- Email: [email protected]
Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers
Accepted Manuscript Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers Kinga Jurczuk, Andrzej Galeski PII: DOI: Reference:
S0014-3057(16)30313-5 http://dx.doi.org/10.1016/j.eurpolymj.2016.04.031 EPJ 7341
To appear in:
European Polymer Journal
Received Date: Accepted Date:
10 February 2016 25 April 2016
Please cite this article as: Jurczuk, K., Galeski, A., Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers, European Polymer Journal (2016), doi: http://dx.doi.org/10.1016/ j.eurpolymj.2016.04.031
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers. Kinga Jurczuk* and Andrzej Galeski Centre of Molecular and Macromolecular Studies, Polish Academy of Sciences, Sienkiewicza 112, 90363 Lodz, Poland *Corresponding author: K. Jurczuk (e-mail: [email protected]) Abstract The effect of poly(tetrafluoroethylene) (PTFE) nanofibers content on the rheological and mechanical properties of thermoplastic elastomer/low density polyethylene (TPE/LDPE) blends was investigated. The PTFE nanofibers were generated in situ during compounding of crystalline PTFE grains with molten matrix based on TPE/LDPE blends. The studies revealed that PTFE nanofibers produced by solid-state deformation of PTFE crystals formed an entangled network which significantly improved both rheological and mechanical properties of TPE/LDPE blends. The maximum strain hardening coefficient was found to depend on the concentration of the PTFE nanofibers. The material based on TPE/LDPE blend containing 5 wt.% of PTFE nanofibers (75/20/5) exhibited strain hardening coefficient about three times higher as compared to conventional TPE/LDPE blend (60/40). Also, the mechanical properties of the blends depended on the content of PTFE nanofibers. Higher the content of PTFE nanofibers the stiffer TPE/LDPE blend. At 8 wt.% of PTFE nanofibers (60/32/8) the modulus of elasticity was two times higher in comparison to 60/40 blend. Keywords: Polymer blends, thermoplastic elastomer, poly(tetrafluoroethylene) nanofibers, rheological properties, mechanical properties
1
1. Introduction Thermoplastic elastomers, TPE, belong to a unique class of polymeric materials that combine properties of thermoset rubbers such as elasticity and high flexibility with processability and recyclability of thermoplastics [1]. Two main groups of TPE can be distinguished: block copolymers and thermoplastic/elastomer blends. Block copolymers are further divided into copolymers with amorphous rigid segments (styrenic block copolymers, SBC) and copolymers with crystalline rigid segments (thermoplastic
polyurethanes,
TPU,
thermoplastic
copolyesters,
COPE,
and
thermoplastic polyamides, PEBA) [2,3]. Thermoplastic/elastomer blends are further divided into blends where the elastomeric phase is not vulcanized (thermoplastic polyolefins, TPO) and blends where elastomeric phase is dynamically vulcanized during mixing (thermoplastic vulcanizates, TPV) [5]. Resent advances in metallocene polymerization catalysts allowed to develop two new classes of TPE, i.e. polyolefin elastomers (POE) and polyolefin plastomers (POP) [5]. The ease of processing and fabrication of TPE compared with cross-linkable elastomers allows to use them in variety of markets and application sectors including automotive and transportation, building and construction, footwear, appliance, housewares, medical and safety, agriculture and off road, electrical and industrial controls, food and beverage, hydraulics and pneumatics, plumbing and irrigation, marine etc. [2,6]. Worldwide demand for TPE is forecast to rise 5.5 percent per year to reach 5.8 million metric tons in 2017, valued at more than $20 billion. Unfortunately, neat TPE resins usually do not meet the existing end-use requirements, and do need to be modified by means of filling. Various fillers, both organic or inorganic, micro- or nano-sized (e.g. calcium carbonate, silica, nanosilica, talc, organoclay, glass fibers, carbon fibers, aramid fibers, cellulose fibers, silk fibers
2
and carbon nanotubes) have been used to improve the strength, stiffness and dimensional stability of TPEs [7-19]. In order to obtain maximum reinforcement, the filler need to be uniformly dispersed and exhibit a good adhesion to the TPE matrix. The higher the filler content, the more difficult to achieve its good dispersion and usually a surface modification of the filler and/or polymer functionalization, as well as addition of compatibilizers, are required [10,14,16,19-22]. Promising alternative to conventionally filled polymers is a preparation of polymer fiber-reinforced composites directly during mixing. Several studies on melt blending of TPE and a thermotropic liquid crystalline polymers (TLCP) have been reported [23-28]. The interest was mainly focused on the fact that under the appropriate conditions for melt blending of TLCP and TPE, the dispersed drops of TLCP can be deformed into microfibrils and frozen in the matrix after rapid cooling resulting in so-called in situ composite [29]. Stability of the morphology of TLCP/TPE blend is an important factor with regard to reprocessing. To preserve the shape of extended threads of TLCP after deformation, i.e. acting against capillary instabilities leading to the TLCP fibers breakup into droplets, still is a big challenge. The concept of converting instead of adding was also used to produce socalled microfibrillar reinforced composites (MFC) [30-34]. MCF material e.g. isotactic polypropylene, PP, reinforced with fibrous poly(ethylene terephthalate) (PET) can be obtained by three-step processing method comprising melt blending of PP with PET and extrusion; spinning of the PP/PET blend into filament of 30 µm diameter; and remelting of PP matrix without melting of microfibrous PET inclusions. The preparation process of MFC has an advantage as compared to traditionally prepared fiber-reinforced composites, i.e. does not require the dispersion stage of ready-made fibers into the matrix.
3
Recently, a one-step method for the preparation of polymer fiber-reinforced micro- and nanocomposites have been shown [35-37]. This was based on reports about the application of PTFE emulsion and powders at a low concentration, below 1% as nucleating agent for crystallization of other polymers [38-41]. Much larger concentration up to 20% of PTFE was used by Jurczuk et al. [35, 42] for production of nanocomposites by transforming PTFE grains into nanofibers by shearing via various viscous molten polymers. Blending with the use of compounders resulted in generation of PTFE nanofibers and thus formation of “all-polymer nanocomposites”. Simplicity of this process consists in a generation of polymer nanofibers in situ during compounding of a molten polymer matrix with solid crystalline particles of another polymer. The grains of highly crystalline polymer with low density of chain entanglements, e.g. nascent poly(tetrafluoroethylene) (PTFE) or disentangled polypropylene (PP) during shearing of the matrix undergo fibrillation. This process is in fact a plastic deformation of polymer crystals embedded in molten matrix. This one-step method allows to form micro- or nanofibers-reinforced materials with significantly improved rheological and mechanical properties [35-37,42]. The aim of this paper is to report the preparation method of TPE/LDPE blends reinforced with PTFE nanofibers. The effect of content of in situ generated PTFE nanofibers on the rheological and mechanical properties of TPE/LDPE blends is here analyzed and discussed.
2. Experimental 2.1 Materials Nascent PTFE powder, Teflon 7C with an average particle size of 28 µm and melting temperature of 346 oC, was supplied by DuPont. Two grades of TPE were
4
used in the study: (1) ethylene-octene copolymer, Infuse 9007 with a melt flow index, MFI = 0.5 g/10min (190 oC, 2.16 kg) and a density, d = 0.866 g/cm3, obtained from Dow Chemical Company; and (2) ethylene-butene copolymer, Koattro KT AR05 with MFI = 0.5 g/10min (190 oC, 2.16 kg) and d = 0.890 g/cm3, supplied by LyondellBasell. Low density polyethylene, LDPE Icorene N2100 with MFI = 2.0 g/10min (190 oC, 2.16 kg) and d = 0.923 g/cm3, was obtained from ICO Polymers (A. Schulman Inc.). 2.2 Preparation of TPE/LDPE blends The TPE/LDPE blends having various component ratio were prepared by mixing in a Brabender internal batch mixer W50E at 170 oC for 10 min with a rotor speed of 100 rpm. In order to reduce an aggregation of PTFE grains during mixing, a masterbatch composed of 20 wt. % of PTFE and 80 wt. % of LDPE was prepared. Afterwards, the LDPE/PTFE masterbatch was diluted in TPE matrix to obtain 3, 5 and 8 wt. % of PTFE in the final material. In must be mentioned that the mixing temperature applied, i.e. 170 oC, was above melting temperature of TPE and LDPE but much below the melting temperature of PTFE. For further studies, 1 mm thick films were prepared by compression molding at 140 oC followed by cooling between to metal blocks kept at room temperature. Samples are further described by the following codes: (TPE content)/(LDPE content)/(PTFE content).
2.3 Characterization 2.3.1 Morphology of TPE/LDPE blends The bulk morphology of materials was studied by means of scanning electron microscopy (SEM) using a microscope JEOL JSM 6010 LV/LA. In order to reveal the internal structure of samples, two separate methods were applied: fracture of
5
samples at the temperature of liquid nitrogen and etching of polymeric matrix. Samples for etching were first cut with an ultramicrotome (Power Tome XL, RMC Boeckeler Instruments, Inc.) equipped with a glass knife, in order to expose a flat and smooth cross-section surface. Then the exposed surface of materials was etched for 72 hours at room temperature in permanganic etchant [43]. It must be mentioned that only polyolefins are susceptible to permanganic etching while PTFE is resistant. The purpose of such prolonged etching was to reveal large fragments of PTFE nanofibers. Afterwards, exposed surfaces of samples were coated with a fine gold layer by ion sputtering (JEOL JFC-1200) and examined with SEM in a high vacuum mode at the accelerating voltage of 15 kV using secondary electron detector (SEI). 2.3.2 Differential scanning calorimetry Thermal properties of the TPE/LDPE blends and neat polymers were determined by differential scanning calorimetry (DSC) using an indium-calibrated calorimeter TA 2920 (TA Instruments). Samples of weight 6-10 mg, cut out from 1 mm thick films, were heated at the rate of 10 oC/min to 180 oC, annealed for 3 min and then cooled at the rate of 10 oC/min to room temperature. The entire thermal treatment was performed under nitrogen flow. 2.3.3 Wide-angle X-ray scattering (WAXS) The crystalline structure of materials was probed by wide-angle X-ray scattering (WAXS). The x-ray radiation was Cu Kα lamp powered by Philips 1840 generator operating at 30 kV and 50 mA. The samples for WAXS measurements were cut from 1 mm-thick films. Two-dimensional scattering patterns with the use of Pilatus 300 K solid state detector of resolution 172 x 172 µm2 (Dectris, Switzerland) were registered at room temperature. The WAXS profiles on the horizontal line were obtained from 2-D WAXS patterns.
6
2.3.4 Rheological properties The rheological behavior of TPE/LDPE blends and neat polymers was examined using a strain-controlled rheometer (ARES LS2, TA Instruments) with a torque transducer (0.02-2000 g·cm) and a normal force transducer (2-2000 g). Three types of experiments were performed: (a) The small-amplitude oscillatory measurements were carried out using 25 mm parallel-plate fixtures with a gap distance of 0.8 mm. The disk-shaped samples were cut out from 1 mm thick films. The boundaries of the linear regime, over which the storage modulus, G’, and loss modulus, G’’, are independent of a strain amplitude were determined for each sample by running strain amplitude sweeps from 0.01 to 10 % at an angular frequency, ω, of 1 rad s-1. The main oscillatory tests were performed at 140 oC in the ω range from 0.1 to 100 rad s-1. The average of three measurements for each sample is reported. (b) Uniaxial extension tests were performed using extensional viscosity fixture (EVF) attached to the rheometer. The 18 mm × 10 mm × 0.7 mm rectangular specimens were prepared by hot-compression molding in the standard mold, provided by TA Instruments, at 140 oC, then left under pressure of 100 bars for 30 min to relax possible molecular orientation with simultaneous cooling to room temperature. Molten samples were uniaxially extended at 130 oC and Hencky strain rate, , of 5.0 s-1. Each sample was tested at least eight times, and the results were averaged. (c) Steady simple shear measurements were executed using a 25 mm cone-plate fixtures with a cone angle of 0.1 rad and a gap distance of 0.046 mm. Disk-shaped samples were cut out from 1 mm thick films. The shear tests were performed at 130 o
C and a shear rate, , of 0.01 s-1.
2.3.5 Mechanical properties
7
Tensile properties of TPE/LDPE blends and neat polymers were measured at room temperature with a tensile testing machine (Instron 5582). The oar-shaped samples according to ISO 527-2 1BA standard were cut out from 1 mm thick films. Tests were performed at a cross-head speed of 12.5 mm/min which corresponds to a strain rate of 50 %/min. The selected tensile parameters were determined by averaging values obtained for at least five specimens.
3. Results and discussion 3.1 Morphology of TPE/LDPE blends In many applications the properties of TPE are tuned by blending with LDPE. In this study, a primary idea was to use PTFE alone for modifying the properties of TPE resins. Unfortunately, PTFE particles strongly tend to aggregate during compounding with thermoplastic polymers. So to avoid this phenomenon we used LDPE in the powder form to produce LDPE/PTFE masterbatch, which was further diluted with TPE resin. In consequence we obtained the TPE/LDPE blend with PTFE. It is known that during compounding of nascent PTFE powder with thermoplastic polymer, highly crystalline PTFE powder particles, subjected to shearing, can readily undergo large deformation, without melting step, and transform into nanofibers provided that the critical conditions for fibrillation are fulfilled [35]. The ability of PTFE to fibrillate is related to the fact that this highly crystalline polymer possesses large crystals with significantly reduced degree of chain entanglements in the amorphous phase. Fibrillation process, being in reality a solid-state deformation of PTFE, depends on many factors including interfacial tension, SL, between solid PTFE particles and molten polymer matrix, which in turns affect on a shear stress transfer from the matrix to PTFE particles. The value of interfacial tension should be
8
as low as possible in order to achieve effective stress transfer. The more effective stress transfer from the polymer matrix to crystalline PTFE inclusions the higher deformation ratios can be achieved. Fig. 1 presents SEM micrograph of fractured surface of the LDPE/PTFE masterbatch. A plentitude of PTFE nanofibers with the transversal sizes (i.e. diameter) ranging from 100 to 900 nm are visible. However, it turned out after detailed analysis of SEM images that thicker PTFE fibers are in fact bundles consisted of several thinner PTFE nanofibers (usually more than two nanofibers). Apparently such bundles are formed from PTFE grains, which consists of several large crystals each one capable of single fiber drawing [35]. It is estimated that the LDPE matrix was subjected to shear with the rate of 960 s-1 during preparation of the masterbatch, i.e. average shear rate in a chamber of Brabender batch mixer W50E operating at 100 rpm. The interfacial tension, SL, between solid PTFE and molten LDPE was at the level of 8.8 mN/m at 170 oC, i.e. the compounding temperature applied for preparation of the LDPE/PTFE masterbatch. The SL between PTFE and LDPE was calculated according to the following formula [44,45]:
SL = SV + LV – F
(1)
F = 2 (dLVdSV)1/2
(2)
where
based on the following parameters: SV = 25.5mN/m at 20 oC, -d/dT = 0.053 mN/mK and dLV = 18.6 mN/m for PTFE [46]; LV = 34.3 mN/m at 20 oC, -d/dT = 0.060 mN/mK and dLV = 32.0 mN/m at 20 oC for LDPE [44].
9
Fig. 1. SEM image of cryo-fractured surface of LDPE/PTFE masterbatch 80/20. The SL value is not very low, however, it was sufficiently low to ensure effective shear stress transfer from LDPE matrix to the particles of nascent PTFE powder and initiate their deformation and transformation into nanofibers. However, the compounding conditions were insufficient to obtain the PTFE nanofibers with very high deformation ratios, and as it was expected a plentitude of PTFE fibers underwent further deformation and were pulled out from LDPE matrix during fracture of the sample prepared for SEM investigation. Fig. 2 shows an exemplary SEM micrographs of the cryogenic fractured surface of Koattro/LDPE blend containing 3 wt. % of PTFE nanofibers. Only the ends of PTFE nanofibers are visible on the fractured surface of Koattro/LDPE blend indicating that PTFE nanofibers generated during preparation of the masterbatch underwent further deformation during dilution of the masterbatch in the TPE matrix and were strained strongly to the hardening stage where further deformation is difficult and requires higher stress. Very strong similar nanofibrillation of PTFE grains was evident also in Infuse/LDPE blends (results not shown here). 10
Fig. 2. SEM images of cryogenic fractured surfaces of Koattro/LDPE blend containing 3 wt.% of PTFE nanofibers (85/12/3).
It was impossible to determine the diameter and length of generated PTFE nanofibers directly from SEM images on fractured surface of TPE/LDPE blends. For this reason the etching of TPE/LDPE was applied. Fig. 3 shows SEM image of the free surface of TPE/LDPE blend with 3 wt.% of PTFE after 72 hours permanganic etching of a layer of Koattro/LDPE matrix as well as size distribution diagram of PTFE nanofibers.
11
Fig. 3. SEM image of free surface of Koattro/LDPE blend containing 3 wt.% of PTFE nanofibers (85/12/3) after 72 hours etching with size distribution diagram of PTFE nanofibers.
Detailed analysis of SEM images revealed that thicker PTFE fibers (i.e. 220-260 nm thick) are located in bundles in comparison with thinner ones (i.e. 100-140 nm). The length of PTFE nanofibers ranging from 30 µm up to few hundreds of µm. During mixing, the grains of nascent PTFE were deformed to very high deformation ratios. Simultaneously, generated nanofibers underwent entangling process. Entanglement knots of PTFE nanofibers work as local constrains for further deformation. However, the fibers which were split out from the bundle and underwent further deformation resulting in much thinner nanofibers in comparison with those present in the bundles. WAXS diffraction profiles of samples of neat Koattro, LDPE and their blends are presented in Fig. 4. It is indicated that Koattro crystallizes in the hexagonal crystallographic units of form I characteristic for poly(1-butene) [47].
Fig. 4. WAXS profiles of ethylene-butene copolymer Koattro, LDPE, Koattro/LDPE blend (60/40) and Koattro/LDPE blend containing 8 wt.% of PTFE nanofibers (60/32/8). Main diffraction peaks are marked on the plots. Diffraction peak from (107) planes of form IV of PTFE is also indicated. 12
This crystallographic form reflects x-ray at 2 = 10.3° and 17.8° assigned to the (110) and (300) peaks while at 2 = 20.7° there are two overlapped peaks from (220) and (211) planes. No signs of crystallization of ethylene component of Koattro is detected in Fig. 4. The diffraction profiles for 60/40 blend indicate that 1-butene and ethylene elements form separate crystals and they do not co-crystallize in the blends. Propably they will crystallize separately in different temperature ranges. The broad diffraction peak for 60/40 blend centered at 21.1o is composed of three reflections: (110) of ethylene orthorhombic form, and (220) and (211) of 1-butene form I. Amorphous hallo is visible on all diffraction profiles for Koattro/LDPE blends as a very broad and low diffraction because it is composed with ethylene and 1-butene amorphous phases with peaks centered at largely different 2 angles. When PTFE nanofibers are present in the blend (60/32/8) the diffraction peak (107) of form IV of PTFE is superimposed on (300) peak of 1-butene form I crystals. The addition of PTFE nanofibers does not influence significantly either crystal sizes of ethylene and 1-butene components (no significant peak broadening of respective X-ray peaks) or degree of crystallinity.
13
Fig. 5. X-ray diffraction profiles of ethylene-octene copolymer Infuse, LDPE, Infuse/LDPE blend (60/40) and Infuse/LDPE blend containing 8 wt. % of PTFE nanofibers (60/32/8). Main diffraction peaks are marked on the plots. Diffraction peak from (107) planes of form IV of PTFE is also indicated.
WAXS diffraction profiles of samples of Infuse and LDPE materials are shown in Fig. 5. It is seen that all of them form orthorhombic polyethylene type crystals, possibly co-crystallizing in the blends. The addition of PTFE nanofibers does not influence significantly either crystal sizes (no significant peak broadening) or degree of crystallinity. Amorphous hallo is visible on all diffraction profiles with a peak at around 2 = 19.5o. If PTFE nanofibers are present in the Infuse/LDPE blend (60/32/8) the amorphous hallo is superimposed with the diffraction peak (107) of form IV of PTFE. 3.2 Thermal properties of TPE/LDPE blends Thermal properties of TPE/LDPE blends and neat polymers were examined. Typical DSC heating and cooling thermograms for Koattro- and Infuse-based materials are shown in Fig. 6 and Fig. 7, respectively. Two endothermic peaks during melting and two exothermic peaks on crystallization were observed for all Koattro/LDPE blends indicating the formation of separate crystals. This observation is expected taking into account the conclusions, based on x-ray diffraction presented in Fig. 4, of separate crystallization of 1-butene crystals of Koattro and ethylene crystals of LDPE. The low temperature peaks are attributed to melting and crystallization of Koattro 1-butene crystals, whereas the higher temperature peaks are associated with melting and crystallization of LDPE crystals, respectively. Addition of 15 wt.% of LDPE in to Koattro matrix caused the increase in the onset of crystallization temperature by 40
o
C. However no significant changes were observed in
Koattro/LDPE blends containing PTFE nanofibers. 14
Fig. 6. DSC thermograms of neat Koattro, neat LDPE and Koattro/LDPE blends with PTFE nanofibers. The thermograms have been shifted vertically for better visualization.
Fig. 7. DSC thermograms of neat Infuse, neat LDPE and Infuse/LDPE blends with PTFE nanofibers. The thermograms have been shifted for better visualization.
Different thermal behavior was observed for Infuse-based materials (see Fig. 7). A single endothermic peak during melting and single exothermic peak during crystallization was evident for all Infuse/LDPE blends. It appeared, therefore, that cocrystallization of Infuse and LDPE took place, which is known for other polyolefin
15
blends such as LDPE with HDPE [48]. This observation is also supported by x-ray diffraction studies presented in Figure 5 of co-crystallization of ethylene elements of LDPE and Infuse. A gradual increase of melting temperature by 6.7
o
C and
crystallization temperature by 10.2 oC with the increasing content of LDPE was observed. Also, the onset of crystallization has been shifted to higher temperatures with increasing amount of LDPE. The temperature increase is caused by longer unperturbed ethylene sequences in LDPE than in Infuse ethylene copolymer. In case of Infuse/LDPE blend containing PTFE nanofibers, two melting peaks and broad crystallization peak with a small shoulder were observed (see DSC thermograms for 60/32/8 composition on Fig. 7). This is an result of PTFE nanofibers ability to accelerate the crystallization of LDPE [38,39].
3.4 Rheological properties The effect of PTFE nanofibers on viscoelastic behavior of TPE/LDPE blends was examined by means of small-amplitude oscillatory shear. Since microstructure of a molten polymeric materials undergoing shear is not significantly deformed, the oscillatory shear measurements is a useful tool to explore the effect of polymeric materials structure on their moduli and viscosity [49]. Exemplary curves of storage modulus, G’, and loss modulus, G’’, as a function of angular frequency, ω, for neat Infuse and Infuse/LDPE blends with 3 wt.% of PTFE nanofibers are presented in Fig. 8.
16
Fig. 8. Storage modulus, G’, and loss modulus, G’’, as a function of angular frequency, ω, for neat Infuse, Infuse/LDPE blend, and Infuse/LDPE blend containing 3 wt.% of PTFE nanofibers. For neat Infuse, both moduli, G’ and G’’, decreases with decreasing ω, and in low frequency range G’’ was larger than G’ demonstrating the viscous nature of Infuse (note log-log scale in Fig. 8). Similar behavior was observed for the Infuse blend with 15 wt.% of LDPE, while Infuse/LDPE blend containing 3 wt.% of PTFE nanofibers showed significantly different viscoelastic response as compared to that of neat Infuse and Infuse/LDPE blend. It appeared that PTFE nanofibers caused the increase of both moduli in the low frequency range but the increase of G’ was more pronounced than G’’. It can be also seen that the G’ curve for 85/12/3 blend tended 17
to reach a low-frequency plateau (solid-like response). We have shown previously that such tendency is attributed to the presence of physical entangled network of solid PTFE nanofibers in a polymer matrix [50]. This can be accurately seen in Fig. 9, presenting the loss factor, tanδ (=G’’/G’) for Infuse-based materials as a function of angular frequency, ω. Both for the neat Infuse and its blend with 15 wt.% of LDPE, the tanδ increases with decreasing ω reaching 3.6 and 3.2 at 0.1 rad s -1, respectively. In the case of 85/12/3 blend, a broad peak on the tan δ curve appeared in the low ω range as a result of slow relaxation of Infuse chains exerted by entangled network of PTFE nanofibers.
Fig. 9. Loss factor, tanδ, as a function of angular frequency, ω, for neat Infuse, Infuse/LDPE 85/15 blend and 85/12/3 blend containing 3 wt.% of PTFE nanofibers.
Shear
experiments
are
not
sufficient
for
comprehensive
rheological
characterization of polymeric materials. This is particularly true for all the processing operations where the extension of a polymer melt has to be considered. Uniaxial extension of molten polymeric materials is often used to predict their rheological behavior outside the regime of linear viscoelasticity (LVE) which is very sensitive to structural differences within the materials. Fig. 10 presents a time-dependence of 18
tensile stress growth coefficient,
, for Infuse-based materials recorded during
uniaxial extension. It can be seen that for neat Infuse, the tensile stress growth coefficient,
, gradually increases with time, which is a consequence of
stretching of polymer chains, and ultimately reaches plateau when the chains are nearly fully stretched [51].
Fig. 10. The tensile stress growth coefficient, as a function of time, t, for neat Infuse, Infuse/LDPE blends and Infuse/LDPE blends with 3 and 5 wt.% of PTFE nanofibers, recorded during uniaxial extension at 130 oC and at Hencky strain rate, , of 5 s-1. Solid line represents the threefold of the shear stress growth coefficient, for neat Infuse measured at the shear rate, , of 0.01 s-1. The
curve superimposed on the shear stress growth coefficient,
curve multiplied by a factor of 3 confirming that neat Infuse exhibited LVE response at the applied Hencky strain rate, . It is known that linear polymers fulfill Trouton’s law, i.e. the ratio of
to
equals to 3 in the linear regime of
deformation [52], and therefore they exhibit no strain hardening themselves, i.e. no significant increase of
above the LVE region can be observed during uniaxial
extensional deformation. Pronounced strain hardening can be induced by
19
introduction of long-chain branches into linear polymer [53-56]. Despite the fact that Infuse is ethylene-octene copolymer with a small amount of short-chain branches, no deviation of
from the LVE response was observed within the experimental
window. Such behavior is similar to that of linear low density polyethylene, LLDPE, which show the strain hardening only at low strain rates, e.g. below 0.1 s-1, whereas no strain hardening can be observed at high strain rates [57,58]. The rheological behavior similar to that of neat Infuse was observed for the Infuse/LDPE blend containing 15 wt. % of LDPE. Neat LDPE exhibit strain hardening itself due to the presence of long-chain branches, which makes a local obstruction for the chain stretching in the direction of extensional deformation [53,59,60]. However, the amount of LDPE added into Infuse matrix was insufficient to induce the strain hardening. In contrast to neat Infuse and 85/15 blend, the presence of 3 wt. % of PTFE nanofibers in the form of entangled network in the Infuse/LDPE blend 85/12/3 caused, after a certain period of time of the uniaxial extensional deformation, a rapid increase of
above the LVE region. The entangled network of PTFE
nanofibers was deformed during uniaxial extension, but entanglement knots, resisted to stretching, caused strain hardening of the material. Further enhancement of the strain hardening can be obtained by either increasing the content of PTFE nanofibers or LDPE, however, the upward deviation of
from LVE response is much more
pronounced for the 75/20/5 blend containing 5 wt.% of PTFE nanofibers as compared to that of the 75/25 and 60/40 blends. Similar rheological behavior during extensional deformation was observed for Koattro-based materials (results not shown here). To estimate the degree of strain hardening effect, the time-dependent strain hardening coefficient,
, defined as a ratio of
20
to
, was
determined. Fig. 11 presents the bar plots of the maximum strain hardening coefficients,
, i.e. the values achieved at maximum Hencky strain for all
studied materials.
Fig. 11. The maximum strain hardening coefficients, , for neat TPEs, TPE/LDPE blends and the blends with PTFE nanofibers.
It is known that molten LDPE exhibit the strain hardening phenomenon under extensional flow due to presence of the long chain branching [61]. So, the all TPE/LDPE blends shows higher values of The
compared to neat TPE matrices.
increases with increasing content of LDPE in the blend reaching 2,5
for Koattro/LDPE 60/40 blend. It appeared that the entangled network of PTFE nanofibers caused further improvement of the strain hardening phenomenon in all studied TPE/LDPE blends. However, it seems that the magnitude of
depended on the type of the TPE
matrix used. The strongest effect of PTFE nanofibers on the strain hardening was observed for the Infuse-based blends, where 5 wt.% of PTFE nanofibers were sufficient to cause a four and sixfold increase of Infuse/LDPE 75/25 blend and neat Infuse.
21
as compared to that of
Mechanical properties Mechanical properties of selected materials were examined by tensile drawing. Tensile parameters, i.e. Yield stress, σs, stress at break, σb, elongation at break, εb, and modulus of elasticity, E, of neat TPEs, their blends with LDPE and PTFE nanofibers are presented in Fig. 12 and Fig. 13.
Fig.12 Selected tensile parameters and modulus of elasticity for neat Koattro, its blends with LDPE and PTFE nanofibers. All Koattro/LDPE blends exhibit higher values of σs and E compared to neat TPE matrix. The higher the content of LDPE in Koattro/LDPE blends the higher the σs and E were obtained. For Koattro/LDPE 60/40 blend, the σs and E are about three times higher in comparison with neat Koattro. It can also be seen that σb drastically decreased for blends containing 25 wt.% LDPE or more, as well as εb which apparently results from worse compatibility of the blend containing higher content of LDPE. Presence of PTFE nanofibers in Koattro/LDPE blends resulted in further 22
increase of σs and E, whereas σb and εb decrease. The highest value of E, i.e. was obtained for Koattro/LDPE blend containing 8 wt.% of PTFE nanofibers, which is about three and a half higher than that for neat Koattro.
Fig.13 Selected tensile parameters and modulus of elasticity for neat Infuse, its blends with LDPE and PTFE nanofibers.
In case of Infuse-based materials, blending with LDPE also improved mechanical properties by increasing σ s and E (see Fig. 13), however better compatibility between Infuse and LDPE resulted in no decrease in elongation. It need to be mentioned that the specimens for Infuse/LDPE blends having LDPE content up to 40 wt. % did not break down during tensile drawing conditions, i.e. ε b obtained was above 2500 %. Modification of the blend system through entangled network of PTFE nanofibers enabled further improvement of materials stiffness. The best results were obtained for Infuse/LDPE blend containing 8 wt.% of PTFE nanofibers, where σs increased by eight times and E by twelve and a half time compared to neat Infuse matrix as well as two times higher compared to reference Infuse/LDPE 60/40 blend. However, the material containing 8 wt.% of PTFE nanofibers drastically lost its drawability, εb of only 300 % was obtained. It is seen that the TPE/LDPE blends containing PTFE nanofibers were significantly stiffer and slightly stronger compared to two component blend systems 23
but brittle with lower elongation at break, which decreased with increasing amount of PTFE nanofibers. Similar behavior was observed in conventional composite of polyamide 6, PA6, with glass fibers, GF, which exhibit increased modulus of elasticity and tensile strength whereas elongation at break drastically decreased [62].
4. Conclusions In this paper the effect of in situ generated PTFE nanofibers on the rheological and mechanical properties of TPE/LDPE blends was examined. Formation of PTFE nanofibers has been observed in all studied TPE/LDPE blends, provided that sufficient shear stress was imposed on the molten materials during compounding. The entangled network of PTFE nanofibers drastically improved the rheological properties of studied TPE/LDPE blends. The strain hardening was found to depend on concentration of the PTFE nanofibers and increased with increasing content of PTFE. The most pronounced effect was observed for the Infuse/LDPE blend containing 5 wt.% of PTFE nanofibers (75/20/5), where the maximum strain hardening coefficient was three times in comparison with conventional TPE/LDPE blend (60/40). Also, PTFE nanofibers improved the mechanical properties of TPE/LDPE blends. It appeared the stiffness and drawability depend on the content of PTFE nanofibers. Higher the concentration of PTFE nanofibers in TPE/LDPE blend the higher the modulus of elasticity and lower elongation at break. The modulus of elasticity of the material based on Infuse/LDPE blend containing 8 wt.% of PTFE nanofibers (60/32/8), was two times higher in comparison with 60/40 blend.
24
Acknowledgments The project of National Science Centre, Poland (DEC-2012/04/A/ST5/00606) is gratefully acknowledged for the partial financial support of the studies. Also, the project of National Centre for Research and Development, Poland (INNOTECHK3/IN3/24/227425/NCBR/14) and the statutory fund of the Centre of Molecular and Macromolecular Studies, Polish Academy of Sciences, are acknowledged.
References [1] Holden G, Understanding thermoplastic elastomers, Hanser, Munich (2000). [2] Spontak RJ and Patel NP., Curr. Opin. Colloid Interface Sci. 5:334-341 (2000). [3] Schmaltz H, Abetz V and Lange R, Compos. Sci. Technol. 63:1179-1186 (2003). [4] Coran AY, Das B and Patel RP, US Patent 4130535 (1978). [5] Scheirs J and Kaminsky W, Metallocene-Based Polyolefins. Preparation, Properties and Technology. Vol. 2. John Wiley & Sons, New York (2000). [6] Amin S and Amin M, Rev. Adv. Mater. Sci. 29:15-30 (2011). [7] Jancar J and Dibenedetto AT, J. Mater. Sci. 30:1601-1608 (1995). [8] Katbab AA, Nazockdast H, and Bazgir S, J. Appl. Polym. Sci. 75:1127-1137 (2000). [9] Bazgir S, Katbab AA, and Nazockdast H, J. Appl. Polym. Sci. 92:2000-2007 (2004). [10] Karakaya N, Ersoy OG, Oral MA, Gonul T, and Deniz V, Polym. Eng. Sci. 50:677-688 (2010). [11] Sreekanth MS, Joseph S, Mhaske ST, Mahanwar PA, and Bambole VA, J. Thermoplast. Compos. Mater. 24:317-331 (2011). [12] Bajsic EG, Rek V, and Pavic BO, J. Elastom. Plast. 45:501-522 (2012). [13] Shonaike GO, and Matsuo T, Compos. Struct. 32:445-451 (1995). [14] Ibarra L, and Panos D, Polym. Inter. 43:251-259 (1997). [15] Anuar H, and Zuraida A, Compos. Part B Eng. 42:462-465 (2011). [16] Saikrasun S, Amornsakchai T, Sirisinha C, Meesiri W, and Bualek-Limcharoen S, Polymer 40:6437-6442 (1999). [17] Vajrasthira C, Amornsakchai T, and Bualek-Limcharoen S, J. Appl. Polym. Sci. 87:1059-1067 (2003).
25
[18] Akhtar S, De PP, and De S, J. Appl. Polym. Sci. 32:5123-46 (1986). [19] Austin JR, and Kontopoulou M, Polym. Eng. Sci. 46:1491-1501 (2006). [20] Liu Y, and Kontopoulou M, J. Vinyl Addit. Technol. 13:147-150 (2007). [21] Li C, Deng H, Wang K, Zhang Q, Chen F, and Fu Q, J. Appl. Polym. Sci. 121:2104-2112 (2011). ]22] Coffey AB, and O’Bradaigh CM, J. Mater. Sci. 42:8053-8061 (2007). [23] Verhoogt H, Langelaan C, VanDam J, and DeBoer AP, Polym. Eng. Sci. 33:754-763 (1993). [24] Verhoogt H, Willems CRJ, VanDam J, and DeBoer AP, Polym. Eng. Sci. 34:453-60 (1994). [25] Machiels AGC, Denys KFJ, VanDam J, and DeBoer AP, Polym. Eng. Sci. 36:2451-2466 (1996). [26] Saikrasun S, Bualek-Limcharoen S, Kohjiya S, and Urayama K, J. Appl. Polym. Sci. 89:2676-2685 (2003). [27] Saikrasun S, Bualek-Limcharoen S, Kohjiya S, and Urayama K, J. Polym. Sci. Part B Polym. Phys. 43:135-144 (2005). [28] Saikrasun S, and Amornsakchai T, J. Appl. Polym. Sci. 103:917-927 (2007). [29] Shin BY, and Chung IJ, Polym. J. 21:851-861 (1989). [30 ] Fakirov S, Evstatiev M, and Petrovich S, Macromolecules 26:5219-5226 (1993). [31] Evstatiev M, Schultz JM, Petrovich S, Georgiev G, Fakirov S, and Friedrich K, J. Appl. Polym. Sci. 67:723-737 (1998). [32] Friedrich K, Evstatiev M, Fakirov S, Evstatiev O, Ishii M, and Harrass M, Compos. Sci. Technol. 65:107-116 (2005). [33] Fakirov S, Bhattacharyya D, Shields RJ, Colloid Surf. A Physicochem. Eng. Asp. 313-314:2-8 (2008). [34] Fakirov S, Compos. Sci. Technol. 89:211-225 (2013). [35] Jurczuk K, Galeski A, and Piorkowska E, Polymer 54:4617-4628 (2013). [36] Krajenta J, Pawlak A, and Galeski A, Polimery 60:64-66 (2015). [37] Antonioli D, Sparnacci K, Laus M, Boarino L, and Righetti MC, J. Appl. Polym. Sci. 132:42401 (2015). [38] Masirek R, Piorkowska E, and Galeski A, Polish Patent PL209925. [39] Bernland K, and Smith P, J. Appl. Polym. Sci. 114:281-289 (2009). [40] Van der Meer DW, Milazzo D, Sanguineti A, and Vancso GJ, Polym. Eng. Sci. 45:458-468 (2005). [41] Ali MABM, Nobukawa S, and Yamaguchi M, Pure Appl. Chem. 83:1819-1830 (2011).
26
[42] Jurczuk K, Galeski A, and Piorkowska E, J. Rheol. 58:589-605 (2014). [43] Galeski A, Bartczak Z, and Kazimierczak T, Polymer 51:5780-5787 (2010). [44] Owens DK, and Wendt RC, J. Appl. Polym. Sci. 13:1741-1747 (1969). [45] Mark JE, Physical properties of polymers. Handbook. 2nd edition, Springer, New York (2007). [46] Dettre RH, and Johnson Jr. RE, J. Colloid Interf. Sci. 31:568-569 (1969). [47] Al-Hussein M, and Strobl G, Macromolecules 35:8515-8520 (2002)]. [48] Fonseca CA, and Harrison IR, Thermochimica Acta 313:37-42 (1998). [49] Larson RG, The structure and rheology of complex fluids. Oxford University Press, New York (1999). [50] Gupta RK, Polymer and composite rheology. 2nd edition, Marcel Dekker, Inc., New York (2000). [51] Marrucci G, and Ianniruberto G, Macromolecules 37:3934-3942 (2004). [52] Münstedt H, Rheol. Acta 14:1077-1088 (1975). [53] Laun HM, and Münstedt H, Rheol. Acta 17:415-425 (1978). [54] Ishizuka O, and Koyama K, Polymer 21:164-170 (1983). [55] Kurzbeck S, Oster F, Münstedt H, Nguyen TQ, and Gensler R, J. Rheol. 43:359–375 (1999). [56] Kasehagen LJ, and Macosko CW, J. Rheol. 42:1303–1327 (1998). [57] Münstedt H, KurzbeckS, and Egersdörfer L, Rheol. Acta 37:21-29 (1998). [58] Münstedt H, Kurzbeck S, and Stange J, Polym. Eng. Sci. 46:1190-1195 (2006). [59] Münstedt H, and Auhl D, J. Non-Newtonian Fluid Mech. 128:62–69 (2005). [60] Wagner MH, Bastian H, Hachmann P, Meissner J, Kurzbeck S, Münstedt H, and Langouche F, Rheol. Acta 39:97–109 (2000). [61] Gabriel C, and Mϋnstedt, J. Rheol. 47:619-630 (2003). [62] Thomason JL, Compos. Part A 39:1618-1624 (2008).
27
Graphical Abstract Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers. Kinga Jurczuk* and Andrzej Galeski
TPE/LDPE blends reinforced with PTFE nanofibers produced by solid-state deformation of PTFE crystals.
28
Higlights: TPE/LDPE blends reinforced with PTFE nanofibers; PTFE nanofibers generated in situ during compounding of crystalline PTFE grains with molten matrix based on TPE/LDPE blends; Strain hardening of molten TPE/LDPE blends improved by entangled network of PTFE nanofibers; Significant improvement of the mechanical properties TPE/LDPE blends containing PTFE nanofibers;
29
S0014-3057(16)30313-5 http://dx.doi.org/10.1016/j.eurpolymj.2016.04.031 EPJ 7341
To appear in:
European Polymer Journal
Received Date: Accepted Date:
10 February 2016 25 April 2016
Please cite this article as: Jurczuk, K., Galeski, A., Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers, European Polymer Journal (2016), doi: http://dx.doi.org/10.1016/ j.eurpolymj.2016.04.031
This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers. Kinga Jurczuk* and Andrzej Galeski Centre of Molecular and Macromolecular Studies, Polish Academy of Sciences, Sienkiewicza 112, 90363 Lodz, Poland *Corresponding author: K. Jurczuk (e-mail: [email protected]) Abstract The effect of poly(tetrafluoroethylene) (PTFE) nanofibers content on the rheological and mechanical properties of thermoplastic elastomer/low density polyethylene (TPE/LDPE) blends was investigated. The PTFE nanofibers were generated in situ during compounding of crystalline PTFE grains with molten matrix based on TPE/LDPE blends. The studies revealed that PTFE nanofibers produced by solid-state deformation of PTFE crystals formed an entangled network which significantly improved both rheological and mechanical properties of TPE/LDPE blends. The maximum strain hardening coefficient was found to depend on the concentration of the PTFE nanofibers. The material based on TPE/LDPE blend containing 5 wt.% of PTFE nanofibers (75/20/5) exhibited strain hardening coefficient about three times higher as compared to conventional TPE/LDPE blend (60/40). Also, the mechanical properties of the blends depended on the content of PTFE nanofibers. Higher the content of PTFE nanofibers the stiffer TPE/LDPE blend. At 8 wt.% of PTFE nanofibers (60/32/8) the modulus of elasticity was two times higher in comparison to 60/40 blend. Keywords: Polymer blends, thermoplastic elastomer, poly(tetrafluoroethylene) nanofibers, rheological properties, mechanical properties
1
1. Introduction Thermoplastic elastomers, TPE, belong to a unique class of polymeric materials that combine properties of thermoset rubbers such as elasticity and high flexibility with processability and recyclability of thermoplastics [1]. Two main groups of TPE can be distinguished: block copolymers and thermoplastic/elastomer blends. Block copolymers are further divided into copolymers with amorphous rigid segments (styrenic block copolymers, SBC) and copolymers with crystalline rigid segments (thermoplastic
polyurethanes,
TPU,
thermoplastic
copolyesters,
COPE,
and
thermoplastic polyamides, PEBA) [2,3]. Thermoplastic/elastomer blends are further divided into blends where the elastomeric phase is not vulcanized (thermoplastic polyolefins, TPO) and blends where elastomeric phase is dynamically vulcanized during mixing (thermoplastic vulcanizates, TPV) [5]. Resent advances in metallocene polymerization catalysts allowed to develop two new classes of TPE, i.e. polyolefin elastomers (POE) and polyolefin plastomers (POP) [5]. The ease of processing and fabrication of TPE compared with cross-linkable elastomers allows to use them in variety of markets and application sectors including automotive and transportation, building and construction, footwear, appliance, housewares, medical and safety, agriculture and off road, electrical and industrial controls, food and beverage, hydraulics and pneumatics, plumbing and irrigation, marine etc. [2,6]. Worldwide demand for TPE is forecast to rise 5.5 percent per year to reach 5.8 million metric tons in 2017, valued at more than $20 billion. Unfortunately, neat TPE resins usually do not meet the existing end-use requirements, and do need to be modified by means of filling. Various fillers, both organic or inorganic, micro- or nano-sized (e.g. calcium carbonate, silica, nanosilica, talc, organoclay, glass fibers, carbon fibers, aramid fibers, cellulose fibers, silk fibers
2
and carbon nanotubes) have been used to improve the strength, stiffness and dimensional stability of TPEs [7-19]. In order to obtain maximum reinforcement, the filler need to be uniformly dispersed and exhibit a good adhesion to the TPE matrix. The higher the filler content, the more difficult to achieve its good dispersion and usually a surface modification of the filler and/or polymer functionalization, as well as addition of compatibilizers, are required [10,14,16,19-22]. Promising alternative to conventionally filled polymers is a preparation of polymer fiber-reinforced composites directly during mixing. Several studies on melt blending of TPE and a thermotropic liquid crystalline polymers (TLCP) have been reported [23-28]. The interest was mainly focused on the fact that under the appropriate conditions for melt blending of TLCP and TPE, the dispersed drops of TLCP can be deformed into microfibrils and frozen in the matrix after rapid cooling resulting in so-called in situ composite [29]. Stability of the morphology of TLCP/TPE blend is an important factor with regard to reprocessing. To preserve the shape of extended threads of TLCP after deformation, i.e. acting against capillary instabilities leading to the TLCP fibers breakup into droplets, still is a big challenge. The concept of converting instead of adding was also used to produce socalled microfibrillar reinforced composites (MFC) [30-34]. MCF material e.g. isotactic polypropylene, PP, reinforced with fibrous poly(ethylene terephthalate) (PET) can be obtained by three-step processing method comprising melt blending of PP with PET and extrusion; spinning of the PP/PET blend into filament of 30 µm diameter; and remelting of PP matrix without melting of microfibrous PET inclusions. The preparation process of MFC has an advantage as compared to traditionally prepared fiber-reinforced composites, i.e. does not require the dispersion stage of ready-made fibers into the matrix.
3
Recently, a one-step method for the preparation of polymer fiber-reinforced micro- and nanocomposites have been shown [35-37]. This was based on reports about the application of PTFE emulsion and powders at a low concentration, below 1% as nucleating agent for crystallization of other polymers [38-41]. Much larger concentration up to 20% of PTFE was used by Jurczuk et al. [35, 42] for production of nanocomposites by transforming PTFE grains into nanofibers by shearing via various viscous molten polymers. Blending with the use of compounders resulted in generation of PTFE nanofibers and thus formation of “all-polymer nanocomposites”. Simplicity of this process consists in a generation of polymer nanofibers in situ during compounding of a molten polymer matrix with solid crystalline particles of another polymer. The grains of highly crystalline polymer with low density of chain entanglements, e.g. nascent poly(tetrafluoroethylene) (PTFE) or disentangled polypropylene (PP) during shearing of the matrix undergo fibrillation. This process is in fact a plastic deformation of polymer crystals embedded in molten matrix. This one-step method allows to form micro- or nanofibers-reinforced materials with significantly improved rheological and mechanical properties [35-37,42]. The aim of this paper is to report the preparation method of TPE/LDPE blends reinforced with PTFE nanofibers. The effect of content of in situ generated PTFE nanofibers on the rheological and mechanical properties of TPE/LDPE blends is here analyzed and discussed.
2. Experimental 2.1 Materials Nascent PTFE powder, Teflon 7C with an average particle size of 28 µm and melting temperature of 346 oC, was supplied by DuPont. Two grades of TPE were
4
used in the study: (1) ethylene-octene copolymer, Infuse 9007 with a melt flow index, MFI = 0.5 g/10min (190 oC, 2.16 kg) and a density, d = 0.866 g/cm3, obtained from Dow Chemical Company; and (2) ethylene-butene copolymer, Koattro KT AR05 with MFI = 0.5 g/10min (190 oC, 2.16 kg) and d = 0.890 g/cm3, supplied by LyondellBasell. Low density polyethylene, LDPE Icorene N2100 with MFI = 2.0 g/10min (190 oC, 2.16 kg) and d = 0.923 g/cm3, was obtained from ICO Polymers (A. Schulman Inc.). 2.2 Preparation of TPE/LDPE blends The TPE/LDPE blends having various component ratio were prepared by mixing in a Brabender internal batch mixer W50E at 170 oC for 10 min with a rotor speed of 100 rpm. In order to reduce an aggregation of PTFE grains during mixing, a masterbatch composed of 20 wt. % of PTFE and 80 wt. % of LDPE was prepared. Afterwards, the LDPE/PTFE masterbatch was diluted in TPE matrix to obtain 3, 5 and 8 wt. % of PTFE in the final material. In must be mentioned that the mixing temperature applied, i.e. 170 oC, was above melting temperature of TPE and LDPE but much below the melting temperature of PTFE. For further studies, 1 mm thick films were prepared by compression molding at 140 oC followed by cooling between to metal blocks kept at room temperature. Samples are further described by the following codes: (TPE content)/(LDPE content)/(PTFE content).
2.3 Characterization 2.3.1 Morphology of TPE/LDPE blends The bulk morphology of materials was studied by means of scanning electron microscopy (SEM) using a microscope JEOL JSM 6010 LV/LA. In order to reveal the internal structure of samples, two separate methods were applied: fracture of
5
samples at the temperature of liquid nitrogen and etching of polymeric matrix. Samples for etching were first cut with an ultramicrotome (Power Tome XL, RMC Boeckeler Instruments, Inc.) equipped with a glass knife, in order to expose a flat and smooth cross-section surface. Then the exposed surface of materials was etched for 72 hours at room temperature in permanganic etchant [43]. It must be mentioned that only polyolefins are susceptible to permanganic etching while PTFE is resistant. The purpose of such prolonged etching was to reveal large fragments of PTFE nanofibers. Afterwards, exposed surfaces of samples were coated with a fine gold layer by ion sputtering (JEOL JFC-1200) and examined with SEM in a high vacuum mode at the accelerating voltage of 15 kV using secondary electron detector (SEI). 2.3.2 Differential scanning calorimetry Thermal properties of the TPE/LDPE blends and neat polymers were determined by differential scanning calorimetry (DSC) using an indium-calibrated calorimeter TA 2920 (TA Instruments). Samples of weight 6-10 mg, cut out from 1 mm thick films, were heated at the rate of 10 oC/min to 180 oC, annealed for 3 min and then cooled at the rate of 10 oC/min to room temperature. The entire thermal treatment was performed under nitrogen flow. 2.3.3 Wide-angle X-ray scattering (WAXS) The crystalline structure of materials was probed by wide-angle X-ray scattering (WAXS). The x-ray radiation was Cu Kα lamp powered by Philips 1840 generator operating at 30 kV and 50 mA. The samples for WAXS measurements were cut from 1 mm-thick films. Two-dimensional scattering patterns with the use of Pilatus 300 K solid state detector of resolution 172 x 172 µm2 (Dectris, Switzerland) were registered at room temperature. The WAXS profiles on the horizontal line were obtained from 2-D WAXS patterns.
6
2.3.4 Rheological properties The rheological behavior of TPE/LDPE blends and neat polymers was examined using a strain-controlled rheometer (ARES LS2, TA Instruments) with a torque transducer (0.02-2000 g·cm) and a normal force transducer (2-2000 g). Three types of experiments were performed: (a) The small-amplitude oscillatory measurements were carried out using 25 mm parallel-plate fixtures with a gap distance of 0.8 mm. The disk-shaped samples were cut out from 1 mm thick films. The boundaries of the linear regime, over which the storage modulus, G’, and loss modulus, G’’, are independent of a strain amplitude were determined for each sample by running strain amplitude sweeps from 0.01 to 10 % at an angular frequency, ω, of 1 rad s-1. The main oscillatory tests were performed at 140 oC in the ω range from 0.1 to 100 rad s-1. The average of three measurements for each sample is reported. (b) Uniaxial extension tests were performed using extensional viscosity fixture (EVF) attached to the rheometer. The 18 mm × 10 mm × 0.7 mm rectangular specimens were prepared by hot-compression molding in the standard mold, provided by TA Instruments, at 140 oC, then left under pressure of 100 bars for 30 min to relax possible molecular orientation with simultaneous cooling to room temperature. Molten samples were uniaxially extended at 130 oC and Hencky strain rate, , of 5.0 s-1. Each sample was tested at least eight times, and the results were averaged. (c) Steady simple shear measurements were executed using a 25 mm cone-plate fixtures with a cone angle of 0.1 rad and a gap distance of 0.046 mm. Disk-shaped samples were cut out from 1 mm thick films. The shear tests were performed at 130 o
C and a shear rate, , of 0.01 s-1.
2.3.5 Mechanical properties
7
Tensile properties of TPE/LDPE blends and neat polymers were measured at room temperature with a tensile testing machine (Instron 5582). The oar-shaped samples according to ISO 527-2 1BA standard were cut out from 1 mm thick films. Tests were performed at a cross-head speed of 12.5 mm/min which corresponds to a strain rate of 50 %/min. The selected tensile parameters were determined by averaging values obtained for at least five specimens.
3. Results and discussion 3.1 Morphology of TPE/LDPE blends In many applications the properties of TPE are tuned by blending with LDPE. In this study, a primary idea was to use PTFE alone for modifying the properties of TPE resins. Unfortunately, PTFE particles strongly tend to aggregate during compounding with thermoplastic polymers. So to avoid this phenomenon we used LDPE in the powder form to produce LDPE/PTFE masterbatch, which was further diluted with TPE resin. In consequence we obtained the TPE/LDPE blend with PTFE. It is known that during compounding of nascent PTFE powder with thermoplastic polymer, highly crystalline PTFE powder particles, subjected to shearing, can readily undergo large deformation, without melting step, and transform into nanofibers provided that the critical conditions for fibrillation are fulfilled [35]. The ability of PTFE to fibrillate is related to the fact that this highly crystalline polymer possesses large crystals with significantly reduced degree of chain entanglements in the amorphous phase. Fibrillation process, being in reality a solid-state deformation of PTFE, depends on many factors including interfacial tension, SL, between solid PTFE particles and molten polymer matrix, which in turns affect on a shear stress transfer from the matrix to PTFE particles. The value of interfacial tension should be
8
as low as possible in order to achieve effective stress transfer. The more effective stress transfer from the polymer matrix to crystalline PTFE inclusions the higher deformation ratios can be achieved. Fig. 1 presents SEM micrograph of fractured surface of the LDPE/PTFE masterbatch. A plentitude of PTFE nanofibers with the transversal sizes (i.e. diameter) ranging from 100 to 900 nm are visible. However, it turned out after detailed analysis of SEM images that thicker PTFE fibers are in fact bundles consisted of several thinner PTFE nanofibers (usually more than two nanofibers). Apparently such bundles are formed from PTFE grains, which consists of several large crystals each one capable of single fiber drawing [35]. It is estimated that the LDPE matrix was subjected to shear with the rate of 960 s-1 during preparation of the masterbatch, i.e. average shear rate in a chamber of Brabender batch mixer W50E operating at 100 rpm. The interfacial tension, SL, between solid PTFE and molten LDPE was at the level of 8.8 mN/m at 170 oC, i.e. the compounding temperature applied for preparation of the LDPE/PTFE masterbatch. The SL between PTFE and LDPE was calculated according to the following formula [44,45]:
SL = SV + LV – F
(1)
F = 2 (dLVdSV)1/2
(2)
where
based on the following parameters: SV = 25.5mN/m at 20 oC, -d/dT = 0.053 mN/mK and dLV = 18.6 mN/m for PTFE [46]; LV = 34.3 mN/m at 20 oC, -d/dT = 0.060 mN/mK and dLV = 32.0 mN/m at 20 oC for LDPE [44].
9
Fig. 1. SEM image of cryo-fractured surface of LDPE/PTFE masterbatch 80/20. The SL value is not very low, however, it was sufficiently low to ensure effective shear stress transfer from LDPE matrix to the particles of nascent PTFE powder and initiate their deformation and transformation into nanofibers. However, the compounding conditions were insufficient to obtain the PTFE nanofibers with very high deformation ratios, and as it was expected a plentitude of PTFE fibers underwent further deformation and were pulled out from LDPE matrix during fracture of the sample prepared for SEM investigation. Fig. 2 shows an exemplary SEM micrographs of the cryogenic fractured surface of Koattro/LDPE blend containing 3 wt. % of PTFE nanofibers. Only the ends of PTFE nanofibers are visible on the fractured surface of Koattro/LDPE blend indicating that PTFE nanofibers generated during preparation of the masterbatch underwent further deformation during dilution of the masterbatch in the TPE matrix and were strained strongly to the hardening stage where further deformation is difficult and requires higher stress. Very strong similar nanofibrillation of PTFE grains was evident also in Infuse/LDPE blends (results not shown here). 10
Fig. 2. SEM images of cryogenic fractured surfaces of Koattro/LDPE blend containing 3 wt.% of PTFE nanofibers (85/12/3).
It was impossible to determine the diameter and length of generated PTFE nanofibers directly from SEM images on fractured surface of TPE/LDPE blends. For this reason the etching of TPE/LDPE was applied. Fig. 3 shows SEM image of the free surface of TPE/LDPE blend with 3 wt.% of PTFE after 72 hours permanganic etching of a layer of Koattro/LDPE matrix as well as size distribution diagram of PTFE nanofibers.
11
Fig. 3. SEM image of free surface of Koattro/LDPE blend containing 3 wt.% of PTFE nanofibers (85/12/3) after 72 hours etching with size distribution diagram of PTFE nanofibers.
Detailed analysis of SEM images revealed that thicker PTFE fibers (i.e. 220-260 nm thick) are located in bundles in comparison with thinner ones (i.e. 100-140 nm). The length of PTFE nanofibers ranging from 30 µm up to few hundreds of µm. During mixing, the grains of nascent PTFE were deformed to very high deformation ratios. Simultaneously, generated nanofibers underwent entangling process. Entanglement knots of PTFE nanofibers work as local constrains for further deformation. However, the fibers which were split out from the bundle and underwent further deformation resulting in much thinner nanofibers in comparison with those present in the bundles. WAXS diffraction profiles of samples of neat Koattro, LDPE and their blends are presented in Fig. 4. It is indicated that Koattro crystallizes in the hexagonal crystallographic units of form I characteristic for poly(1-butene) [47].
Fig. 4. WAXS profiles of ethylene-butene copolymer Koattro, LDPE, Koattro/LDPE blend (60/40) and Koattro/LDPE blend containing 8 wt.% of PTFE nanofibers (60/32/8). Main diffraction peaks are marked on the plots. Diffraction peak from (107) planes of form IV of PTFE is also indicated. 12
This crystallographic form reflects x-ray at 2 = 10.3° and 17.8° assigned to the (110) and (300) peaks while at 2 = 20.7° there are two overlapped peaks from (220) and (211) planes. No signs of crystallization of ethylene component of Koattro is detected in Fig. 4. The diffraction profiles for 60/40 blend indicate that 1-butene and ethylene elements form separate crystals and they do not co-crystallize in the blends. Propably they will crystallize separately in different temperature ranges. The broad diffraction peak for 60/40 blend centered at 21.1o is composed of three reflections: (110) of ethylene orthorhombic form, and (220) and (211) of 1-butene form I. Amorphous hallo is visible on all diffraction profiles for Koattro/LDPE blends as a very broad and low diffraction because it is composed with ethylene and 1-butene amorphous phases with peaks centered at largely different 2 angles. When PTFE nanofibers are present in the blend (60/32/8) the diffraction peak (107) of form IV of PTFE is superimposed on (300) peak of 1-butene form I crystals. The addition of PTFE nanofibers does not influence significantly either crystal sizes of ethylene and 1-butene components (no significant peak broadening of respective X-ray peaks) or degree of crystallinity.
13
Fig. 5. X-ray diffraction profiles of ethylene-octene copolymer Infuse, LDPE, Infuse/LDPE blend (60/40) and Infuse/LDPE blend containing 8 wt. % of PTFE nanofibers (60/32/8). Main diffraction peaks are marked on the plots. Diffraction peak from (107) planes of form IV of PTFE is also indicated.
WAXS diffraction profiles of samples of Infuse and LDPE materials are shown in Fig. 5. It is seen that all of them form orthorhombic polyethylene type crystals, possibly co-crystallizing in the blends. The addition of PTFE nanofibers does not influence significantly either crystal sizes (no significant peak broadening) or degree of crystallinity. Amorphous hallo is visible on all diffraction profiles with a peak at around 2 = 19.5o. If PTFE nanofibers are present in the Infuse/LDPE blend (60/32/8) the amorphous hallo is superimposed with the diffraction peak (107) of form IV of PTFE. 3.2 Thermal properties of TPE/LDPE blends Thermal properties of TPE/LDPE blends and neat polymers were examined. Typical DSC heating and cooling thermograms for Koattro- and Infuse-based materials are shown in Fig. 6 and Fig. 7, respectively. Two endothermic peaks during melting and two exothermic peaks on crystallization were observed for all Koattro/LDPE blends indicating the formation of separate crystals. This observation is expected taking into account the conclusions, based on x-ray diffraction presented in Fig. 4, of separate crystallization of 1-butene crystals of Koattro and ethylene crystals of LDPE. The low temperature peaks are attributed to melting and crystallization of Koattro 1-butene crystals, whereas the higher temperature peaks are associated with melting and crystallization of LDPE crystals, respectively. Addition of 15 wt.% of LDPE in to Koattro matrix caused the increase in the onset of crystallization temperature by 40
o
C. However no significant changes were observed in
Koattro/LDPE blends containing PTFE nanofibers. 14
Fig. 6. DSC thermograms of neat Koattro, neat LDPE and Koattro/LDPE blends with PTFE nanofibers. The thermograms have been shifted vertically for better visualization.
Fig. 7. DSC thermograms of neat Infuse, neat LDPE and Infuse/LDPE blends with PTFE nanofibers. The thermograms have been shifted for better visualization.
Different thermal behavior was observed for Infuse-based materials (see Fig. 7). A single endothermic peak during melting and single exothermic peak during crystallization was evident for all Infuse/LDPE blends. It appeared, therefore, that cocrystallization of Infuse and LDPE took place, which is known for other polyolefin
15
blends such as LDPE with HDPE [48]. This observation is also supported by x-ray diffraction studies presented in Figure 5 of co-crystallization of ethylene elements of LDPE and Infuse. A gradual increase of melting temperature by 6.7
o
C and
crystallization temperature by 10.2 oC with the increasing content of LDPE was observed. Also, the onset of crystallization has been shifted to higher temperatures with increasing amount of LDPE. The temperature increase is caused by longer unperturbed ethylene sequences in LDPE than in Infuse ethylene copolymer. In case of Infuse/LDPE blend containing PTFE nanofibers, two melting peaks and broad crystallization peak with a small shoulder were observed (see DSC thermograms for 60/32/8 composition on Fig. 7). This is an result of PTFE nanofibers ability to accelerate the crystallization of LDPE [38,39].
3.4 Rheological properties The effect of PTFE nanofibers on viscoelastic behavior of TPE/LDPE blends was examined by means of small-amplitude oscillatory shear. Since microstructure of a molten polymeric materials undergoing shear is not significantly deformed, the oscillatory shear measurements is a useful tool to explore the effect of polymeric materials structure on their moduli and viscosity [49]. Exemplary curves of storage modulus, G’, and loss modulus, G’’, as a function of angular frequency, ω, for neat Infuse and Infuse/LDPE blends with 3 wt.% of PTFE nanofibers are presented in Fig. 8.
16
Fig. 8. Storage modulus, G’, and loss modulus, G’’, as a function of angular frequency, ω, for neat Infuse, Infuse/LDPE blend, and Infuse/LDPE blend containing 3 wt.% of PTFE nanofibers. For neat Infuse, both moduli, G’ and G’’, decreases with decreasing ω, and in low frequency range G’’ was larger than G’ demonstrating the viscous nature of Infuse (note log-log scale in Fig. 8). Similar behavior was observed for the Infuse blend with 15 wt.% of LDPE, while Infuse/LDPE blend containing 3 wt.% of PTFE nanofibers showed significantly different viscoelastic response as compared to that of neat Infuse and Infuse/LDPE blend. It appeared that PTFE nanofibers caused the increase of both moduli in the low frequency range but the increase of G’ was more pronounced than G’’. It can be also seen that the G’ curve for 85/12/3 blend tended 17
to reach a low-frequency plateau (solid-like response). We have shown previously that such tendency is attributed to the presence of physical entangled network of solid PTFE nanofibers in a polymer matrix [50]. This can be accurately seen in Fig. 9, presenting the loss factor, tanδ (=G’’/G’) for Infuse-based materials as a function of angular frequency, ω. Both for the neat Infuse and its blend with 15 wt.% of LDPE, the tanδ increases with decreasing ω reaching 3.6 and 3.2 at 0.1 rad s -1, respectively. In the case of 85/12/3 blend, a broad peak on the tan δ curve appeared in the low ω range as a result of slow relaxation of Infuse chains exerted by entangled network of PTFE nanofibers.
Fig. 9. Loss factor, tanδ, as a function of angular frequency, ω, for neat Infuse, Infuse/LDPE 85/15 blend and 85/12/3 blend containing 3 wt.% of PTFE nanofibers.
Shear
experiments
are
not
sufficient
for
comprehensive
rheological
characterization of polymeric materials. This is particularly true for all the processing operations where the extension of a polymer melt has to be considered. Uniaxial extension of molten polymeric materials is often used to predict their rheological behavior outside the regime of linear viscoelasticity (LVE) which is very sensitive to structural differences within the materials. Fig. 10 presents a time-dependence of 18
tensile stress growth coefficient,
, for Infuse-based materials recorded during
uniaxial extension. It can be seen that for neat Infuse, the tensile stress growth coefficient,
, gradually increases with time, which is a consequence of
stretching of polymer chains, and ultimately reaches plateau when the chains are nearly fully stretched [51].
Fig. 10. The tensile stress growth coefficient, as a function of time, t, for neat Infuse, Infuse/LDPE blends and Infuse/LDPE blends with 3 and 5 wt.% of PTFE nanofibers, recorded during uniaxial extension at 130 oC and at Hencky strain rate, , of 5 s-1. Solid line represents the threefold of the shear stress growth coefficient, for neat Infuse measured at the shear rate, , of 0.01 s-1. The
curve superimposed on the shear stress growth coefficient,
curve multiplied by a factor of 3 confirming that neat Infuse exhibited LVE response at the applied Hencky strain rate, . It is known that linear polymers fulfill Trouton’s law, i.e. the ratio of
to
equals to 3 in the linear regime of
deformation [52], and therefore they exhibit no strain hardening themselves, i.e. no significant increase of
above the LVE region can be observed during uniaxial
extensional deformation. Pronounced strain hardening can be induced by
19
introduction of long-chain branches into linear polymer [53-56]. Despite the fact that Infuse is ethylene-octene copolymer with a small amount of short-chain branches, no deviation of
from the LVE response was observed within the experimental
window. Such behavior is similar to that of linear low density polyethylene, LLDPE, which show the strain hardening only at low strain rates, e.g. below 0.1 s-1, whereas no strain hardening can be observed at high strain rates [57,58]. The rheological behavior similar to that of neat Infuse was observed for the Infuse/LDPE blend containing 15 wt. % of LDPE. Neat LDPE exhibit strain hardening itself due to the presence of long-chain branches, which makes a local obstruction for the chain stretching in the direction of extensional deformation [53,59,60]. However, the amount of LDPE added into Infuse matrix was insufficient to induce the strain hardening. In contrast to neat Infuse and 85/15 blend, the presence of 3 wt. % of PTFE nanofibers in the form of entangled network in the Infuse/LDPE blend 85/12/3 caused, after a certain period of time of the uniaxial extensional deformation, a rapid increase of
above the LVE region. The entangled network of PTFE
nanofibers was deformed during uniaxial extension, but entanglement knots, resisted to stretching, caused strain hardening of the material. Further enhancement of the strain hardening can be obtained by either increasing the content of PTFE nanofibers or LDPE, however, the upward deviation of
from LVE response is much more
pronounced for the 75/20/5 blend containing 5 wt.% of PTFE nanofibers as compared to that of the 75/25 and 60/40 blends. Similar rheological behavior during extensional deformation was observed for Koattro-based materials (results not shown here). To estimate the degree of strain hardening effect, the time-dependent strain hardening coefficient,
, defined as a ratio of
20
to
, was
determined. Fig. 11 presents the bar plots of the maximum strain hardening coefficients,
, i.e. the values achieved at maximum Hencky strain for all
studied materials.
Fig. 11. The maximum strain hardening coefficients, , for neat TPEs, TPE/LDPE blends and the blends with PTFE nanofibers.
It is known that molten LDPE exhibit the strain hardening phenomenon under extensional flow due to presence of the long chain branching [61]. So, the all TPE/LDPE blends shows higher values of The
compared to neat TPE matrices.
increases with increasing content of LDPE in the blend reaching 2,5
for Koattro/LDPE 60/40 blend. It appeared that the entangled network of PTFE nanofibers caused further improvement of the strain hardening phenomenon in all studied TPE/LDPE blends. However, it seems that the magnitude of
depended on the type of the TPE
matrix used. The strongest effect of PTFE nanofibers on the strain hardening was observed for the Infuse-based blends, where 5 wt.% of PTFE nanofibers were sufficient to cause a four and sixfold increase of Infuse/LDPE 75/25 blend and neat Infuse.
21
as compared to that of
Mechanical properties Mechanical properties of selected materials were examined by tensile drawing. Tensile parameters, i.e. Yield stress, σs, stress at break, σb, elongation at break, εb, and modulus of elasticity, E, of neat TPEs, their blends with LDPE and PTFE nanofibers are presented in Fig. 12 and Fig. 13.
Fig.12 Selected tensile parameters and modulus of elasticity for neat Koattro, its blends with LDPE and PTFE nanofibers. All Koattro/LDPE blends exhibit higher values of σs and E compared to neat TPE matrix. The higher the content of LDPE in Koattro/LDPE blends the higher the σs and E were obtained. For Koattro/LDPE 60/40 blend, the σs and E are about three times higher in comparison with neat Koattro. It can also be seen that σb drastically decreased for blends containing 25 wt.% LDPE or more, as well as εb which apparently results from worse compatibility of the blend containing higher content of LDPE. Presence of PTFE nanofibers in Koattro/LDPE blends resulted in further 22
increase of σs and E, whereas σb and εb decrease. The highest value of E, i.e. was obtained for Koattro/LDPE blend containing 8 wt.% of PTFE nanofibers, which is about three and a half higher than that for neat Koattro.
Fig.13 Selected tensile parameters and modulus of elasticity for neat Infuse, its blends with LDPE and PTFE nanofibers.
In case of Infuse-based materials, blending with LDPE also improved mechanical properties by increasing σ s and E (see Fig. 13), however better compatibility between Infuse and LDPE resulted in no decrease in elongation. It need to be mentioned that the specimens for Infuse/LDPE blends having LDPE content up to 40 wt. % did not break down during tensile drawing conditions, i.e. ε b obtained was above 2500 %. Modification of the blend system through entangled network of PTFE nanofibers enabled further improvement of materials stiffness. The best results were obtained for Infuse/LDPE blend containing 8 wt.% of PTFE nanofibers, where σs increased by eight times and E by twelve and a half time compared to neat Infuse matrix as well as two times higher compared to reference Infuse/LDPE 60/40 blend. However, the material containing 8 wt.% of PTFE nanofibers drastically lost its drawability, εb of only 300 % was obtained. It is seen that the TPE/LDPE blends containing PTFE nanofibers were significantly stiffer and slightly stronger compared to two component blend systems 23
but brittle with lower elongation at break, which decreased with increasing amount of PTFE nanofibers. Similar behavior was observed in conventional composite of polyamide 6, PA6, with glass fibers, GF, which exhibit increased modulus of elasticity and tensile strength whereas elongation at break drastically decreased [62].
4. Conclusions In this paper the effect of in situ generated PTFE nanofibers on the rheological and mechanical properties of TPE/LDPE blends was examined. Formation of PTFE nanofibers has been observed in all studied TPE/LDPE blends, provided that sufficient shear stress was imposed on the molten materials during compounding. The entangled network of PTFE nanofibers drastically improved the rheological properties of studied TPE/LDPE blends. The strain hardening was found to depend on concentration of the PTFE nanofibers and increased with increasing content of PTFE. The most pronounced effect was observed for the Infuse/LDPE blend containing 5 wt.% of PTFE nanofibers (75/20/5), where the maximum strain hardening coefficient was three times in comparison with conventional TPE/LDPE blend (60/40). Also, PTFE nanofibers improved the mechanical properties of TPE/LDPE blends. It appeared the stiffness and drawability depend on the content of PTFE nanofibers. Higher the concentration of PTFE nanofibers in TPE/LDPE blend the higher the modulus of elasticity and lower elongation at break. The modulus of elasticity of the material based on Infuse/LDPE blend containing 8 wt.% of PTFE nanofibers (60/32/8), was two times higher in comparison with 60/40 blend.
24
Acknowledgments The project of National Science Centre, Poland (DEC-2012/04/A/ST5/00606) is gratefully acknowledged for the partial financial support of the studies. Also, the project of National Centre for Research and Development, Poland (INNOTECHK3/IN3/24/227425/NCBR/14) and the statutory fund of the Centre of Molecular and Macromolecular Studies, Polish Academy of Sciences, are acknowledged.
References [1] Holden G, Understanding thermoplastic elastomers, Hanser, Munich (2000). [2] Spontak RJ and Patel NP., Curr. Opin. Colloid Interface Sci. 5:334-341 (2000). [3] Schmaltz H, Abetz V and Lange R, Compos. Sci. Technol. 63:1179-1186 (2003). [4] Coran AY, Das B and Patel RP, US Patent 4130535 (1978). [5] Scheirs J and Kaminsky W, Metallocene-Based Polyolefins. Preparation, Properties and Technology. Vol. 2. John Wiley & Sons, New York (2000). [6] Amin S and Amin M, Rev. Adv. Mater. Sci. 29:15-30 (2011). [7] Jancar J and Dibenedetto AT, J. Mater. Sci. 30:1601-1608 (1995). [8] Katbab AA, Nazockdast H, and Bazgir S, J. Appl. Polym. Sci. 75:1127-1137 (2000). [9] Bazgir S, Katbab AA, and Nazockdast H, J. Appl. Polym. Sci. 92:2000-2007 (2004). [10] Karakaya N, Ersoy OG, Oral MA, Gonul T, and Deniz V, Polym. Eng. Sci. 50:677-688 (2010). [11] Sreekanth MS, Joseph S, Mhaske ST, Mahanwar PA, and Bambole VA, J. Thermoplast. Compos. Mater. 24:317-331 (2011). [12] Bajsic EG, Rek V, and Pavic BO, J. Elastom. Plast. 45:501-522 (2012). [13] Shonaike GO, and Matsuo T, Compos. Struct. 32:445-451 (1995). [14] Ibarra L, and Panos D, Polym. Inter. 43:251-259 (1997). [15] Anuar H, and Zuraida A, Compos. Part B Eng. 42:462-465 (2011). [16] Saikrasun S, Amornsakchai T, Sirisinha C, Meesiri W, and Bualek-Limcharoen S, Polymer 40:6437-6442 (1999). [17] Vajrasthira C, Amornsakchai T, and Bualek-Limcharoen S, J. Appl. Polym. Sci. 87:1059-1067 (2003).
25
[18] Akhtar S, De PP, and De S, J. Appl. Polym. Sci. 32:5123-46 (1986). [19] Austin JR, and Kontopoulou M, Polym. Eng. Sci. 46:1491-1501 (2006). [20] Liu Y, and Kontopoulou M, J. Vinyl Addit. Technol. 13:147-150 (2007). [21] Li C, Deng H, Wang K, Zhang Q, Chen F, and Fu Q, J. Appl. Polym. Sci. 121:2104-2112 (2011). ]22] Coffey AB, and O’Bradaigh CM, J. Mater. Sci. 42:8053-8061 (2007). [23] Verhoogt H, Langelaan C, VanDam J, and DeBoer AP, Polym. Eng. Sci. 33:754-763 (1993). [24] Verhoogt H, Willems CRJ, VanDam J, and DeBoer AP, Polym. Eng. Sci. 34:453-60 (1994). [25] Machiels AGC, Denys KFJ, VanDam J, and DeBoer AP, Polym. Eng. Sci. 36:2451-2466 (1996). [26] Saikrasun S, Bualek-Limcharoen S, Kohjiya S, and Urayama K, J. Appl. Polym. Sci. 89:2676-2685 (2003). [27] Saikrasun S, Bualek-Limcharoen S, Kohjiya S, and Urayama K, J. Polym. Sci. Part B Polym. Phys. 43:135-144 (2005). [28] Saikrasun S, and Amornsakchai T, J. Appl. Polym. Sci. 103:917-927 (2007). [29] Shin BY, and Chung IJ, Polym. J. 21:851-861 (1989). [30 ] Fakirov S, Evstatiev M, and Petrovich S, Macromolecules 26:5219-5226 (1993). [31] Evstatiev M, Schultz JM, Petrovich S, Georgiev G, Fakirov S, and Friedrich K, J. Appl. Polym. Sci. 67:723-737 (1998). [32] Friedrich K, Evstatiev M, Fakirov S, Evstatiev O, Ishii M, and Harrass M, Compos. Sci. Technol. 65:107-116 (2005). [33] Fakirov S, Bhattacharyya D, Shields RJ, Colloid Surf. A Physicochem. Eng. Asp. 313-314:2-8 (2008). [34] Fakirov S, Compos. Sci. Technol. 89:211-225 (2013). [35] Jurczuk K, Galeski A, and Piorkowska E, Polymer 54:4617-4628 (2013). [36] Krajenta J, Pawlak A, and Galeski A, Polimery 60:64-66 (2015). [37] Antonioli D, Sparnacci K, Laus M, Boarino L, and Righetti MC, J. Appl. Polym. Sci. 132:42401 (2015). [38] Masirek R, Piorkowska E, and Galeski A, Polish Patent PL209925. [39] Bernland K, and Smith P, J. Appl. Polym. Sci. 114:281-289 (2009). [40] Van der Meer DW, Milazzo D, Sanguineti A, and Vancso GJ, Polym. Eng. Sci. 45:458-468 (2005). [41] Ali MABM, Nobukawa S, and Yamaguchi M, Pure Appl. Chem. 83:1819-1830 (2011).
26
[42] Jurczuk K, Galeski A, and Piorkowska E, J. Rheol. 58:589-605 (2014). [43] Galeski A, Bartczak Z, and Kazimierczak T, Polymer 51:5780-5787 (2010). [44] Owens DK, and Wendt RC, J. Appl. Polym. Sci. 13:1741-1747 (1969). [45] Mark JE, Physical properties of polymers. Handbook. 2nd edition, Springer, New York (2007). [46] Dettre RH, and Johnson Jr. RE, J. Colloid Interf. Sci. 31:568-569 (1969). [47] Al-Hussein M, and Strobl G, Macromolecules 35:8515-8520 (2002)]. [48] Fonseca CA, and Harrison IR, Thermochimica Acta 313:37-42 (1998). [49] Larson RG, The structure and rheology of complex fluids. Oxford University Press, New York (1999). [50] Gupta RK, Polymer and composite rheology. 2nd edition, Marcel Dekker, Inc., New York (2000). [51] Marrucci G, and Ianniruberto G, Macromolecules 37:3934-3942 (2004). [52] Münstedt H, Rheol. Acta 14:1077-1088 (1975). [53] Laun HM, and Münstedt H, Rheol. Acta 17:415-425 (1978). [54] Ishizuka O, and Koyama K, Polymer 21:164-170 (1983). [55] Kurzbeck S, Oster F, Münstedt H, Nguyen TQ, and Gensler R, J. Rheol. 43:359–375 (1999). [56] Kasehagen LJ, and Macosko CW, J. Rheol. 42:1303–1327 (1998). [57] Münstedt H, KurzbeckS, and Egersdörfer L, Rheol. Acta 37:21-29 (1998). [58] Münstedt H, Kurzbeck S, and Stange J, Polym. Eng. Sci. 46:1190-1195 (2006). [59] Münstedt H, and Auhl D, J. Non-Newtonian Fluid Mech. 128:62–69 (2005). [60] Wagner MH, Bastian H, Hachmann P, Meissner J, Kurzbeck S, Münstedt H, and Langouche F, Rheol. Acta 39:97–109 (2000). [61] Gabriel C, and Mϋnstedt, J. Rheol. 47:619-630 (2003). [62] Thomason JL, Compos. Part A 39:1618-1624 (2008).
27
Graphical Abstract Thermoplastic elastomers reinforced with poly(tetrafluoroethylene) nanofibers. Kinga Jurczuk* and Andrzej Galeski
TPE/LDPE blends reinforced with PTFE nanofibers produced by solid-state deformation of PTFE crystals.
28
Higlights: TPE/LDPE blends reinforced with PTFE nanofibers; PTFE nanofibers generated in situ during compounding of crystalline PTFE grains with molten matrix based on TPE/LDPE blends; Strain hardening of molten TPE/LDPE blends improved by entangled network of PTFE nanofibers; Significant improvement of the mechanical properties TPE/LDPE blends containing PTFE nanofibers;
29