ISSN 0965545X, Polymer Science, Ser. A, 2013, Vol. 55, No. 10, pp. 595–602. © Pleiades Publishing, Ltd., 2013. Original Russian Text © E.Yu. Rozova, I.S. Kuryndin, V.K. Lavrent’ev, G.K. Elyashevich, 2013, published in Vysokomolekulyarnye Soedineniya, Ser. A, 2013, Vol. 55, No. 10, pp. 1255–1262.
STRUCTURE AND PROPERTIES
Structure and Mechanical Properties of Porous Films Based on Polyethylenes of Different Molecular Masses E. Yu. Rozova*, I. S. Kuryndin, V. K. Lavrent’ev, and G. K. Elyashevich Institute of Macromolecular Compounds, Russian Academy of Sciences, Bol’shoi pr. 31 (V.O.), St. Petersburg, 199004 Russia *email:
[email protected] Received November 6, 2012; Revised Manuscript Received April 22, 2013
Abstract—The structure and mechanical properties of porous films of linear polyethylenes with different molecular masses obtained via a process including four successive stages—melt extrusion, isometric anneal ing, uniaxial extension, and thermal fixation—are studied. The molecularmass range in which microporous structures containing throughflow channels can form is determined. The effects of molecularmass charac teristics of polyethylene on the morphological features of the structures, porosities, sizes of throughflow channels, specific surfaces, and permeabilities of the porous films are analyzed. DOI: 10.1134/S0965545X13090046
INTRODUCTION Polymer microporous films are substantially supe rior to other porous materials: Namely, they are rela tively easy manufacturing and are characterized by a lower cost of the raw materials and a combination of high transport characteristics and good mechanical properties. Uniform polymer films possessing a net work of throughflow channels are widely applied as microporous membranes for filtration of particles in liquids, separators for chemical current sources, dia phragms for electrolytic capacitors, mechanically strength and elastic porous supports for preparation of composite systems with various functional properties, and nano and microcontainers for delivery of photo active components [1–6]. Among the known polymers, polyolefins are the most promising materials for preparation of microporous films because they can form viscous and chemically stable melts. Being flexiblechain poly mers, polyolefins, depending on crystallization condi tions, are able to form a large number of various supramolecular structures determining the properties and, hence, application fields of the films [7]. There are a number of methods to prepare porous structures from flexiblechain polymers (e.g., [4]). One of the most efficient methods for the preparation of porous films of polyolefins, such as polyethylene and polypropylene, is a process based on extrusion of their melts followed by thermal and orientation treat ments [8–10]. The study of the effects of melt crystal lization and annealing conditions on the characteris tics of porous films obtained via this process showed that these characteristics depend on the molecular mass of a polymer [11–14]. However, the data reported in [11–14] were obtained for different grades
of polyolefins; as a result, their comparison presents a problem. In this study, a process including four key stages— extrusion of the polymer melt, annealing of extruded samples under isometric conditions, uniaxial exten sion, and thermal fixation [9]—was used to prepare porous films of linear polyethylene, and the effects of the molecularmass characteristics of the polymer on the formation of the porous structures of films and their transport and mechanical properties were inves tigated. RESEARCH OBJECTS AND ANALYTICAL METHODS Porous films of linear lowpressure PE of commer cial grades were obtained and investigated. Their molecularmass characteristics were determined by GPC (Table 1). At the first stage of sample preparation (Fig. 1), the films were formed on a laboratory extruder (SCAMIA) at a melt temperature of 200°C, a flatslit die clear Table 1. Molecularmass characteristics of PE samples
595
Sample PE1
Brand I0760 (Uzbekistan)
Mw × 10–4 Mw /Mn 7
2.7
PE2 HIZEX 2208J (Japan)
12
4.3
PE3 Stavrolen 276 (Russia)
17
5.0
PE4 ПЭНД 288 (Russia)
25
9.6
PE5 Borealis (FB1350) (Sweden)
28.5
13.0
596
ROZOVA et al. (c)
L02
(b)
b2
Orientation direction b1 L01
(a)
Fig. 1. Structure formation during preparation of porous PE films during the (a) extrusion, (b) annealing, and (c) uniaxialtension (poreformation) stages. Parameters L01 and L02 are the long Xray periods in the extruded and annealed samples, respectively, and b1 and b2 are the lamella thicknesses in the extruded and annealed samples, respectively.
ance of 1.5 mm, and a linear meltflow rate at the die outlet of 0.06 m/s. The degree of melt extension was defined by spindraw ratio λ calculated via the formula λ = v2/v1, where v1 is the meltflow rate at the die outlet and v2 is the takeup speed of the film. During variation in the spindraw ratio in the range 30–200, films were obtained from granules of all indi cated types of PE. The second stage of the process (annealing of the extruded films) was performed for 45 min at 130°C under isometric conditions. At the third stage, porous structures were formed as a result of uniaxial extension of the films obtained via the first two stages, that is, extrusion and annealing. The sizes and numbers of appearing pores depend on the degree and rate of extension. In the present study, this stage was realized for all samples at a degree of extension of 200% and a rate of extension of 400%/min. The porous structures of the films were fixed through keeping at elevated temperatures (110°C) in the extended state for 1 h. Overall porosity Р of the films was measured gravi metrically and calculated via the formula Р = [(ρ – ρf)/ρ] × 100% Here, ρ is the density of linear lowpressure PE and ρf is the density of the porous film determined through weighing. For granules of PEs of the used grades, the ρ values differed by less than 1% and amounted to 950 kg/m3. Permeabilities G of porous samples for liquids were determined from the flow rate of ethanol (liquid spon taneously wetting PE) under pressure through a filtra tion cell. The permeabilities were determined as fol lows: G = V/(SτpК), where V is the volume of liquid passing through a porous film with area S for time τ under pressure p and К is the resistance coefficient of the filtration cell, which depends on the design of the cell and character
izes the ratio of the flow rate of the liquid through the cell with a porous sample to that for the cell without any sample. For the used cell, K = 0.2. Size distributions of throughflow channels were obtained via filtration porosimetry by Poiseuille method [2] with the use of a nonwetting liquid for which the porous film is impermeable in the absence of pressure. An ethanol–water (30 : 70) mixture was used as such a liquid [11]. At a certain threshold of the applied pressure, the sample becomes permeable through its largest pores. The flow rate increases with further increases in pressure because increasingly smaller pores are gradually involved in the permeation process. The poresize distribution was calculated from the pressure dependence of the liquid flow rate. The specific surfaces of porous samples were deter mined through the method of lowtemperature nitro gen sorption–desorption (BET method [15]) on a SorbtometrM device (Novosibirsk, Russia). Smallangle Xray scattering curves were measured with a Kratky camera with an entrance slit of 150 μm; the primary beam divergence was 6.5'. Wideangle scattering studies were performed on a DRON 2.0 setup (Burevestnik, St. Petersburg, Russia) in the transmission mode with CuKα radiation. To determine the characters and degrees of orientation of the sam ples, azimuthal curves of (110) reflection intensity were obtained. The surface structures of porous films were investi gated by scanning electron microscopy on a SCAN JSM35 device (JEOL, Japan). The mechanical characteristics of samples were calculated from the stress–strain curves obtained on a 2166 P5 tensile testing machine (Tochpribor, Ivanovo, Russia). Strength, elastic modulus, and rela tive elongation at break were measured during uniaxial tension of samples at a rate of 100%/min. Elastic recovery ER100 and work recovery AR/AD were deter mined under cyclic loading of the films until the degree of extension reached 100% at a rate of 100%/min. POLYMER SCIENCE
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RESULTS AND DISCUSSION ϕ
Figure 1 illustrates the formation of porous struc tures of PE films obtained as described in [9]. At the stage of extrusion (Fig. 1a), a structure comprising a system of planar folded crystals (lamellae) arranged in parallel to each other and predominantly perpendicu larly to the direction of melt flow is formed [16, 17]. The azimuthal distribution of the axes of the lamellas in the extruded samples is biaxial. The lamellas are connected by a relatively small number of tie chains. The thicknesses and degrees of perfection of lamellas, the numbers of tie chains, and their length distribu tions on the surfaces of lamellae depend on the condi tions of polymer crystallization. As a result of annealing of the crystallized samples at the second stage of the process (Fig. 1b), the crystal structure improves, namely, the lamella thickness increases owing to the drawing of tie chains from amorphous regions into crystallites; that is, these chains become strained. Accordingly, the number and length of tie chains decrease, while the fraction of stretched tie chains increases. The mechanical prop erties of these structures are described as “rigid–elas tic” and are characterized by high elastic moduli and the capability to undergo large reversible elastic defor mations. As was shown in [14], the formation of these structures is a necessary condition for the appearance of pores at the stage of uniaxial extension of the films. During extension of hard elastic samples along their orientation direction (Fig. 1c), the distance between regions of neighboring lamellae connected by bridges of tie chains remains almost the same, while the parts of lamellae unconnected to each other move apart and become bent. This situation leads to the appearance of voids (pores) between lamellas. With increases in the number and sizes of pores due to an increase in the degree of extension, pores coalen scence with each other and form throughflow chan nels (via the percolation mechanism) and the sample becomes permeable to liquids [17]. In this study, the structures of annealed films (hard elastic samples) of PE that have different molecular masses and that are characterized by Xray long period L02; by the degree of crystallinity, χ; and by the degree of orientation described by disorientation angle ϕ between the direction of chains in lamellar crystallites and the direction of melt flow (Fig. 2) were examined via Xray scattering. Small angle Xray scattering experiments showed that degree of crystallinity χ is almost molecular massindependent and equal to ~70%. This parameter is determined by the annealing temperature, which was the same for all samples. At the same time, long period L02 and lamella thickness b2 increase with an increase in molecular mass (Table 2). POLYMER SCIENCE
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Orientation direction
Dependence of Basic Structural Parameters of the Films on the Molecular Mass of Linear PE
Fig. 2. Schematic arrangement of lamellas in the annealed samples.
Wideangle Xray diffractograms of the annealed samples are fourpoint patterns corresponding to biaxially oriented structures. As seen from Table 2, dis orientation angles ϕ for samples PE2, PE3, and PE4 do not depend on molecular mass. (All samples were prepared at the same spindraw ratio and the same degree of extension, that is, under the same ori entation conditions.) However, PE1 and PE5 sam ples have larger ϕ values (Table 2) and, accordingly, lower degrees of orientation (Fig. 2). Polymer PE5, having the highest molecular mass among the investi gated samples, is characterized by the broadest molec ularmass distribution, that is, the highest polydisper sity Mw/Mn (Table 1). According to gelpermeation chromatography, a lowmolecularmass “wing” is present in an appreciable amount in PE5 together with its highmolecularmass fraction. In this case, small crystallites are additionally formed in the inter lamellar space during melt crystallization along with a large lamellar crystal structure [6, 18]. This phenome non prevents uniform orientation of lamellae with respect to the direction of melt flow and explains the observed increase in disorientation angle ϕ for PE5 and, hence, the decrease in the degree of orientation. Sample PE1, having a narrow molecularmass distri bution ( λt, throughflow chan nels appear during subsequent uniaxial tension. The resulting porous films become permeable to liquids, and their permeabilities increase with an increase in λ. It is known that percolation (in this case, the appear ance of throughflow channels) is usually attained if the content of a new phase (overall porosity) is at least 28–30%. Indeed, porosities considerably exceeding this value are observed for permeable PE2, PE3, and PE4 samples (Table 4). The appearance of throughflow channels in PE2, PE3, and PE4 samples is observed at different spin draw ratios λ (Fig. 5). Moreover, the higher the molec ular mass of the polymer, the lower the λ value may be set for its extrusion for the subsequent formation of a permeable microporous structure. With a further increase in λ, permeability increases for all samples. Throughout the studied range of λ values, the levels of permeability were higher for porous films prepared from a higher molecular mass polymer than those for lower molecular mass polymers (Fig. 5). The observed increase in permeability may be related to two processes: an increase in the number of pores and an increase in pore size. Size distributions of throughflow channels were found with the use of fil tration porosimetry (Fig. 6). To determine the effect of molecular mass on pore size, porous films obtained under identical conditions, namely, at λ = 100, were selected. It is seen from Fig. 6 that these distributions are shifted toward larger sizes with an increase in molecular mass. An increase in molecular mass leads to increases in both the most probable (corresponding to the maximum of distribution function) size and the POLYMER SCIENCE
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601
increases in molecular mass (Table 2) result in the for mation of larger pores at the stage of uniaxial exten sion. This conclusion is in good agreement with the data on the sizes of pores obtained by filtration poro simetry (Table 5).
(a)
0.5
Mechanical Properties of Porous Films
1.0
To ascertain the effect of molecular masses of poly mers on the mechanical properties of their polymer films, samples obtained at λ = 100 were investigated (Table 6). As follows from Table 6, breaking strength and elastic modulus increase with an increase in molecular mass, while elongation at break decreases. However, film elasticity remains high.
(b)
0.5
CONCLUSIONS
1.0
(c)
0.5
200
400
600 d, nm
Fig. 6. Size distributions of throughflow channels for films of (a) PE2, (b) PE3, and (c) PE4. Parameter δn/n0 is the ratio of the number of throughflow pores in the size range from di to di + 1 to the number of pores with sizes corresponding to the most probable value of the distribu tion function n0.
maximum size of throughflow channels (Table 5). These changes in permeabilities and sizes of through flow channels are caused above all by specific features of the structures formed at the stages of melt extrusion and annealing. Increases in the long periods and, hence, the thicknesses of amorphous regions with Table 5. Effect of molecular mass on the sizes of through flow channels
PE2 PE3 PE4
Table 6. Mechanical properties of porous films Sample
Pore size, nm Sample
Our studies have made it possible to determine the range of molecular masses of linear PE within which porous structures containing throughflow channels can form in films. The experimental data show that, if the molecular mass of the polymer is below ∼105, an insufficient number of entanglements and, as a conse quence, slipping of chains during melt flow (the melt viscosity for PE1 is 20 times lower than that for PE4) prevent the formation of the spatial network of lamel lae, while insufficient interconnecting and a poor ori entation of lamellas do not promote the appearance of throughflow channels during extension of such sam ples. As to the highest molecular mass sample, PE5, it may be assumed from the molecularmass depen dences of the characteristics of porous samples that the formation of many pores and throughflow chan nels in this sample would be possible if its broad molecularmass distribution did not contain a large “wing” of short chains. The latters crystallize apart from long chains (the socalled crystallization induced segregation [18]), and their crystallites fill the interlamellar space between lamellae, where pores can form during the moving apart of lamellae at the stage of extension [6]. On the basis of dependences of permeability on spindraw ratio λ obtained for samples with different molecular masses, threshold values λt, at which
Breaking Elastic Relative elonga strength, MPa modulus, MPa tion at break, %
the most probable
average
maximum
PE2
60
400
300
180 200 220
210 230 260
1000 1200 1600
PE3
115
650
105
PE4
125
750
52
Note: Errors for these values were determined from five measure ments and amounted to 5%.
Note: The measurement error is 10%. POLYMER SCIENCE
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throughflow channels appear in porous films, were determined. Size distributions of throughflow chan nels were measured, and it was shown that both aver age and maximum sizes of throughflow pores increase with an increase in molecular mass. The used process makes it possible to prepare porous films com bining high permeability with mechanical strength and elasticity. ACKNOWLEDGMENTS This work was supported by the Russian Founda tion for Basic Research, project no. 100300421. REFERENCES 1. M. Mulder, Basic Principles of Membrane Technology (Kluwer Academic, Dordrecht, 1996; Mir, Moscow, 1999). 2. R. E. Kesting, Synthetic Polymeric Membranes (Wiley, New York, 1985; Khimiya, Moscow, 1991). 3. S. S. Zhang, J. Power Sources, No. 1, 351 (2007). 4. A. L. Volynskii and N. F. Bakeev, Highly Dispersed Ori ented State of Polymers (Khimiya, Moscow, 1984) [in Russian]. 5. V. P. Shibaev, A. Bobrovskii, G. K. Elyashevich, E. Rozova, F. Shimkins, V. Shirinyan, A. Bubnov, M. Kaspar, V. Hamplova, and M. Glogarova, Liq. Cryst. 35, 533 (2008). 6. I. Novak, G. K. Elyashevich, I. Chodak, A. S. Olif irenko, M. Steviar, M. Spirkova, N. Saprykina, E. Vla sova, and A. Kleinova, Eur. Polym. J. 44, 2702 (2008).
7. V. A. Marikhin and L. P. Myasnikova, Supramolecular Structure of Polymers (Khimiya, Leningrad, 1977) [in Russian]. 8. I. Brazinsky, W. M. Cooper, and A. S. Gould, US Patent No. 4,138,459 (1979). 9. G. K. Elyashevich, E. Yu. Rozova, and E. A. Karpov, Patent of the Russian Federation No. 2140936 (1997). 10. E. Kamei, H. Ashitaka, and T. Takahashi, US Patent No. 5,173,235 (1992). 11. G. K. Elyashevich, A. G. Kozlov, and E. Yu. Rozova, Polym. Sci., Ser. A 40, 567 (1998). 12. G. K. Elyashevich, I. S. Kuryndin, V. K. Lavrentyev, A. Yu. Bobrovsky, and V. Buko, Phys. Solid State 54, 1907 (2012). 13. S.Y. Lee, S.Y. Park, and H.S. Song, Polymer 47, 3540 (2006). 14. I. K. Park and H. D. Noether, Colloid Polym. Sci. 253, 824 (1975). 15. S. Gregg and K. S. W. Sing, Adsorption, Surface Area and Porosity (Academic, New York, 1982; Mir, Mos cow, 1984). 16. G. K. Elyashevich, E. A. Karpov, and A. G. Kozlov, Macromol. Symp. 147, 91 (1999). 17. M. Raab, J. Scudla, A. G. Kozlov, V. K. Lavrentyev, and G. K. Elyashevich, J. Appl. Polym. Sci. 80, 214 (2001). 18. L. Mandelkern, Crystallization of Polymers (McGraw Hill, New York, 1964; Khimiya, Moscow, 1966).
Translated by A. Yakimanskii
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