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fluoroelastomer thermoplastic elastomeric blends: Influence of interaction and morphology on physical properties
Accepted Manuscript Novel Nanostructured Polyamide 6/Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties Shib Shankar Banerjee, Anil K. Bhowmick PII:
S0032-3861(13)00943-9
DOI:
10.1016/j.polymer.2013.10.001
Reference:
JPOL 16532
To appear in:
Polymer
Received Date: 10 May 2013 Revised Date:
11 September 2013
Accepted Date: 1 October 2013
Please cite this article as: Banerjee SS, Bhowmick AK, Novel Nanostructured Polyamide 6/ Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties, Polymer (2013), doi: 10.1016/j.polymer.2013.10.001. 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.
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Graphical Abstract
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Novel Nanostructured Polyamide 6/Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties
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Shib Shankar Banerjee and Anil K. Bhowmick
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Novel Nanostructured Polyamide 6/Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties Shib Shankar Banerjee and Anil K. Bhowmick*
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Department of Material Science and Engineering, Indian Institute of Technology, Patna – 800013, India
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ABSTRACT: Novel Polyamide 6 (PA6)/Fluoroelastomer nanostructured thermoplastic elastomeric blends were developed in the present work. The influence of interaction between
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the components and morphology on physical properties of the blends was analyzed. Scanning Electron Microscopy and Atomic Force Microscopy studies, solubility and theoretical analysis of complex modulus clearly indicated that PA6 was the continuous matrix in which fluorocarbon elastomer was present in nanoscale. Low torque ratio (0.34) of rubber/plastic, high mixing speed and long mixing time had an important role in developing
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the nanostructured morphology of the blend. Tensile strength of the thermoplastic elastomer was about 39.0 MPa which was much higher than that reported earlier and showed significant improvement with increasing PA6 content. Large shifting of the glass transition temperature
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of the rubber and the plastic phases towards the lower temperature compared to those pristine polymers was also observed. The above properties were explained with the help of interaction
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between the components and morphology.
Key words: Polyamide 6, Fluoroelastomer, thermoplastic elastomer,
*Corresponding Author. Tel/Fax: +91 612 2277380. E-mail address: [email protected] (Anil K. Bhowmick)
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1. Introduction Thermoplastic elastomers (TPEs) have appeared as one of the most important classes of speciality materials having immense potential for industrial revolution. They have unique and
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extraordinary properties such as excellent physico-mechanical properties, chemical resistance, wider processing windows, ability to bond with multiple thermoplastics, light weight and excellent price-performance ratio, all of which make these an ideal candidate for
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enormous industrial applications [1-4]. Polyamide based thermoplastic elastomers are reported to have great importance because of their excellent strength, stiffness, low friction,
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high melting point, and chemical and wear resistance [5]. Thermoplastic elastomers based on polyamides, prepared by melt mixing process are expected to have excellent heat and oil resistance [6]. Several studies of different properties of blends of polyamide and rubbers are available in the literature. Crisenza et al. studied direct 3D visualization of the phaseseparated morphology of thermoplastic elastomers from chlorinated polyethylene/nylon
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terpolyamide blends [7]. Jha and Bhowmick investigated the reactive nylon-6/acrylate thermoplastic elastomeric blends and their interaction on mechanical and dynamic mechanical properties [8]. Oderkerk and Groeninckx reported the morphology development
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by reactive compatibilization of nylon-6/EPDM blends with a high amount of rubber fraction and extensively studied the effect of viscosity ratio of rubber/thermoplastic on the blend
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morphology [9].
Coran and Patel extensively investigated the thermoplastic elastomers from various rubberplastic blends and correlated the properties with morphology, surface energy mismatch and crystallinity of the matrix phase of the blends [10-11]. Few authors investigated the interaction between immiscible polymers through the addition of reactive components [1214]. Mechanical properties, dynamic response and interaction between components (both reactive and non-reactive systems) of rubber-plastic blends are main concern with the rubber
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modified plastics systems. Many researchers reported non-covalent bonding interaction in polymer blends (mainly H-bonding interaction) through infrared spectroscopic analysis [1518].
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Among different classes of thermoplastic elastomers, blends of rubbers and plastics are gaining importance because the required property can be easily monitored by varying the ratio of the constituents, viscosity of the components and also by incorporation of additives
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like compatibilizer, crosslinking agents, fillers, etc. Even though many polyamide based TPEs are commercially available, development of polyamide-fluoroelastomer TPEs has not
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been reported in literature. Also, polyamide shows interesting processing problems and its blends with rubber without compatibilizer exhibit difficulty in processing and poor mechanical characteristics.
The present investigation reports our observations on the blends of polyamide 6 and fluoroelastomer by melt-blending process. The influence of interaction and blend ratio on
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dynamic mechanical and mechanical properties was investigated. FTIR study on H-bonding interaction between polyamide 6 and fluoroelastomer was highlighted. The molecular weight of polyamide is one of the most important parameters governing the melt viscosity and hence
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processing characteristics, interfacial adhesion and morphology of the blend. The rubber particle size in rubber toughened plastics can be reduced by an order of magnitude by
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increasing the molecular weight of polyamide keeping other process conditions constant [1922]. Our objective is to develop a thermoplastic elastomeric blend with nanodimensional elastomeric particles dispersed in a plastic matrix. To the best of our knowledge, this is the first time that a nanostructured polyamide based thermoplastic elastomeric blends has been prepared by melt blending process. Furthermore, the effects of nanostructured morphology on the physical properties of the blends are reported.
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2. Experimental Section 2.1. Materials Polyamide 6 (tradename Akulon) in granular form (viscosity number 245cm3/g, density 1130
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kg/m3, melting point 220 °C) was supplied by DSM, Nederland. Viton A (a copolymer of vinylidene fluoride (VF2) and hexafluoropropylene (HFP), density 1810 kg m-3 at 25 °C, 66% F, Mooney Viscosity, ML
1+10
at 121°C = 20) abbreviated as FKM was procured from
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DuPont Dow Elastomers, Geneva, Switzerland. The formulation of the mixes is given in
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Table I. The molecular structure of polyamide 6 and FKM is given in Fig. 1
2.2. Preparation of Thermoplastic Elastomeric Blend Compositions PA6 granules were dried in vacuum oven at 100 °C for 24 h before blending. The blends were compounded in batches of 50 g polymer in a Haake Rhecord internal mixer with roller
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type rotor. Blending was allowed to proceed for 7 min at a temperature of 240 °C and 100 rpm rotor speed. After 2 min melting of PA6 in the internal mixer, FKM was added and mixed for another 5 min under the same conditions. The changes of mixing torque and energy
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with time were recorded for each composition. After complete mixing in the internal mixer, the resulting blends were quickly removed and passed through a two-roll mill having close
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nip-gap at room temperature to make sheets and were cut into small pieces. One blend composition {40 PA6/60FKM (w/w)} which showed best thermoplastic elastomeric behaviour was also processed in the following two ways: i. Similar procedure as above was followed except high viscosity PA6 was replaced by low viscosity PA6. ii. Mixing was allowed to proceed for 3.5 min at a temperature of 240 °C and 50 rpm rotor speed.
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The procedure followed after complete mixing was same as before.
2.3. Molding of Thermoplastics Elastomers
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Test specimens (1-1.5 mm thick) were prepared by means of micro-injection molding machine (Haake MiniJet II) at 260 °C cylinder temperature and 60 °C mold temperature. Injection pressure and holding pressure were 700 bar and 350 bar respectively. The injection
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and holding times were 5 sec each.
2.4.1. Solubility Measurement
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2.4. Characterization
Approximately 0.5 g (accurately weighed) of each blend, put inside a packet of Whatman-41 filter
paper, was immersed in 50 ml of methyl ethyl ketone (MEK) solvent at room
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temperature and the solvent was renewed every 24 h. After complete extraction as measured from constant residual weight, the samples were removed and dried to constant weight using a vacuum oven at 70 °C. The weight percent of entrapped FKM into polyamide matrix was
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calculated from the final and the initial weight of the rubber in the blends. Similar procedure for solubility in formic acid was followed to calculate the weight percent of unextracted
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polyamide.
2.4.2. IR Spectroscopy in Attenuated Total Reflectance (ATR) Mode IR spectra (in ATR mode) of the pristine polymers and the blends were taken on smooth film using a Perkin–Elmer (model 400) spectrophotometer with a resolution of 4 cm-1 and sixteen scans were averaged for each spectra.
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2.4.3. Microscopy 2.4.3.1.
Field Emission Scanning Electron Microscope
Morphology of various blends was recorded from FESEM (S4800 Hitachi microscope) at an
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acceleration voltage of 15.0 kV at a working distance of 6 mm. A thin layer of Platinum was sputter coated for 60 sec on the smooth sample surface to avoid charging on exposure to electron beam during FESEM analysis. Measurement of particle diameters (di) from the
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photomicrographs gave the rubber particle size distributions for different compositions of the
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blend. From these data, the number average diameter (< dn >) and weight average diameter
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() were determined using the following equations:
The inter-particle distance between dispersed phases play an important role on the mechanical properties of polymer blends. Wu [23] first proposed the concept of the average thickness of the ligament between two particles i.e. the inter-particle distance that decides whether a blend will be tough or brittle. This concept has been applied for several rubber-
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plastic blends and rubber-toughened plastics [24-25]. For the cubic array of spherical
is the volume fraction of rubber and
2.4.3.2.
Atomic Force Microscopy
is the rubber particle size.
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where,
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particles, the inter-particle distance can be obtained from the following equation.
Intermittent Contact Mode Atomic Force Microscopy, ACAFM (Agilent 5500 Scanning Probe Microscope) images were obtained for supporting the nanostructures of the blends. The
2.4.4. Mechanical Tests
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resonance frequency of the tip was 130-280 kHz and the force constant was 48 N/m.
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Tensile test was carried out in a Zwick Universal Testing Machine (UTM) model Z010 at room temperature at a test speed of 200 mm/min. The samples were prepared following ISO
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527-2-5A specification. The average of three tests is reported here. For tension set measurement, the samples were extended up to 50 % in the tensile direction at a rate of 50 mm/min and kept at that position for 10 min at room temperature. It was relaxed back to unstressed condition and the percentage change in dimension in tensile direction was measured after 24 h and reported as tension set. Shore D hardness of samples was obtained by DIN ISO 7619 test method at room temperature.
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2.4.5. Dynamical Mechanical Analysis The dynamic mechanical properties of the blends and individual polymers were measured using a DMA of TA Instruments (model Q800) in tension mode. All the samples (height 16
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mm, width 9.8 mm, thickness 1mm) were analyzed at a constant frequency of 1 Hz, at a heating rate of 2 °C/min and a strain amplitude of 30 µm over a temperature range of - 100
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°C to + 100 °C.
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2.4.6. Differential Scanning Calorimetry
Thermal analysis of the pristine polymers and the blends was carried out using a Perkin– Elmer Differential Scanning Calorimeter (DSC 8500) at a heating and cooling rate of 5 °C min-1 in an inert atmosphere (N2 atmosphere). Standard aluminum pans were used to analyze the sample. The experiment was conducted from – 70 °C to 240 °C. The data of second
Results and Discussion
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3.1. Mixing torque
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3.
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heating cycle was used to eliminate the thermal history.
Torque during polymer mixing is a complex combination of shear and elongational flow in actual mixing environment. Mixing torque behavior of the pristine polymers and the blends is shown in Fig. 2. The torque-time curves for all the blends in Fig. 2 gave two peaks. The first peak was attributed to the increase in mixing torque (i.e. viscosity) value due to resistance exerted on the rotor by the unmolten polymer followed by a decrease in viscosity due to the complete melting of polyamide 6. Upon the addition of FKM into polyamide, the torque value again increased which gave the second peak. A gradual decline in torque value for the
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blend was observed with the rise in stock temperature and both reach a steady-state value after a certain time of mixing. The uniform torque value after a specified time suggested good level of mixing of the two components. The extents of mixing and mixing time are important
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parameters for immiscible polymer blends, since final properties of the blends depend on these parameters. From Fig. 2, it was clear that the equilibrium mixing torque was maximum for PA6 and lowest for FKM. This was attributed to the high melt viscosity of the PA6
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3.1.1. Theoretical analysis of mixing torque
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compared to FKM under the specified mixing conditions.
In order to assess the interaction between the blend components, theoretical calculation of mixing torque was investigated. We have calculated the theoretical value of torque based on
Here,
and
) phases using the following equation
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volume fraction of plastic ( ) and rubber (
are the torque values of the plastic and the rubber respectively determined
from Haake rheomixer at 240 °C and 100 rpm rotor speed. Table 2 highlights the theoretical
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and experimental values of the torque, which vary as a function of rubber proportions. It was observed that the experimental value of the torque of PA6/FKM blends was much higher than
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the theoretical value in the entire composition range. Melt mixing of PA6 and FKM results in an increase in molecular weight due to interaction and therefore, the viscosity of the mixture, which raises the torque during mixing. It is interesting to note that positive deviation of torque increases with increase of the plastic content and maximum positive deviation of torque (~ 81 %)) is observed when polyamide content is 80 percent by weight in the blend. It indicates that interaction of FKM and PA6 increases with increasing the PA6 content. This phenomena is also observed from the solubility characteristics and discussed later.
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3.2. Solubility Characteristics In order to determine the extent of interaction between PA6 and FKM during melt blending
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process and the amount of weight percent of FKM entrapped into hard polyamide matrix, solubility measurements have been done using their respective solvents. The data are presented in Fig. 3. The amount of unextracted FKM by MEK solvent has been taken as the
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entrapped fluoroelastomer in the polyamide matrix and this is valid when FKM part is uncured in the blend systems. The blend composition of 80 PA6/20 FKM (w/w) showed
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higher amount of unextracted rubber compared to other compositions (Fig. 3). This result revealed that maximum amount of rubber encapsulation had taken place when rubber content was near 20 percent by weight in the blend. This result supports the findings from the mixing torque measurement. It may also be assumed that some of the FKM chains were immobilized close to the boundary between the two phases. As a consequence, these rubber particles were
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embedded in the boundary region (This phenomena is also revealed from dynamic mechanical analysis and discussed later). From the solubility in formic acid, it was observed that the structure of the blend was completely collapsed and the material in the form of
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powder was found, after removal of polyamide from the blends. This result gave a direct evidence that polyamide 6 formed the continuous matrix in the whole composition regime.
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From Fig. 3, it was clear that at lower weight percent of rubber, there was no unextracted polyamide but at high weight percent of rubber in the blend, unextracted polyamide was found. As polyamide is hard matrix phase in the blend and mainly non-covalent bonded interaction occurs between the components (confirmed by FT-IR analysis and discussed later), lower amount of dispersed phase was unable to tightly hold the hard matrix part in its solvent. At higher weight percent of the dispersed phase, however, some amount of unextracted polyamide was found.
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3.3. Infrared Spectroscopic Analysis Infrared spectroscopic technique is extensively used to identify specific interactions that
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occur in various polymer systems [26-27]. It was apparent from the mixing torque and solubility measurements that an appreciable amount of interaction occurred between PA6 and FKM. Preparation of melt-blending of polyamide with various functionalized polymers was
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reported by a few researchers [28-30]. The amine and carboxyl end groups of polyamide 6 have ability to bond with other reactive polymers having suitable functional groups at their
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chain. The fluoroelastomer used in this blend contains C-F, C-F2 and C-F3 groups in their chain. Hence, H-bonds with N-H group of polyamide 6 can be formed. The characteristic peak positions of the pristine polymers are assigned in Table 3. Polyamide 6 gave N-H stretching intensity at 3296 cm-1, but the N-H stretching intensity was shifted to 3290 cm-1 after blending (Fig. 4). Similarly, C-F and C-F3 stretching of pure FKM appeared at 1397 cmand 1353 cm-1 respectively. C-F and C-F3 combined stretching appeared at 1022 cm-1. In the
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blend systems, the intensities were shifted to 1375 cm-1 (C-F stretching), 1338 cm-1 (C-F3 stretching) and 1018 cm-1 (C-F and C-F3 stretching) respectively (Fig. 4). The reason for the
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shift might be due to interactions between the blend components and formation of H-bond between N-H group of PA6 and C-F containing groups of FKM. The probable scheme of H-
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bond formation is given in Fig. 5.
3.4. Microstructure Analysis The properties of multiphase polymeric blends strongly depend on interfacial chemistry and microstructure. The morphology of PA6/FKM blend at different compositions is highlighted in Fig. 6. The distribution of the rubber particle size for each composition is also given in the
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adjoining Fig. 6. For the distribution of particle size, each rubber particle was approximated as a circle and the diameter of the particle was measured manually in the SDM image analysis software (available with S4800 Hitachi microscope), because they could not
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identified by the threshold of black-white images. Table 4 reflects the blend morphology in
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terms of number average diameter dn, weight average diameter dw, particle size distribution
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and inter-particle distance of the dispersed phase. Interestingly, the dimension of the
dispersed FKM phase in polyamide matrix was found in the nanometric level in the whole
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composition regime of the blends. Weight average diameter of the dispersed phase was in
the range of 89-110 nm. The average particle size of the rubber increased with increase in
its weight percent in the blends and narrow particle size distribution was observed with a
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typical value of 1.05. As shown in Table 4, it is clear that inter-particle distance based on
weight average particle diameter increased as the dispersed phase volume fraction decreased. Maximum inter-particle distance (60.61 nm) was observed when the dispersed phase was
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near 20 percent by weight in the blends. It is interesting to note that continuous polyamide phase with dispersed fluorocarbon elastomer was observed for the composition range of 80
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PA6/20 FKM to 20 PA6/80 FKM (w/w). This is the first time that nanostructure has been reported for a thermoplastic elastomeric composition. The prominent factors that decide the morphology of blends are viscosity ratio (viscosity of disperse phase/viscosity of matrix), blend composition, elasticity, shear rate during mixing, interfacial modification and additives
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such as fillers, plasticizers, etc. It was reported that the melt viscosity differences between the components and composition ratios determined the morphology of the blends of isotactic polypropylene and ethylene-propylene rubbers for the same processing conditions [31]. The
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development of nano-dimensions of the elastomer phase in blend systems could be explained in the light of torque ratio of the components and very high melt-viscosity of the matrix
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component (both were measured by mixing torque behavior). It was reported that transient breakup process of the disperse phase during melt-mixing was more dependent on matrix viscosity than the viscosity ratio [32]. We have taken PA6 having different viscosity values. These were then used to prepare PA6/FKM blends to show the effect of viscosity ratio of the component on the properties of nanostructured blends. Here, we measured the morphology after 7 min of mixing followed by injection molding of the samples. AFM photograph of 40 PA6 (low viscosity)/60 FKM (w/w) is shown in Fig. 7d. The particle size of the rubber phase
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was found to be 0.30 µm. It was found that under the above specified condition, the torque value was 6.4 Nm for PA6, 0.8 Nm for low viscosity PA6 and 2.2 Nm for FKM (Fig. 2). The ratio of the torque values was 0.34 for FKM/ PA6 and 2.75 for FKM/ PA6 (low viscosity).
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This indicated that high melt-viscosity of PA6 compared to FKM under the specified mixing conditions, was one of the crucial parameters for the development of the nanostructured blends. The generation of increased stress during mixing and molding operations, due to
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highly viscous PA6 gave extensive break-up of the dispersed phase and reduced its particle size compared to low viscosity PA6. It is clear from AFM phase image (Fig. 7). However,
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when mixing speed and time were reduced to their half value compared to the specified conditions for 40 PA6 (high viscosity)/60 FKM (w/w), the dimension of the dispersed phase was 0.22 µm (Fig. 7c). Favis and Chalifoux explained that minimum particle size of the dispersed phase (0.5 µm) for PP/PC blends was obtained when torque ratio was approximately 0.25 [33]. So, the torque ratio of the components and viscosity of each matrix
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play an important role in the achievement of nanostructure blends. Increasing the concentration of the dispersed phase increases the number of particles and hence particleparticle interaction and subsequent coalescence. As a result, the particle size increases. If we
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consider the collision of two identical dispersed particles in shear flow during melt mixing in an internal mixer, coalescence may occur in the following way: the dispersed particles
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deform upon collision and then matrix material is expelled from the regime between them. The basic influencing parameters which determine the extent of coalescence are the volume fraction of the disperse phase, viscosity of each phase, radius of particle and the interfacial tension which is the major driving force for coalescence [26]. In order to support the nanostructured morphology of PA6/FKM blends, further investigation was done by Atomic Force Microscopy. Fig. 7a-b reflects the phase morphology of 40 PA6/60 FKM and 60 PA6/40 FKM (w/w) blends. The light yellow colour in the phase image
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corresponds to the soft rubber particles, which are dispersed in the hard polyamide matrix. The dimension of this dispersed phase was also found in nanometer range (60-120 nm), which supports and coincides with the findings from FESEM analysis.
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In order to further substantiate the dimension of the dispersed phase, FKM phase was preferentially etched by MEK. Fig. 8 exhibits the dark holes dispersed throughout the polyamide matrix for the entire blend compositions. The average diameter of the dispersed
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holes was calculated by measuring the diameter of about 50 randomly selected small holes for each composition of the blends. We have considered both minor and major axis to
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measure the diameter of the dispersed hole in Fig. 8. It is clear from Fig. 9 that the weight average diameter of the dispersed particle is almost the same as the average diameter of dispersed hole. However, the small holes are observed to merge in Fig. 8.
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3.5. Mechanical Properties
The basic differences among elastomer, plastic and thermoplastic elastomer are directly reflected in the mechanical properties of the blends. Typical tensile stress vs. tensile strain
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curves for the blend compositions are presented in Fig. 10. Table 5 shows the mechanical properties for the same and the pristine polymers. The stress-strain curve for 20 PA6/80 FKM
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(w/w) blend exhibited low modulus and considerable elongation, typical behavior for unvulcanized or undervulcanized rubber. Conversely, the curve for 80 PA6/20 FKM (w/w) blend composition showed high modulus with a narrow yield point and a consistent plastic deformation plateau, which reflected the typical behavior for a thermoplastic material. Finally, the curves for the compositions of 40 PA6/60 FKM (w/w) and 60 PA6/40 FKM (w/w) displayed typical shape of thermoplastic elastomers. The curves were marked by the absence of an yield point before failure (Fig. 10). As mentioned in Table 5, tensile strength,
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Young’s modulus, tension set, toughness and hardness increased with increase of PA6 content in the blends. Interestingly, the blend composition of 80 PA6/20 FKM (w/w) exhibited very high toughness value compared to other compositions in the blends (Table 5).
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So, this particular composition may behave as a rubber-toughened plastic. Surprisingly, for the PA6/FKM blends, a very high value of tensile strength (27.9-39.0 MPa) was observed, when they were compared with polyamide systems reported in the literature [7-8] probably
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due to interaction and injection molding adopted here. Also, significant increase of tensile strength was found with an increase of polyamide content in the blends, which was not
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reported earlier for chlorinated polyethylene/nylon terpolymer and nylon/acrylate systems [78]. This significant improvement of tensile strength and very high value of toughness in the blends could be explained on the basis of nanostructured morphology in the PA6/FKM blends. In order to support this, we prepared blends of 40 PA6/60 FKM (w/w) in the following two ways: (1) by changing the mixing conditions (dimension of dispersed phase
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was 0.22 µm), and (2) by using low viscous polyamide under the specified conditions (dimension of dispersed phase was 0.30 µm). As shown in Table 5 and Figure 10, tensile strength and toughness are found to be low when dimension of elastomer was in micron
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range. It was proven that PA6 and FKM have H-bonding interaction. The higher the surface area, higher will be interaction. Nanostructure of elastomers generates higher surface area
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than the microstructured blends, which results in higher tensile strength and toughness.
3.6. Dynamic Mechanical Analysis Fig. 11 represents the temperature dependence of tan δ of pristine polymers and their blends over a temperature range of - 100 to 100 °C. In the above temperature range, PA6 mainly revealed two main transitions, α and β, which appeared at 64 °C and - 63 °C, respectively.
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The α-transition, designated as its Tg (64 °C), was attributed to motion within the amorphous phase and mainly dependent on crystallinity of the materials. Conversely, the β–transition was due to carbonyl groups of polyamide forming H-bond with absorbed water [34]. In the
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case of FKM, main transition was found at - 4 °C and ascribed to the glassy to rubbery transition or Tg. The blend composition, 40 PA6/60 FKM (w/w) showed three main damping peaks- one at - 13.5 °C corresponding to the glassy to rubbery transition of FKM, another
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broad peak at 50.5 °C due to Tg of PA6, and the third broad peak at - 61.5 °C arising due to PA6. Interestingly, the peak at - 4 °C was shifted towards the lower temperature and the β
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transition temperature for polyamide has a tendency to move towards the Tg of rubber. The Tg of the plastic was also shifted towards the lower temperature. However, the shift in the peak positions of the blend revealed interaction and small amount of miscibility between the phase components.
The DMA results of PA6/FKM blends are presented in a tabular form (Table 6) in terms of
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(i) glass transition temperature of the rubber phase, Tgr, (ii) glass transition temperature of the plastic phase, Tgp, (iii) (tanδ)max of the rubber phase with respect to weight percent of rubber in the blends and (iv) storage modulus, E′ at 30 °C of the blends. (tanδ)max and Tgr of rubber
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phase decreased with increasing the plastic content in the blends, which could be explained on the basis of interaction between FKM and PA6 phases. It was also found that Tgp of the
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plastic phase decreased with increasing the rubber content in the blends. Storage modulus, however, showed an increasing trend with an increase of the plastic content in the blends. McCrum reported that the amplitude of dynamic transition of a component of a composite was directly attributed to the relative quantity of the component itself [35], which was further confirmed from our previous work. [8].
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Let us visualize each rubber domain as a combination of two separate phases- bulk rubber ‘active’ in the dynamic transition and immobilized rubber particles close to the boundary between two phases with restricted mobility. Thus, it might be said that with increasing the
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plastic content in the blends, the relative quantity of the bulk rubber ‘active’ for the dynamic transition decreased and immobilization of the rubber particle to the phase boundary increased, thus, resulting in a reduction of (tanδ)max value of the rubber phase. The second
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result, i.e. the decrease in Tg of the rubber phase, can be ascribed in the light of thermal stresses built-up in each of the rubber particles, as the sample was cooled from the melt
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during mixing and molding operations. A state of triaxial tension was expected due to greater thermal contraction of rubber compared to the glassy matrix [36-37]. This thermal stress could only be developed when there was sufficient adhesion between the two phases. As a consequence of this, the free volume and chain mobility of the rubber increased, resulting in a decrease of the glass transition temperature. Such type of contrasting feature was also found
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for nylon/acrylate blends [8]. Interestingly, we observed that Tg of FKM was shifted by 9.5 °C in the PA6/FKM blends, whereas there was a shift of only 1.5 °C for the nylon/acrylate blends having the same composition (both cases in the negative direction compared to their
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pristine glass transition temperature). As the fluoroelastomer has maximum tendency to shrink (3.17 %) compared to other elastomers [38], its free volume increment is much more
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prominent for fluorocarbon rubber based blends compared to acrylate rubber counterpart. To the author’s knowledge, this large shift of Tg of rubber has not been reported before for rubber-plastic blends.
The third result, i.e. the decrease in Tg of the plastic phase, could be explained on the basis of enhancing segmental motion of the glassy polymer chains when attached to a more mobile component i.e., lower Tg. As PA6 was mixed with FKM chains which were rubbery and
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mobile in nature, the flexibility and mobility of the polyamide chains are increased due to
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mutual interaction, leading to a decrease in its Tg value.
3.7. Differential Scanning Calorimetric Studies
In order to assess the influence of thermal stresses built-up in the FKM particles during mixing and molding with PA6, DSC experiment was done. We have taken the data from
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second heating to eliminate the thermal history. Fig 12 shows the DSC thermograms of
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pristine polymers and blend compositions.
PA6 showed melting a endotherm at 220 °C and glass transition at 47 °C. FKM revealed glass transition at - 25.5 °C. The blend composition, 40 PA6/60 FKM (w/w) showed base line shift in three regions- one at - 30.5 °C corresponding to the glass transition of FKM, another at 42.5 °C due to Tg of PA6, and the third at 208 °C arising due to melting of PA6. Table 6
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summarized the thermal analysis data of PA6, FKM and blend systems. It was interesting to note that the glass transition temperature of the rubber phase of the blend was lower than that of the control FKM and crystalline melting point of PA6 decreased substantially with the
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addition of FKM. This phenomena is in line with the findings from dynamic mechanical analysis confirming that thermal stresses built-up in the FKM particles are permanent and
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responsible for lowering the glass transition of the rubber in the blend systems.
3.8. Theoretical Analysis of Complex Modulus In order to assess the behaviour of the two-phase blend from the component property data, the existing theoretical Kerner model was used [39]. The complex moduli of the blends at 30 °C (a temperature between the two Tg values of the components) as well as the theoretical values of the modulus of the blends based on Kerner’s model are plotted in Fig. 13. The
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Kerner’s equation for a binary blend where one component is dispersed in a matrix material
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is
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where
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E* is the complex modulus of the blend, Em* is the complex modulus of the matrix, Ei* is the
complex modulus of the dispersed phase. ν is the Poisson ratio of the blend, and νi and νm are
the Poisson ratio of the dispersed and the matrix phases respectively. The Poisson ratio of the blend, which varies with temperature, has been calculated from the following equation [40]
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It was clear from Fig. 13 that the experimental complex moduli for the PA6/FKM blends, especially at higher plastic content, are very close and almost coincide with those obtained from Kerner’s rigid matrix–soft filler model, suggesting the formation of polyamide as the continuous matrix. This result from the theoretical analysis of complex modulus is in good
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accord with the findings from the morphology and the solubility measurement of the blends
4. Conclusions
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that polyamide is the continuous matrix in the blends systems.
Nanostructured blends of PA6 and FKM were developed and fully characterized. The
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following conclusions are drawn:
(1) The mixing torque and IR spectroscopic analysis indicated interaction between the polymers.
polyamide matrix.
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(2) The solubility measurement revealed that rubber particles are encapsulated into hard
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(3) FESEM, AFM, solubility and theoretical analysis of complex modulus provided strong evidence for continuous polyamide phase in the blends. (4) The low torque ratio (0.34) of rubber/plastic, high mixing speed and long mixing time played an important role in developing the nanostructured morphology. (5) Very high value of tensile strength of the blends was observed when they were compared with polyamide systems reported in literature and high value of toughness was also observed when FKM was near 20 weight percent in the blend.
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(6) In the dynamic mechanical analysis, it was observed that the (tanδ)max as well as Tg of rubber phase decreased with increase of plastic content as a layer of restricted chain mobility was formed near the phase boundary and a state of triaxial tension was generated during
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cooling of the samples. (7) Micro-structural analysis showed the nanoheterogeneity of the blend which is reported first time for thermoplastic elastomer blends and also narrow particle size distribution of the
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dispersed phase.
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References
[1] Bhowmick AK, Stephens HL. Handbook of Elastomers, 2nd ed. New York, Marcel
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Dekker; 2001.
[2] Holden G, Kricheldorf HR, Quirk RP. Thermoplastic Elastomers. Munich, Hanser Verlag; 2004.
[3] Chen Y, Kushner AM, Williams GA, Guan Z. Nat Chem 2012;4:467-72. [4] Amin S, Amin M. Rev Adv Mater Sci 2011;29:15-30. [5] Wang Z, Zhang X, Zhang Y, Zhang Y, Zhou W. J Appl Polym Sci 2003;87:205762.
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[6] Coran AY, Patel R. Rubber Chem Technol 1980;53:781-94. [7] Crisenza T, Butt HJ, Koynov K, Simonutti R. Macromol Rapid Commun 2012;33:114-9.
[9] Oderkerk J, Groeninckx G. Polymer 2002;43:2219-28.
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[8] Jha A, Bhowmick AK. Rubb Chem Technol 1997;70:798-814.
[10] Coran AY, Patel, R. Rubber Chem Technol 1981;54:91-100.
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[11] Coran AY, Patel R. Rubber Chem Technol 1981;54:892-903.
[12] Bhowmick AK, Chiba T, Inoue T. J Appl Polym Sci 1993;50:2055-64.
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[13] Roychoudhury N, Bhowmick AK. J. Appl Polym Sci 1989;38:1091-109. [14] Barlow JW, Paul DR. Polym Eng Sci 1984;24:525-34.
[15] Chen L, Qin Y, Wang, X, Zhao X, Wang F. Polymer 2011;52:4873-80. [16] Suttiwijitpukdee
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Sato
H,
Unger
M,
Ozaki
Y.
Macromolecules
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[17] Chien RH, Lai CT, Hong JL. J. Phys. Chem C 2011;115:20732–739. [18] Han SH, Kim JK, Pryamitsyn V, Ganesan V. Macromolecules 2011;44:4970–76. [19] Oshinski AJ, Keskkula H, Paul DR. Polymer 1996;37:4891-907.
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[20] Li LP, Yin B, Zhou Y, Gong L, Yang MB, Xie BH, Chen C. Polymer 2012;53:3043-51.
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[21] Laura DM, Keskkula H, Barlow JW, Paul DR. Polymer 2003;44:3347-61. [22] Jiang W, Yu D, Jiang B. Polymer 2004;45:6427-30. [23] Wu S. Polymer 1985;26:1855-63. [24] Liu ZH, Li RKY, Tjong SC, Qi ZN, Wang FS, Choy CL. Polymer 1998;39:44336. [25] Liu ZH, Zhang XD, Zhu XG, Li RKY, Qi ZN, Wang FS, Choy CL. Polymer 1998;39:5035-45.
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[26] Paul DR, Bucknall CB. Polymer Blends. Vol. 1. New York, Willy; 2000. [27] Ren Z, Cheng S, Zhang G, Ma D, Yang X. J Phys Chem B 2008;112:1926-34. [28] Jha A, Bhowmick AK. J Appl Polym Sci 2000;78:1001–08.
[30] Thomas S, Groeninckx G. Polymer 1999;40:5799–5819. [31] Danesi S, Porter RS. Polymer 1978;19:448-57.
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[32] Kang J, Smith TG, Bigio DI. AIChE J 1996;42:649-59.
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[29] Wu JH, Li CH, Chiu HT, Shong ZJ. J Appl Polym Sci 2008;108:4114-21.
[33] Favis BD, Chalifoux JP. Polym Eng Sci 1987;27:1591-600.
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[34] Bell J, Murayama T. J Appl Polym Sci 1968;12:1795–9. [35] McCrum NG. J Polym Sci 1958;27:555-78.
[36] Beck RH, Gratch S, Newman S, Rusch K. J Polym Sci Part (B) 1968;6:707-09. [37] Bohn L. Angew Makromol Chem 1971;20:129–140.
[38] Beatty JR. Rubb Chem Technol 1978;51:1044-1059.
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[39] Paul DR, Newman S. Polymer Blends. Vol. 1. New York, Academic Press; 1978. [40] Mazich KA, Jr Pllumer HK, Samus MA, Jr Killogoar PC, Jr. J Appl Polym Sci 1989;37:1877–88.
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List of Figures
Fig. 1. Structure of (a) PA6 and (b) FKM.
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Fig. 2. Torque vs. time of mixing of pristine polymers and blends at 240 °C and 100 rpm rotor speed.
Fig. 3. Weight percentage of unextracted FKM and PA6 in the blends. Fig. 4. Infrared spectra of 60 PA6/40 FKM (w/w) showing the peaks in the range from (a) 4000-3000 cm-1 and (b) 1500-650 cm-1 respectively. Fig. 5. Scheme of H-bond between N-H group of PA6 and C-F containing groups of FKM.
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Fig. 6. FESEM micrograph of PA6/FKM blends (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) and the distribution of the rubber particle size of the corresponding images.
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Fig. 7. AFM phase image of (a) 40 PA6/60 FKM (w/w) [scan area 0.5 × 0.5 µm2], (b) 60 PA6/40 FKM (w/w) [scan area 0.5 × 0.5 µm2] , (c) 40 PA6 /60 FKM_50 rpm (w/w) [scan area 2 × 2 µm2] and (d) 40 PA6 (low viscosity) /60 FKM (w/w) [scan area 2 × 2 µm2]. Fig. 8. FESEM micrograph of PA6/FKM blends after extraction in MEK solvent (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) respectively.
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Fig. 9. Plots of weight average diameter of rubber particle and average dispersed hole vs. weight percentage of polyamide.
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Fig. 10. Tensile stress vs. tensile strain plots for PA6/FKM blend compositions. Fig. 11. Temperature dependence loss tangent (tan δ) of pure polymers and PA6/FKM blends.
Fig. 12. DSC thermograms of pristine polymers (a, b) and blend compositions (c, d).
Fig. 13. Experimental and theoretical plot of complex modulus of PA6/FKM blends as a
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function of volume fraction of polyamide at 30 °C.
(a) PA6
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(b) FKM
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Fig. 1. Structure of (a) PA6 and (b) FKM.
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Fig. 2. Torque vs. time of mixing of pristine polymers and blends at 240 °C and 100 rpm
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rotor speed.
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Fig. 3. Weight percentage of unextracted FKM and PA6 in the blends.
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Fig. 4. Infrared spectra of 60 PA6/40 FKM (w/w) showing the peaks in the range from (a) 4000-3000 cm-1 and (b) 1500-650 cm-1 respectively.
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Fig. 5. Scheme of H-bond between N-H group of PA6 and C-F containing groups of FKM.
(a)
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Fig. 6. FESEM micrograph of PA6/FKM blends (a) 20/80 (w/w), (b) 40/60 (w/w), (c) (d) 60/40(w/w) and (d) 80/20 (w/w) (c) and the distribution of the rubber particle size of the corresponding images.
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Fig. 7. AFM phase image of (a) 40 PA6/60 FKM (w/w) [scan area 0.5 × 0.5 µm2], (b) 60 PA6/40 FKM (w/w) [scan area 0.5 × 0.5 µm2] , (c) 40 PA6 /60 FKM_50 rpm (w/w) [scan area 2 × 2 µm2] and (d) 40 PA6 (low viscosity) /60 FKM (w/w) [scan area 2 × 2 µm2].
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(b)
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(d)
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Fig. 8. FESEM micrograph of PA6/FKM blends after extraction in MEK solvent (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) respectively.
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Fig. 9. Plots of weight average diameter of rubber particle and average dispersed hole vs. weight percentage of polyamide.
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Fig. 10. Tensile stress vs. tensile strain plots for PA6/FKM blend compositions.
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Fig. 11. Temperature dependence loss tangent (tan δ) of pure polymers and PA6/FKM blends.
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Fig. 12. DSC thermograms of pristine polymers (a, b) and blend compositions (c, d).
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Fig. 13. Experimental and theoretical plot of complex modulus of PA6/FKM blends as a
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function of volume fraction of polyamide at 30 °C.
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List of Tables Table 1 Formulation of mixes. Table 2 Experimental and theoretical value of torque of PA6/FKM blends at 5 min of
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mixing. Table 3 Peak Position and their Assignment in the IR Spectra (in ATR Mode) of (a) PA6 and (b) FKM.
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Table 4 Morphology of PA6/FKM blends.
Table 5 Mechanical properties of pristine polymers and PA6/FKM blends.
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Table 6 Dynamic Mechanical and DSC Analysis of PA6 /FKM Blends
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Table 1 Formulation of mixes
100
FKM
a
Comments
(Pure Components)
a
FKM
-
100
80 PA6/20 FKM (w/w)
80
20
(Variation of
60 PA6/40 FKM (w/w)
60
40
plastic-rubber ratio)
40 PA6/60 FKM (w/w)
40
60
20 PA6/80 FKM (w/w)
20
80
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PA6
PA6
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Sample Composition
a
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Weight percent.
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Table 2 Experimental and theoretical value of torque of PA6/FKM blends at 5 min of mixing. Experimental
Theoretical
Deviation/increment
torque (Nm)
torque (Nm)
of torque
80 PA6/20 FKM
8.6
4.76
3.84
60 PA6/40 FKM
7.0
4.06
2.94
40 PA6/60 FKM
4.2
3.13
1.07
20 PA6/80 FKM
2.9
2.20
0.70
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(Nm)
Deviation/increment
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Blend Composition (w/w)
of torque (%)
80.67 72.41 34.18 31.82
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Table 3
Wave number (cm-1) (a) PA6 3296 2927, 2853 1633 1538 1462 (b) FKM 1454 1397 1353 1022 800
Functional group
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N-H stretching absorption of amide C-H stretching >C=O stretching of amide (Amide I band) N-H bending absorption of amide ((Amide II band) C-H bending
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>CH2 scissor vibration of vinylidenefluoride unit C-F stretching vibration C-F3 stretching vibration C-F and CF3 stretching C-F deformation
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Peak Position and their Assignment in the IR Spectra (in ATR Mode) of (a) PA6 and (b) FKM
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Table 4
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Morphology of PA6/FKM blends dn (nm)
dw (nm)
dw/dn
IDw (nm)
106
110
1.04
6.36
40 PA6/60 FKM
97
102
1.05
16.69
60 PA6/40 FKM
93
97
80 PA6/20 FKM
84
89
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Blend Composition (w/w) 20 PA6/80 FKM
1.04
32.25
1.06
60.61
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Table 5 Mechanical properties of pristine polymers and PA6/FKM blends
Tensile Strength (MPa)
Young's modulus (MPa)
EAB (%)
100
0.71 (±0.05)a
-
388 (±6)
Modulus @ 100 % Elongation (MPa) 0.70 (±0.04)
80
6.51 (±0.2)
46 (±4)
126 (±5)
5.47 (±0.5)
26 (±1)
664
25 (±1)
60
27.90 {21.90}b [19.00]c (±1.1)
189 {162} [150] (±7)
178 {135} [155] (±7)
25.10 {21.00} [17.90] (±1.3)
33 {34} [34] (±1)
4517 {2768} [2863]
49 {48} [48] (±1)
40
39.00 (±1.2)
200 (±5)
151 (±2)
36.60 (±1.4)
35 (±1)
5196
61 (±1)
20
52.10 (±1.0)
249 (±4)
121 (±4)
48.30 (±1.1)
38 (±1)
5875
64 (±1)
0
68.60 (±0.5)
217 (±3)
44 (±1)
-
-
195
70 (±1)
2152
-
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-
values in the parenthesis indicate the standard deviation.
b
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a
Tension Toughness Hardness Set (MJ/m3) (shore D) (%)
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Weight of FKM (%)
values within the second bracket indicate the results of the blend when maxing time and rotor speed were 3.5 min and 50 rpm respectively.
c
values within the third bracket indicate the results of the blend with low viscous PA6.
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Table 6
Dynamic Mechanical and DSC Analysis of PA6 /FKM Blends DMA a
Weight
Tgr
of FKM
(°C)
Tgp
b
DSC ′
E (MPa) at
Tgra
30 °C
(°C)
1.96
3.8
- 25.5
Tanδmax
c
(°C)
(%)
(°C)
-
-
42.0
201
43.5
0.37
189.0
- 27.0
60
-13.5
50.5
0.15
995.8
- 30.5
42.5
208
40
-15.5
53.5
0.07
1970.0
- 31.2
43.0
209
20
-18.0
52.5
0.05
0
-
64.0
-
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-7.5
2614.0
- 41.0
43.5
212
3329.0
-
47.0
220
Glass transition temperature of the plastic phase.
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Tanδ value at the maximum peak position of the rubber phase.
d
(°C)
80
Melting temperature of the plastic phase.
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c
Tmd
- 4.0
Glass transition temperature of the rubber phase.
b
Tgpb
100
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a
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Sample
S0032-3861(13)00943-9
DOI:
10.1016/j.polymer.2013.10.001
Reference:
JPOL 16532
To appear in:
Polymer
Received Date: 10 May 2013 Revised Date:
11 September 2013
Accepted Date: 1 October 2013
Please cite this article as: Banerjee SS, Bhowmick AK, Novel Nanostructured Polyamide 6/ Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties, Polymer (2013), doi: 10.1016/j.polymer.2013.10.001. 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.
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Graphical Abstract
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Novel Nanostructured Polyamide 6/Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties
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Shib Shankar Banerjee and Anil K. Bhowmick
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Novel Nanostructured Polyamide 6/Fluoroelastomer Thermoplastic Elastomeric Blends: Influence of Interaction and Morphology on Physical Properties Shib Shankar Banerjee and Anil K. Bhowmick*
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Department of Material Science and Engineering, Indian Institute of Technology, Patna – 800013, India
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ABSTRACT: Novel Polyamide 6 (PA6)/Fluoroelastomer nanostructured thermoplastic elastomeric blends were developed in the present work. The influence of interaction between
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the components and morphology on physical properties of the blends was analyzed. Scanning Electron Microscopy and Atomic Force Microscopy studies, solubility and theoretical analysis of complex modulus clearly indicated that PA6 was the continuous matrix in which fluorocarbon elastomer was present in nanoscale. Low torque ratio (0.34) of rubber/plastic, high mixing speed and long mixing time had an important role in developing
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the nanostructured morphology of the blend. Tensile strength of the thermoplastic elastomer was about 39.0 MPa which was much higher than that reported earlier and showed significant improvement with increasing PA6 content. Large shifting of the glass transition temperature
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of the rubber and the plastic phases towards the lower temperature compared to those pristine polymers was also observed. The above properties were explained with the help of interaction
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between the components and morphology.
Key words: Polyamide 6, Fluoroelastomer, thermoplastic elastomer,
*Corresponding Author. Tel/Fax: +91 612 2277380. E-mail address: [email protected] (Anil K. Bhowmick)
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1. Introduction Thermoplastic elastomers (TPEs) have appeared as one of the most important classes of speciality materials having immense potential for industrial revolution. They have unique and
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extraordinary properties such as excellent physico-mechanical properties, chemical resistance, wider processing windows, ability to bond with multiple thermoplastics, light weight and excellent price-performance ratio, all of which make these an ideal candidate for
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enormous industrial applications [1-4]. Polyamide based thermoplastic elastomers are reported to have great importance because of their excellent strength, stiffness, low friction,
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high melting point, and chemical and wear resistance [5]. Thermoplastic elastomers based on polyamides, prepared by melt mixing process are expected to have excellent heat and oil resistance [6]. Several studies of different properties of blends of polyamide and rubbers are available in the literature. Crisenza et al. studied direct 3D visualization of the phaseseparated morphology of thermoplastic elastomers from chlorinated polyethylene/nylon
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terpolyamide blends [7]. Jha and Bhowmick investigated the reactive nylon-6/acrylate thermoplastic elastomeric blends and their interaction on mechanical and dynamic mechanical properties [8]. Oderkerk and Groeninckx reported the morphology development
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by reactive compatibilization of nylon-6/EPDM blends with a high amount of rubber fraction and extensively studied the effect of viscosity ratio of rubber/thermoplastic on the blend
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morphology [9].
Coran and Patel extensively investigated the thermoplastic elastomers from various rubberplastic blends and correlated the properties with morphology, surface energy mismatch and crystallinity of the matrix phase of the blends [10-11]. Few authors investigated the interaction between immiscible polymers through the addition of reactive components [1214]. Mechanical properties, dynamic response and interaction between components (both reactive and non-reactive systems) of rubber-plastic blends are main concern with the rubber
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modified plastics systems. Many researchers reported non-covalent bonding interaction in polymer blends (mainly H-bonding interaction) through infrared spectroscopic analysis [1518].
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Among different classes of thermoplastic elastomers, blends of rubbers and plastics are gaining importance because the required property can be easily monitored by varying the ratio of the constituents, viscosity of the components and also by incorporation of additives
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like compatibilizer, crosslinking agents, fillers, etc. Even though many polyamide based TPEs are commercially available, development of polyamide-fluoroelastomer TPEs has not
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been reported in literature. Also, polyamide shows interesting processing problems and its blends with rubber without compatibilizer exhibit difficulty in processing and poor mechanical characteristics.
The present investigation reports our observations on the blends of polyamide 6 and fluoroelastomer by melt-blending process. The influence of interaction and blend ratio on
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dynamic mechanical and mechanical properties was investigated. FTIR study on H-bonding interaction between polyamide 6 and fluoroelastomer was highlighted. The molecular weight of polyamide is one of the most important parameters governing the melt viscosity and hence
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processing characteristics, interfacial adhesion and morphology of the blend. The rubber particle size in rubber toughened plastics can be reduced by an order of magnitude by
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increasing the molecular weight of polyamide keeping other process conditions constant [1922]. Our objective is to develop a thermoplastic elastomeric blend with nanodimensional elastomeric particles dispersed in a plastic matrix. To the best of our knowledge, this is the first time that a nanostructured polyamide based thermoplastic elastomeric blends has been prepared by melt blending process. Furthermore, the effects of nanostructured morphology on the physical properties of the blends are reported.
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2. Experimental Section 2.1. Materials Polyamide 6 (tradename Akulon) in granular form (viscosity number 245cm3/g, density 1130
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kg/m3, melting point 220 °C) was supplied by DSM, Nederland. Viton A (a copolymer of vinylidene fluoride (VF2) and hexafluoropropylene (HFP), density 1810 kg m-3 at 25 °C, 66% F, Mooney Viscosity, ML
1+10
at 121°C = 20) abbreviated as FKM was procured from
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DuPont Dow Elastomers, Geneva, Switzerland. The formulation of the mixes is given in
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Table I. The molecular structure of polyamide 6 and FKM is given in Fig. 1
2.2. Preparation of Thermoplastic Elastomeric Blend Compositions PA6 granules were dried in vacuum oven at 100 °C for 24 h before blending. The blends were compounded in batches of 50 g polymer in a Haake Rhecord internal mixer with roller
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type rotor. Blending was allowed to proceed for 7 min at a temperature of 240 °C and 100 rpm rotor speed. After 2 min melting of PA6 in the internal mixer, FKM was added and mixed for another 5 min under the same conditions. The changes of mixing torque and energy
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with time were recorded for each composition. After complete mixing in the internal mixer, the resulting blends were quickly removed and passed through a two-roll mill having close
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nip-gap at room temperature to make sheets and were cut into small pieces. One blend composition {40 PA6/60FKM (w/w)} which showed best thermoplastic elastomeric behaviour was also processed in the following two ways: i. Similar procedure as above was followed except high viscosity PA6 was replaced by low viscosity PA6. ii. Mixing was allowed to proceed for 3.5 min at a temperature of 240 °C and 50 rpm rotor speed.
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The procedure followed after complete mixing was same as before.
2.3. Molding of Thermoplastics Elastomers
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Test specimens (1-1.5 mm thick) were prepared by means of micro-injection molding machine (Haake MiniJet II) at 260 °C cylinder temperature and 60 °C mold temperature. Injection pressure and holding pressure were 700 bar and 350 bar respectively. The injection
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and holding times were 5 sec each.
2.4.1. Solubility Measurement
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2.4. Characterization
Approximately 0.5 g (accurately weighed) of each blend, put inside a packet of Whatman-41 filter
paper, was immersed in 50 ml of methyl ethyl ketone (MEK) solvent at room
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temperature and the solvent was renewed every 24 h. After complete extraction as measured from constant residual weight, the samples were removed and dried to constant weight using a vacuum oven at 70 °C. The weight percent of entrapped FKM into polyamide matrix was
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calculated from the final and the initial weight of the rubber in the blends. Similar procedure for solubility in formic acid was followed to calculate the weight percent of unextracted
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polyamide.
2.4.2. IR Spectroscopy in Attenuated Total Reflectance (ATR) Mode IR spectra (in ATR mode) of the pristine polymers and the blends were taken on smooth film using a Perkin–Elmer (model 400) spectrophotometer with a resolution of 4 cm-1 and sixteen scans were averaged for each spectra.
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2.4.3. Microscopy 2.4.3.1.
Field Emission Scanning Electron Microscope
Morphology of various blends was recorded from FESEM (S4800 Hitachi microscope) at an
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acceleration voltage of 15.0 kV at a working distance of 6 mm. A thin layer of Platinum was sputter coated for 60 sec on the smooth sample surface to avoid charging on exposure to electron beam during FESEM analysis. Measurement of particle diameters (di) from the
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photomicrographs gave the rubber particle size distributions for different compositions of the
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blend. From these data, the number average diameter (< dn >) and weight average diameter
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(
The inter-particle distance between dispersed phases play an important role on the mechanical properties of polymer blends. Wu [23] first proposed the concept of the average thickness of the ligament between two particles i.e. the inter-particle distance that decides whether a blend will be tough or brittle. This concept has been applied for several rubber-
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plastic blends and rubber-toughened plastics [24-25]. For the cubic array of spherical
is the volume fraction of rubber and
2.4.3.2.
Atomic Force Microscopy
is the rubber particle size.
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where,
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particles, the inter-particle distance can be obtained from the following equation.
Intermittent Contact Mode Atomic Force Microscopy, ACAFM (Agilent 5500 Scanning Probe Microscope) images were obtained for supporting the nanostructures of the blends. The
2.4.4. Mechanical Tests
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resonance frequency of the tip was 130-280 kHz and the force constant was 48 N/m.
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Tensile test was carried out in a Zwick Universal Testing Machine (UTM) model Z010 at room temperature at a test speed of 200 mm/min. The samples were prepared following ISO
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527-2-5A specification. The average of three tests is reported here. For tension set measurement, the samples were extended up to 50 % in the tensile direction at a rate of 50 mm/min and kept at that position for 10 min at room temperature. It was relaxed back to unstressed condition and the percentage change in dimension in tensile direction was measured after 24 h and reported as tension set. Shore D hardness of samples was obtained by DIN ISO 7619 test method at room temperature.
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2.4.5. Dynamical Mechanical Analysis The dynamic mechanical properties of the blends and individual polymers were measured using a DMA of TA Instruments (model Q800) in tension mode. All the samples (height 16
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mm, width 9.8 mm, thickness 1mm) were analyzed at a constant frequency of 1 Hz, at a heating rate of 2 °C/min and a strain amplitude of 30 µm over a temperature range of - 100
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°C to + 100 °C.
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2.4.6. Differential Scanning Calorimetry
Thermal analysis of the pristine polymers and the blends was carried out using a Perkin– Elmer Differential Scanning Calorimeter (DSC 8500) at a heating and cooling rate of 5 °C min-1 in an inert atmosphere (N2 atmosphere). Standard aluminum pans were used to analyze the sample. The experiment was conducted from – 70 °C to 240 °C. The data of second
Results and Discussion
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3.1. Mixing torque
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3.
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heating cycle was used to eliminate the thermal history.
Torque during polymer mixing is a complex combination of shear and elongational flow in actual mixing environment. Mixing torque behavior of the pristine polymers and the blends is shown in Fig. 2. The torque-time curves for all the blends in Fig. 2 gave two peaks. The first peak was attributed to the increase in mixing torque (i.e. viscosity) value due to resistance exerted on the rotor by the unmolten polymer followed by a decrease in viscosity due to the complete melting of polyamide 6. Upon the addition of FKM into polyamide, the torque value again increased which gave the second peak. A gradual decline in torque value for the
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blend was observed with the rise in stock temperature and both reach a steady-state value after a certain time of mixing. The uniform torque value after a specified time suggested good level of mixing of the two components. The extents of mixing and mixing time are important
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parameters for immiscible polymer blends, since final properties of the blends depend on these parameters. From Fig. 2, it was clear that the equilibrium mixing torque was maximum for PA6 and lowest for FKM. This was attributed to the high melt viscosity of the PA6
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3.1.1. Theoretical analysis of mixing torque
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compared to FKM under the specified mixing conditions.
In order to assess the interaction between the blend components, theoretical calculation of mixing torque was investigated. We have calculated the theoretical value of torque based on
Here,
and
) phases using the following equation
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volume fraction of plastic ( ) and rubber (
are the torque values of the plastic and the rubber respectively determined
from Haake rheomixer at 240 °C and 100 rpm rotor speed. Table 2 highlights the theoretical
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and experimental values of the torque, which vary as a function of rubber proportions. It was observed that the experimental value of the torque of PA6/FKM blends was much higher than
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the theoretical value in the entire composition range. Melt mixing of PA6 and FKM results in an increase in molecular weight due to interaction and therefore, the viscosity of the mixture, which raises the torque during mixing. It is interesting to note that positive deviation of torque increases with increase of the plastic content and maximum positive deviation of torque (~ 81 %)) is observed when polyamide content is 80 percent by weight in the blend. It indicates that interaction of FKM and PA6 increases with increasing the PA6 content. This phenomena is also observed from the solubility characteristics and discussed later.
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3.2. Solubility Characteristics In order to determine the extent of interaction between PA6 and FKM during melt blending
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process and the amount of weight percent of FKM entrapped into hard polyamide matrix, solubility measurements have been done using their respective solvents. The data are presented in Fig. 3. The amount of unextracted FKM by MEK solvent has been taken as the
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entrapped fluoroelastomer in the polyamide matrix and this is valid when FKM part is uncured in the blend systems. The blend composition of 80 PA6/20 FKM (w/w) showed
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higher amount of unextracted rubber compared to other compositions (Fig. 3). This result revealed that maximum amount of rubber encapsulation had taken place when rubber content was near 20 percent by weight in the blend. This result supports the findings from the mixing torque measurement. It may also be assumed that some of the FKM chains were immobilized close to the boundary between the two phases. As a consequence, these rubber particles were
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embedded in the boundary region (This phenomena is also revealed from dynamic mechanical analysis and discussed later). From the solubility in formic acid, it was observed that the structure of the blend was completely collapsed and the material in the form of
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powder was found, after removal of polyamide from the blends. This result gave a direct evidence that polyamide 6 formed the continuous matrix in the whole composition regime.
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From Fig. 3, it was clear that at lower weight percent of rubber, there was no unextracted polyamide but at high weight percent of rubber in the blend, unextracted polyamide was found. As polyamide is hard matrix phase in the blend and mainly non-covalent bonded interaction occurs between the components (confirmed by FT-IR analysis and discussed later), lower amount of dispersed phase was unable to tightly hold the hard matrix part in its solvent. At higher weight percent of the dispersed phase, however, some amount of unextracted polyamide was found.
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3.3. Infrared Spectroscopic Analysis Infrared spectroscopic technique is extensively used to identify specific interactions that
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occur in various polymer systems [26-27]. It was apparent from the mixing torque and solubility measurements that an appreciable amount of interaction occurred between PA6 and FKM. Preparation of melt-blending of polyamide with various functionalized polymers was
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reported by a few researchers [28-30]. The amine and carboxyl end groups of polyamide 6 have ability to bond with other reactive polymers having suitable functional groups at their
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chain. The fluoroelastomer used in this blend contains C-F, C-F2 and C-F3 groups in their chain. Hence, H-bonds with N-H group of polyamide 6 can be formed. The characteristic peak positions of the pristine polymers are assigned in Table 3. Polyamide 6 gave N-H stretching intensity at 3296 cm-1, but the N-H stretching intensity was shifted to 3290 cm-1 after blending (Fig. 4). Similarly, C-F and C-F3 stretching of pure FKM appeared at 1397 cmand 1353 cm-1 respectively. C-F and C-F3 combined stretching appeared at 1022 cm-1. In the
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blend systems, the intensities were shifted to 1375 cm-1 (C-F stretching), 1338 cm-1 (C-F3 stretching) and 1018 cm-1 (C-F and C-F3 stretching) respectively (Fig. 4). The reason for the
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shift might be due to interactions between the blend components and formation of H-bond between N-H group of PA6 and C-F containing groups of FKM. The probable scheme of H-
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bond formation is given in Fig. 5.
3.4. Microstructure Analysis The properties of multiphase polymeric blends strongly depend on interfacial chemistry and microstructure. The morphology of PA6/FKM blend at different compositions is highlighted in Fig. 6. The distribution of the rubber particle size for each composition is also given in the
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adjoining Fig. 6. For the distribution of particle size, each rubber particle was approximated as a circle and the diameter of the particle was measured manually in the SDM image analysis software (available with S4800 Hitachi microscope), because they could not
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identified by the threshold of black-white images. Table 4 reflects the blend morphology in
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terms of number average diameter dn, weight average diameter dw, particle size distribution
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and inter-particle distance of the dispersed phase. Interestingly, the dimension of the
dispersed FKM phase in polyamide matrix was found in the nanometric level in the whole
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composition regime of the blends. Weight average diameter of the dispersed phase was in
the range of 89-110 nm. The average particle size of the rubber increased with increase in
its weight percent in the blends and narrow particle size distribution was observed with a
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typical value of 1.05. As shown in Table 4, it is clear that inter-particle distance based on
weight average particle diameter increased as the dispersed phase volume fraction decreased. Maximum inter-particle distance (60.61 nm) was observed when the dispersed phase was
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near 20 percent by weight in the blends. It is interesting to note that continuous polyamide phase with dispersed fluorocarbon elastomer was observed for the composition range of 80
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PA6/20 FKM to 20 PA6/80 FKM (w/w). This is the first time that nanostructure has been reported for a thermoplastic elastomeric composition. The prominent factors that decide the morphology of blends are viscosity ratio (viscosity of disperse phase/viscosity of matrix), blend composition, elasticity, shear rate during mixing, interfacial modification and additives
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such as fillers, plasticizers, etc. It was reported that the melt viscosity differences between the components and composition ratios determined the morphology of the blends of isotactic polypropylene and ethylene-propylene rubbers for the same processing conditions [31]. The
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development of nano-dimensions of the elastomer phase in blend systems could be explained in the light of torque ratio of the components and very high melt-viscosity of the matrix
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component (both were measured by mixing torque behavior). It was reported that transient breakup process of the disperse phase during melt-mixing was more dependent on matrix viscosity than the viscosity ratio [32]. We have taken PA6 having different viscosity values. These were then used to prepare PA6/FKM blends to show the effect of viscosity ratio of the component on the properties of nanostructured blends. Here, we measured the morphology after 7 min of mixing followed by injection molding of the samples. AFM photograph of 40 PA6 (low viscosity)/60 FKM (w/w) is shown in Fig. 7d. The particle size of the rubber phase
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was found to be 0.30 µm. It was found that under the above specified condition, the torque value was 6.4 Nm for PA6, 0.8 Nm for low viscosity PA6 and 2.2 Nm for FKM (Fig. 2). The ratio of the torque values was 0.34 for FKM/ PA6 and 2.75 for FKM/ PA6 (low viscosity).
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This indicated that high melt-viscosity of PA6 compared to FKM under the specified mixing conditions, was one of the crucial parameters for the development of the nanostructured blends. The generation of increased stress during mixing and molding operations, due to
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highly viscous PA6 gave extensive break-up of the dispersed phase and reduced its particle size compared to low viscosity PA6. It is clear from AFM phase image (Fig. 7). However,
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when mixing speed and time were reduced to their half value compared to the specified conditions for 40 PA6 (high viscosity)/60 FKM (w/w), the dimension of the dispersed phase was 0.22 µm (Fig. 7c). Favis and Chalifoux explained that minimum particle size of the dispersed phase (0.5 µm) for PP/PC blends was obtained when torque ratio was approximately 0.25 [33]. So, the torque ratio of the components and viscosity of each matrix
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play an important role in the achievement of nanostructure blends. Increasing the concentration of the dispersed phase increases the number of particles and hence particleparticle interaction and subsequent coalescence. As a result, the particle size increases. If we
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consider the collision of two identical dispersed particles in shear flow during melt mixing in an internal mixer, coalescence may occur in the following way: the dispersed particles
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deform upon collision and then matrix material is expelled from the regime between them. The basic influencing parameters which determine the extent of coalescence are the volume fraction of the disperse phase, viscosity of each phase, radius of particle and the interfacial tension which is the major driving force for coalescence [26]. In order to support the nanostructured morphology of PA6/FKM blends, further investigation was done by Atomic Force Microscopy. Fig. 7a-b reflects the phase morphology of 40 PA6/60 FKM and 60 PA6/40 FKM (w/w) blends. The light yellow colour in the phase image
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corresponds to the soft rubber particles, which are dispersed in the hard polyamide matrix. The dimension of this dispersed phase was also found in nanometer range (60-120 nm), which supports and coincides with the findings from FESEM analysis.
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In order to further substantiate the dimension of the dispersed phase, FKM phase was preferentially etched by MEK. Fig. 8 exhibits the dark holes dispersed throughout the polyamide matrix for the entire blend compositions. The average diameter of the dispersed
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holes was calculated by measuring the diameter of about 50 randomly selected small holes for each composition of the blends. We have considered both minor and major axis to
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measure the diameter of the dispersed hole in Fig. 8. It is clear from Fig. 9 that the weight average diameter of the dispersed particle is almost the same as the average diameter of dispersed hole. However, the small holes are observed to merge in Fig. 8.
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3.5. Mechanical Properties
The basic differences among elastomer, plastic and thermoplastic elastomer are directly reflected in the mechanical properties of the blends. Typical tensile stress vs. tensile strain
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curves for the blend compositions are presented in Fig. 10. Table 5 shows the mechanical properties for the same and the pristine polymers. The stress-strain curve for 20 PA6/80 FKM
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(w/w) blend exhibited low modulus and considerable elongation, typical behavior for unvulcanized or undervulcanized rubber. Conversely, the curve for 80 PA6/20 FKM (w/w) blend composition showed high modulus with a narrow yield point and a consistent plastic deformation plateau, which reflected the typical behavior for a thermoplastic material. Finally, the curves for the compositions of 40 PA6/60 FKM (w/w) and 60 PA6/40 FKM (w/w) displayed typical shape of thermoplastic elastomers. The curves were marked by the absence of an yield point before failure (Fig. 10). As mentioned in Table 5, tensile strength,
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Young’s modulus, tension set, toughness and hardness increased with increase of PA6 content in the blends. Interestingly, the blend composition of 80 PA6/20 FKM (w/w) exhibited very high toughness value compared to other compositions in the blends (Table 5).
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So, this particular composition may behave as a rubber-toughened plastic. Surprisingly, for the PA6/FKM blends, a very high value of tensile strength (27.9-39.0 MPa) was observed, when they were compared with polyamide systems reported in the literature [7-8] probably
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due to interaction and injection molding adopted here. Also, significant increase of tensile strength was found with an increase of polyamide content in the blends, which was not
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reported earlier for chlorinated polyethylene/nylon terpolymer and nylon/acrylate systems [78]. This significant improvement of tensile strength and very high value of toughness in the blends could be explained on the basis of nanostructured morphology in the PA6/FKM blends. In order to support this, we prepared blends of 40 PA6/60 FKM (w/w) in the following two ways: (1) by changing the mixing conditions (dimension of dispersed phase
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was 0.22 µm), and (2) by using low viscous polyamide under the specified conditions (dimension of dispersed phase was 0.30 µm). As shown in Table 5 and Figure 10, tensile strength and toughness are found to be low when dimension of elastomer was in micron
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range. It was proven that PA6 and FKM have H-bonding interaction. The higher the surface area, higher will be interaction. Nanostructure of elastomers generates higher surface area
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than the microstructured blends, which results in higher tensile strength and toughness.
3.6. Dynamic Mechanical Analysis Fig. 11 represents the temperature dependence of tan δ of pristine polymers and their blends over a temperature range of - 100 to 100 °C. In the above temperature range, PA6 mainly revealed two main transitions, α and β, which appeared at 64 °C and - 63 °C, respectively.
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The α-transition, designated as its Tg (64 °C), was attributed to motion within the amorphous phase and mainly dependent on crystallinity of the materials. Conversely, the β–transition was due to carbonyl groups of polyamide forming H-bond with absorbed water [34]. In the
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case of FKM, main transition was found at - 4 °C and ascribed to the glassy to rubbery transition or Tg. The blend composition, 40 PA6/60 FKM (w/w) showed three main damping peaks- one at - 13.5 °C corresponding to the glassy to rubbery transition of FKM, another
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broad peak at 50.5 °C due to Tg of PA6, and the third broad peak at - 61.5 °C arising due to PA6. Interestingly, the peak at - 4 °C was shifted towards the lower temperature and the β
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transition temperature for polyamide has a tendency to move towards the Tg of rubber. The Tg of the plastic was also shifted towards the lower temperature. However, the shift in the peak positions of the blend revealed interaction and small amount of miscibility between the phase components.
The DMA results of PA6/FKM blends are presented in a tabular form (Table 6) in terms of
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(i) glass transition temperature of the rubber phase, Tgr, (ii) glass transition temperature of the plastic phase, Tgp, (iii) (tanδ)max of the rubber phase with respect to weight percent of rubber in the blends and (iv) storage modulus, E′ at 30 °C of the blends. (tanδ)max and Tgr of rubber
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phase decreased with increasing the plastic content in the blends, which could be explained on the basis of interaction between FKM and PA6 phases. It was also found that Tgp of the
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plastic phase decreased with increasing the rubber content in the blends. Storage modulus, however, showed an increasing trend with an increase of the plastic content in the blends. McCrum reported that the amplitude of dynamic transition of a component of a composite was directly attributed to the relative quantity of the component itself [35], which was further confirmed from our previous work. [8].
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Let us visualize each rubber domain as a combination of two separate phases- bulk rubber ‘active’ in the dynamic transition and immobilized rubber particles close to the boundary between two phases with restricted mobility. Thus, it might be said that with increasing the
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plastic content in the blends, the relative quantity of the bulk rubber ‘active’ for the dynamic transition decreased and immobilization of the rubber particle to the phase boundary increased, thus, resulting in a reduction of (tanδ)max value of the rubber phase. The second
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result, i.e. the decrease in Tg of the rubber phase, can be ascribed in the light of thermal stresses built-up in each of the rubber particles, as the sample was cooled from the melt
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during mixing and molding operations. A state of triaxial tension was expected due to greater thermal contraction of rubber compared to the glassy matrix [36-37]. This thermal stress could only be developed when there was sufficient adhesion between the two phases. As a consequence of this, the free volume and chain mobility of the rubber increased, resulting in a decrease of the glass transition temperature. Such type of contrasting feature was also found
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for nylon/acrylate blends [8]. Interestingly, we observed that Tg of FKM was shifted by 9.5 °C in the PA6/FKM blends, whereas there was a shift of only 1.5 °C for the nylon/acrylate blends having the same composition (both cases in the negative direction compared to their
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pristine glass transition temperature). As the fluoroelastomer has maximum tendency to shrink (3.17 %) compared to other elastomers [38], its free volume increment is much more
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prominent for fluorocarbon rubber based blends compared to acrylate rubber counterpart. To the author’s knowledge, this large shift of Tg of rubber has not been reported before for rubber-plastic blends.
The third result, i.e. the decrease in Tg of the plastic phase, could be explained on the basis of enhancing segmental motion of the glassy polymer chains when attached to a more mobile component i.e., lower Tg. As PA6 was mixed with FKM chains which were rubbery and
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mobile in nature, the flexibility and mobility of the polyamide chains are increased due to
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mutual interaction, leading to a decrease in its Tg value.
3.7. Differential Scanning Calorimetric Studies
In order to assess the influence of thermal stresses built-up in the FKM particles during mixing and molding with PA6, DSC experiment was done. We have taken the data from
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second heating to eliminate the thermal history. Fig 12 shows the DSC thermograms of
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pristine polymers and blend compositions.
PA6 showed melting a endotherm at 220 °C and glass transition at 47 °C. FKM revealed glass transition at - 25.5 °C. The blend composition, 40 PA6/60 FKM (w/w) showed base line shift in three regions- one at - 30.5 °C corresponding to the glass transition of FKM, another at 42.5 °C due to Tg of PA6, and the third at 208 °C arising due to melting of PA6. Table 6
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summarized the thermal analysis data of PA6, FKM and blend systems. It was interesting to note that the glass transition temperature of the rubber phase of the blend was lower than that of the control FKM and crystalline melting point of PA6 decreased substantially with the
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addition of FKM. This phenomena is in line with the findings from dynamic mechanical analysis confirming that thermal stresses built-up in the FKM particles are permanent and
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responsible for lowering the glass transition of the rubber in the blend systems.
3.8. Theoretical Analysis of Complex Modulus In order to assess the behaviour of the two-phase blend from the component property data, the existing theoretical Kerner model was used [39]. The complex moduli of the blends at 30 °C (a temperature between the two Tg values of the components) as well as the theoretical values of the modulus of the blends based on Kerner’s model are plotted in Fig. 13. The
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Kerner’s equation for a binary blend where one component is dispersed in a matrix material
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is
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where
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E* is the complex modulus of the blend, Em* is the complex modulus of the matrix, Ei* is the
complex modulus of the dispersed phase. ν is the Poisson ratio of the blend, and νi and νm are
the Poisson ratio of the dispersed and the matrix phases respectively. The Poisson ratio of the blend, which varies with temperature, has been calculated from the following equation [40]
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It was clear from Fig. 13 that the experimental complex moduli for the PA6/FKM blends, especially at higher plastic content, are very close and almost coincide with those obtained from Kerner’s rigid matrix–soft filler model, suggesting the formation of polyamide as the continuous matrix. This result from the theoretical analysis of complex modulus is in good
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accord with the findings from the morphology and the solubility measurement of the blends
4. Conclusions
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that polyamide is the continuous matrix in the blends systems.
Nanostructured blends of PA6 and FKM were developed and fully characterized. The
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following conclusions are drawn:
(1) The mixing torque and IR spectroscopic analysis indicated interaction between the polymers.
polyamide matrix.
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(2) The solubility measurement revealed that rubber particles are encapsulated into hard
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(3) FESEM, AFM, solubility and theoretical analysis of complex modulus provided strong evidence for continuous polyamide phase in the blends. (4) The low torque ratio (0.34) of rubber/plastic, high mixing speed and long mixing time played an important role in developing the nanostructured morphology. (5) Very high value of tensile strength of the blends was observed when they were compared with polyamide systems reported in literature and high value of toughness was also observed when FKM was near 20 weight percent in the blend.
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(6) In the dynamic mechanical analysis, it was observed that the (tanδ)max as well as Tg of rubber phase decreased with increase of plastic content as a layer of restricted chain mobility was formed near the phase boundary and a state of triaxial tension was generated during
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cooling of the samples. (7) Micro-structural analysis showed the nanoheterogeneity of the blend which is reported first time for thermoplastic elastomer blends and also narrow particle size distribution of the
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dispersed phase.
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References
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Dekker; 2001.
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List of Figures
Fig. 1. Structure of (a) PA6 and (b) FKM.
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Fig. 2. Torque vs. time of mixing of pristine polymers and blends at 240 °C and 100 rpm rotor speed.
Fig. 3. Weight percentage of unextracted FKM and PA6 in the blends. Fig. 4. Infrared spectra of 60 PA6/40 FKM (w/w) showing the peaks in the range from (a) 4000-3000 cm-1 and (b) 1500-650 cm-1 respectively. Fig. 5. Scheme of H-bond between N-H group of PA6 and C-F containing groups of FKM.
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Fig. 6. FESEM micrograph of PA6/FKM blends (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) and the distribution of the rubber particle size of the corresponding images.
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Fig. 7. AFM phase image of (a) 40 PA6/60 FKM (w/w) [scan area 0.5 × 0.5 µm2], (b) 60 PA6/40 FKM (w/w) [scan area 0.5 × 0.5 µm2] , (c) 40 PA6 /60 FKM_50 rpm (w/w) [scan area 2 × 2 µm2] and (d) 40 PA6 (low viscosity) /60 FKM (w/w) [scan area 2 × 2 µm2]. Fig. 8. FESEM micrograph of PA6/FKM blends after extraction in MEK solvent (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) respectively.
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Fig. 9. Plots of weight average diameter of rubber particle and average dispersed hole vs. weight percentage of polyamide.
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Fig. 10. Tensile stress vs. tensile strain plots for PA6/FKM blend compositions. Fig. 11. Temperature dependence loss tangent (tan δ) of pure polymers and PA6/FKM blends.
Fig. 12. DSC thermograms of pristine polymers (a, b) and blend compositions (c, d).
Fig. 13. Experimental and theoretical plot of complex modulus of PA6/FKM blends as a
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function of volume fraction of polyamide at 30 °C.
(a) PA6
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(b) FKM
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Fig. 1. Structure of (a) PA6 and (b) FKM.
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Fig. 2. Torque vs. time of mixing of pristine polymers and blends at 240 °C and 100 rpm
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rotor speed.
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ACCEPTED MANUSCRIPT
AC C
EP
Fig. 3. Weight percentage of unextracted FKM and PA6 in the blends.
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
Fig. 4. Infrared spectra of 60 PA6/40 FKM (w/w) showing the peaks in the range from (a) 4000-3000 cm-1 and (b) 1500-650 cm-1 respectively.
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
Fig. 5. Scheme of H-bond between N-H group of PA6 and C-F containing groups of FKM.
(a)
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 6. FESEM micrograph of PA6/FKM blends (a) 20/80 (w/w), (b) 40/60 (w/w), (c) (d) 60/40(w/w) and (d) 80/20 (w/w) (c) and the distribution of the rubber particle size of the corresponding images.
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 7. AFM phase image of (a) 40 PA6/60 FKM (w/w) [scan area 0.5 × 0.5 µm2], (b) 60 PA6/40 FKM (w/w) [scan area 0.5 × 0.5 µm2] , (c) 40 PA6 /60 FKM_50 rpm (w/w) [scan area 2 × 2 µm2] and (d) 40 PA6 (low viscosity) /60 FKM (w/w) [scan area 2 × 2 µm2].
ACCEPTED MANUSCRIPT
RI PT
(b)
TE D
M AN U
SC
(d)
AC C
EP
Fig. 8. FESEM micrograph of PA6/FKM blends after extraction in MEK solvent (a) 20/80 (w/w), (b) 40/60 (w/w), (c) 60/40(w/w) and (d) 80/20 (w/w) respectively.
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
EP
TE D
Fig. 9. Plots of weight average diameter of rubber particle and average dispersed hole vs. weight percentage of polyamide.
TE D
M AN U
SC
RI PT
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AC C
EP
Fig. 10. Tensile stress vs. tensile strain plots for PA6/FKM blend compositions.
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
AC C
Fig. 11. Temperature dependence loss tangent (tan δ) of pure polymers and PA6/FKM blends.
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 12. DSC thermograms of pristine polymers (a, b) and blend compositions (c, d).
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 13. Experimental and theoretical plot of complex modulus of PA6/FKM blends as a
AC C
EP
TE D
function of volume fraction of polyamide at 30 °C.
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List of Tables Table 1 Formulation of mixes. Table 2 Experimental and theoretical value of torque of PA6/FKM blends at 5 min of
RI PT
mixing. Table 3 Peak Position and their Assignment in the IR Spectra (in ATR Mode) of (a) PA6 and (b) FKM.
SC
Table 4 Morphology of PA6/FKM blends.
Table 5 Mechanical properties of pristine polymers and PA6/FKM blends.
AC C
EP
TE D
M AN U
Table 6 Dynamic Mechanical and DSC Analysis of PA6 /FKM Blends
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Table 1 Formulation of mixes
100
FKM
a
Comments
(Pure Components)
a
FKM
-
100
80 PA6/20 FKM (w/w)
80
20
(Variation of
60 PA6/40 FKM (w/w)
60
40
plastic-rubber ratio)
40 PA6/60 FKM (w/w)
40
60
20 PA6/80 FKM (w/w)
20
80
SC
PA6
PA6
RI PT
Sample Composition
a
AC C
EP
TE D
M AN U
Weight percent.
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Table 2 Experimental and theoretical value of torque of PA6/FKM blends at 5 min of mixing. Experimental
Theoretical
Deviation/increment
torque (Nm)
torque (Nm)
of torque
80 PA6/20 FKM
8.6
4.76
3.84
60 PA6/40 FKM
7.0
4.06
2.94
40 PA6/60 FKM
4.2
3.13
1.07
20 PA6/80 FKM
2.9
2.20
0.70
AC C
EP
TE D
M AN U
SC
(Nm)
Deviation/increment
RI PT
Blend Composition (w/w)
of torque (%)
80.67 72.41 34.18 31.82
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Table 3
Wave number (cm-1) (a) PA6 3296 2927, 2853 1633 1538 1462 (b) FKM 1454 1397 1353 1022 800
Functional group
SC
N-H stretching absorption of amide C-H stretching >C=O stretching of amide (Amide I band) N-H bending absorption of amide ((Amide II band) C-H bending
TE D
M AN U
>CH2 scissor vibration of vinylidenefluoride unit C-F stretching vibration C-F3 stretching vibration C-F and CF3 stretching C-F deformation
EP AC C
RI PT
Peak Position and their Assignment in the IR Spectra (in ATR Mode) of (a) PA6 and (b) FKM
ACCEPTED MANUSCRIPT
Table 4
RI PT
Morphology of PA6/FKM blends dn (nm)
dw (nm)
dw/dn
IDw (nm)
106
110
1.04
6.36
40 PA6/60 FKM
97
102
1.05
16.69
60 PA6/40 FKM
93
97
80 PA6/20 FKM
84
89
AC C
EP
TE D
M AN U
SC
Blend Composition (w/w) 20 PA6/80 FKM
1.04
32.25
1.06
60.61
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Table 5 Mechanical properties of pristine polymers and PA6/FKM blends
Tensile Strength (MPa)
Young's modulus (MPa)
EAB (%)
100
0.71 (±0.05)a
-
388 (±6)
Modulus @ 100 % Elongation (MPa) 0.70 (±0.04)
80
6.51 (±0.2)
46 (±4)
126 (±5)
5.47 (±0.5)
26 (±1)
664
25 (±1)
60
27.90 {21.90}b [19.00]c (±1.1)
189 {162} [150] (±7)
178 {135} [155] (±7)
25.10 {21.00} [17.90] (±1.3)
33 {34} [34] (±1)
4517 {2768} [2863]
49 {48} [48] (±1)
40
39.00 (±1.2)
200 (±5)
151 (±2)
36.60 (±1.4)
35 (±1)
5196
61 (±1)
20
52.10 (±1.0)
249 (±4)
121 (±4)
48.30 (±1.1)
38 (±1)
5875
64 (±1)
0
68.60 (±0.5)
217 (±3)
44 (±1)
-
-
195
70 (±1)
2152
-
EP
TE D
M AN U
SC
-
values in the parenthesis indicate the standard deviation.
b
AC C
a
Tension Toughness Hardness Set (MJ/m3) (shore D) (%)
RI PT
Weight of FKM (%)
values within the second bracket indicate the results of the blend when maxing time and rotor speed were 3.5 min and 50 rpm respectively.
c
values within the third bracket indicate the results of the blend with low viscous PA6.
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Table 6
Dynamic Mechanical and DSC Analysis of PA6 /FKM Blends DMA a
Weight
Tgr
of FKM
(°C)
Tgp
b
DSC ′
E (MPa) at
Tgra
30 °C
(°C)
1.96
3.8
- 25.5
Tanδmax
c
(°C)
(%)
(°C)
-
-
42.0
201
43.5
0.37
189.0
- 27.0
60
-13.5
50.5
0.15
995.8
- 30.5
42.5
208
40
-15.5
53.5
0.07
1970.0
- 31.2
43.0
209
20
-18.0
52.5
0.05
0
-
64.0
-
M AN U
SC
-7.5
2614.0
- 41.0
43.5
212
3329.0
-
47.0
220
Glass transition temperature of the plastic phase.
TE D
Tanδ value at the maximum peak position of the rubber phase.
d
(°C)
80
Melting temperature of the plastic phase.
EP
c
Tmd
- 4.0
Glass transition temperature of the rubber phase.
b
Tgpb
100
AC C
a
RI PT
Sample