films were first annealed at 120 0 C for 30 min and then cold and hot stretched at 25 0 C and 120 0 C, respectively. Both annealing and stretching were performed in an lnstron 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.
• the thermal behavior of films was obtained by heating from 50 to 220 0 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
• FTIR Fourier transform infrared spectroscopy
• 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]:
• Wide angle X-ray diffraction is based on the diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the polymer crystalline phase.
• 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 F,y of a crystalline axis / with respect to a reference axisy is given by [69]: where ⁇ ,y is the angle between the unit cell axes / (a, b, or c) and the reference axis/ [00193]
• the Herman orientation functions were derived from the 110 and
• F c orientation of the c-axis
• MD is determined by the combination of data of two planes for the HDPE, which are 110 and 200 [69]:
• the orientation parameter for the b-axis can be calculated from the orthogonality relation:
• SAXS Small angle X-ray scattering
• 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 2Ox to 500kx.
• 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 N 2 gas to a relative humidity (RH) sensor.
• RH relative humidity
• 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]:
• the inset in Figure 37 compares the weighted relaxation spectra of the resins evaluated from dynamic moduli (G', G", ⁇ ) using the NLREG (non linear regularization) software [73] (the vertical dash lines represent the range of frequencies covered during the experiments).
• the area under the spectrum curve represents the zero-shear viscosity of the melt and as expected is larger for the PP compared to the HDPE.
• Figure 38 presents the DSC heating thermograms for the PP and HDPE monolayer films as well as their multilayer film.
• the PP and HDPE exhibit melting peaks around 162 0 C and 129 0 C, respectively, whereas the multilayer film shows two melting peaks at the same temperatures as the single layers.
• the crystallinities of the components in the multilayer film were slightly lower than that measured for the monolayer films.
• the WAXD patterns as well as the diffraction intensity profiles for the PP and HDPE shown in Figure 39 reveal four and two diffractions, respectively, corresponding to the indicated crystallographic planes.
• the 110 and 040 crystalline planes and for the HDPE the 110 and 200 crystalline planes were used to obtain the orientation of the unit crystal cell axes (a, b, and c) with respect to MD, TD, and ND.
• WAXD cannot be employed for the orientation measurement of the HDPE phase in the multilayer film.
• Figure 40 demonstrates the effect of AFR 1 DR, and annealing on the diffraction patterns as well as on the pole figures of the PP and HDPE monolayer films and the PP in the trilayer film.
• the normal to the 110 plane is the bisector of the a and b axes and 040 and 200 are along the b-axis and a-axis of unit crystal cells, respectively [69]. From the WAXD patterns of the PP monolayer ( Figure 40a), it is clear that by increasing DR and AFR or annealing, the arcs become sharper and more concentrated in the center, implying more orientation.
• annealing causes the 110 plane to be more aligned in TD. Similar trends for the influences of DR, AFR, and annealing on the crystalline alignment of the PP component in the trilayer film are observed ( Figure 40b).
• the four off-axes arcs for the 110 plane of the HDPE are observed in Figure 40c that is a typical behavior of the twisted lamellar structure of PE where the a-axis rotates around the b-axis, resulting in the rotation of the reciprocal vector of the 110 plane.
• Increasing air cooling to the films surface i.e. H-AFR
• H-AFR improves orientation of the 110 plane along TD and reduces the alignment of the 200 plane along MD drastically.
• increasing DR and annealing slightly enhance the orientation of the crystallographic planes of the HDPE monolayer.
• an FTIR technique was used and the results are discussed in the following paragraph.
• the c-axis alignment of the HDPE (both in the monolayer and multilayer) is significantly lower than that of the PP and also the c-axis orientation along MD of the PP and HDPE in the multilayer is lower than in the monolayer made under the same conditions, in accordance with the results discussed in Figure 40.
• the HDPE has much higher crystallinity and larger heat of fusion compared to the PP, resulting in a large heat release during its crystallization. This can explain lower orientation of the PP component in the trilayer than the PP monolayer.
• 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 [77].
• the meridian intensity the formation of more lamellae for the HDPE is obvious.
• the contribution of the shish to the crystalline phase is much less than that of lamellae, confirming the results of Somani et al. [78] for PE and PP.
• Annealing shifts the peak of the PP precursor to lower values, indicating an increase of the long period spacing.
• annealing does not impact the peak position of the HDPE, implying that L p remains mainly unchanged.
• the lamellae thickness, / c could be calculated by multiplying L p by the crystalline fraction (see the legend in figure).
• the values of Lp and / c for the PP precursor film are much smaller than those for the HDPE and increase with annealing.
• HDPE-DR 60 973 7 ( 90 0) 51 0 (1 5) 88 6 ( 4 2) 2 9 (0 1) 165 (1 6) 0014 (0003)
• HDPE-DR 90 1138 8 ( 69 8) 51 7 (2 3) 96 0 ( 5 1) 3 2 (02) 21 6 (04) 0021 (0 002)
• PP/HDPE/PP-DR 60 920 0 ( 85 2) 55 0 (5 1) 1404 (20 0) 5 2 (04) 15 6 (1 2) 0024 (0003)
• PP/HDPE/PP-DR 90 1037 3 (126 0) 604 (3 4) 164 5 (18 9) 40 (05) 18 7 (3 0) 0036 (0004)
• 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 (T ⁇ ), 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 0 C without extension was the optimum annealing condition for PP.
• Figure 45 presents SEM micrographs of the surface of the monolayer microporous membranes prepared at the cold and hot stretching extensions of 55% and 75%, respectively. The details about the optimum cold and hot stretching levels will be discussed later.
• the distribution of interlamellar tie chains in the precursor films may not be uniform [80]. Therefore, it is believed that the micropores are first developed in region with a small amount of tie chains. It is clear that the size of the pores in the HDPE microporous membrane is much greater than in the PP membrane. The longer interlamellar microfibrils (bridges) in the HDPE porous membrane compared to the PP porous membrane are believed to be due to the longer tie chains in the precursor film of the former. It should be mentioned that the surface morphology of the PP/HDPE/PP membrane (not shown) was somewhat similar to the surface structure of the PP monolayer exhibited in Figure 45.
• Figure 46 illustrates SEM micrographs of the cross-section of the multilayer porous membrane.
• Figure 46a it is obvious that a porous HDPE layer has been sandwiched between two porous PP layers while the three layers have almost identical thicknesses.
• Figures 46b and 46c present the interface between the layers at different magnifications. Similar to the surface micrographs, it is clearly recognized that the HDPE layer has much larger pores than the PP layer. Additionally, a reasonable adhesion between the layers is seen, which can be explained by the transcrystallization and penetration of the HDPE lamellae into the PP layer observed in the cross-section image of the trilayer precursor film shown in Figure 44.
• Table 7 presents the mechanical properties along MD and TD as well as the puncture resistance along ND for the microporous membranes.
• the porous membranes have nearby similar tensile responses in MD and, as expected, the tensile properties along TD are considerably smaller than in MD.
• the PP microporous membranes show a much lower strain at break along TD than that the HDPE and multilayer 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.
• 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.
• FIG. 50 illustrates the interfacial morphology of the multilayer film membrane prepared with a cold stretching of 55% followed by a hot stretching of 175%. This is 100% more total stretching than for the one shown in Figure 46. Very large pores, particularly for the HDPE, are seen and that can explain the improvement in WVTR by increasing hot stretching.
• the arrows in Figure 50 show the connection between interlamellar microfibrils and the lamellae in the HDPE layer. At such a high level of extension, it is clear that the microfibrils have been connected to the surrounded lamellae by bundles of small blocks. According to Yu [80], at high stretch levels, lamellae located at the end of the microfibrils break into small blocks and tilt along the stretching direction.
• 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 Figures 40 and 41).
• the HDPE showed a twisted lamellar morphology whereas at high AFR an intermediate structure between twisted and flat kebabs was detected.
• WVTR transmission rate
• the trilayer microporous membranes showed a lower permeability than the monolayer membranes, possibly due to the presence of the interface as well as the lower orientation of the PP and HDPE components in the multilayer film compared to the monolayer films.

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