Preparation, characterization and properties of acid functionalized multi-walled carbon nanotube reinforced thermoplastic polyurethane nanocomposites

Materials Science and Engineering B 176 (2011) 1435–1447

Contents lists available at ScienceDirect

Materials Science and Engineering B journal homepage: www.elsevier.com/locate/mseb

Preparation, characterization and properties of acid functionalized multi-walled carbon nanotube reinforced thermoplastic polyurethane nanocomposites Aruna Kumar Barick, Deba Kumar Tripathy ∗ Rubber Technology Centre, Indian Institute of Technology, Kharagpur 721302, West Bengal, India

a r t i c l e

i n f o

Article history: Received 14 December 2010 Received in revised form 29 July 2011 Accepted 1 August 2011 Available online 9 September 2011 Keywords: Thermoplastic polyurethane Multi-walled carbon nanotube Nanocomposites Morphology Thermomechanical properties Melt intercalation

a b s t r a c t The multi-walled carbon nanotube (MWNT) reinforced thermoplastic polyurethane (TPU) nanocomposites were prepared through melt compounding method followed by compression molding. The spectroscopic study indicated that a strong interfacial interaction was developed between carbon nanotube (CNT) and the TPU matrix in the nanocomposites. The microscopic observation showed that the CNTs were homogeneously dispersed throughout the TPU matrix well apart from a few clusters. The results from thermal analysis indicated that the glass transition temperature (Tg ) and storage modulus (E ) of the nanocomposites were increased with increase in CNTs content and their thermal stability were also improved in comparison with pure TPU matrix. The rheological analysis showed the low frequency plateau of shear modulus and the shear thinning behavior of the nanocomposites. The electrical behaviors of the nanocomposites are increased with increase in weight percent (wt%) of CNT loading. The mechanical properties of nanocomposites were substantially improved by the incorporation of CNTs into the TPU matrix. © 2011 Elsevier B.V. All rights reserved.

1. Introduction The CNT has gained tremendous momentum in the present decade because of their wide range of potential applications in several areas since its discovery in 1991 by Iijima [1]. MWNT is considered to be the most promising nanomaterial due to its high flexibility, low mass density, high aspect ratio (typically > 103 ) and exceptional mechanical and electrical properties. The MWNT is an ideal reinforcing nanofiller to develop a new generation high strength, lightweight and multi-functional polymer nanocomposites [2]. In the past decade, several fabrication methods have been successfully commercially implemented to fabricate CNT based TPU nanocomposites, i.e. in situ polymerization of monomers in the presence of CNT [3], solution casting method [4] and melt mixing process [5]. Among these, the melt blending is of particular interest for fabrication of polymer nanocomposites because of the ease with which the process may be scaled up to an industrial standard without hampering the ecology and economy very much. The advantage of this technique is its speed and simplicity, not to mention its compatibility with standard industrial processing techniques. The driving force for dispersion of CNTs into a polymer matrix is high

∗ Corresponding author at: Present address: Veer Surendra Sai University of Technology (VSSUT), Burla 768 018, Sambalpur, Odisha, India. Tel.: +91 3222 283196/663 2430211; fax: +91 3222 255303/663 2430204. E-mail addresses: [email protected], [email protected] (D.K. Tripathy). 0921-5107/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.mseb.2011.08.001

processing temperature and high shearing force developed during melt processing. The important factors for producing superior CNT reinforced polymer nanocomposites using melt intercalation technique are also widely established as (i) effective homogeneous dispersion of CNTs in polymeric matrix and (ii) introduction of strong interfacial interaction between the polymeric matrix and CNTs to constitute efficient interphase load transfer characteristic. Furthermore, very poor nanoscale dispersion and poor interfacial interaction obstructs the complete utilization of carbon nanotubes for significant reinforcing and toughening of polymeric matrix. The problems may be substantially controlled by the functionalization of the CNTs. Therefore, it is a competent and productive way to counteract against the aggregation/agglomeration of the nanotubes that effectively subjects to the homogeneous dispersion of the CNTs within polymer matrix and to control the interfacial shear strength issues [6]. The first ever preparation of polymer nanocomposites using MWNT as nanofiller was reported in literature by Ajayan et al. [7]. Since then significant number of papers has been published regarding the processing and characterization of the fabricated polymer/CNT nanocomposites for high performance and multifunctions [8]. The literature survey reveals that a very little research works have been published about MWNT reinforced TPU nanocomposites through melt blending technique. Chen et al. [9] have studied the TPU/MWNT nanocomposites fabricated by melt extrusion process. It was concluded that the homogeneous dispersion of MWNTs throughout the TPU matrix and the strong interfacial adhesion between the functionalized MWNTs and the TPU

1436

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

matrix are responsible for the significantly enhancement of Young’s modulus and tensile strength with retention of high elongation at break by incorporating up to 9.3 wt% MWNTs. Pötschke et al. [5] have focused on the effect of different types of MWNTs and carbon blacks as well as different melt mixing methods on the electrical volume resistivity and tensile test behavior of the MWNT and carbon black filled TPU nanocomposites. Raja et al. [10] have investigated the influence of functionalized CNT on the thermal and mechanical properties of the TPU nanocomposites prepared by melt mixing process. The tensile modulus and thermal stability of the nanocomposites were improved while the Tg of the TPU was slightly decreased by the addition of functionalized CNTs, which is attributed to the better dispersion of functionalized CNTs and to improve interaction between functionalized CNTs and TPU matrix. Abdullah et al. [11] have comparatively evaluated the mechanical and electrical properties of the carbon fiber (CF) and CNT based TPU nanocomposites prepared by extruder followed by injection molding. The storage modulus, thermal stability and thermal conductivity of the TPU nanocomposites were significantly improved due to good adhesion between the TPU matrix and the fillers. Jiang et al. [12] have employed melt mixing method to prepare nanocomposites based on TPU and MWNT. The result indicated that the unmodified MWNT were uniformly dispersed within the TPU matrix and the microphase separation structures of the TPU nanocomposites were slightly affected by the presence of MWNT. The mechanical properties of the TPU nanocomposites were prominently increased at both room temperature and at 120 ◦ C. In addition, the electrical and thermal properties were greatly improved with increase in MWNT content. Zhang et al. [13] have reported the effect of conductive network formation in a polymer melt on the conductivity of MWNT/TPU nanocomposites. The relationship between the processing conditions and dynamic percolation behavior was established with an objective to reduce the percolation threshold and to improve the conductivity of the TPU nanocomposites. The main objective of this article is to prepare MWNT based novel TPU nanocomposites with tailor-made properties for wide range applications of the particular TPU. The effect of MWNT loading on material properties of TPU/MWNT nanocomposites was investigated. This work substantially differs from the previous studies, because the TPU matrix consists of higher level of non-aromatic hard segments. The acid functionalized MWNT presumably enhance the interfacial interaction with the TPU matrix via hydrogen bonding. The work deals with the preparation of TPU/CNT nanocomposites by a simple melt blending process. The influence of the dispersion of nanotubes on resultant nanostructure morphology was examined by microscopy techniques and X-ray diffraction to ascertain the structure–property relationship of this system. The mechanical properties of the TPU/MWNT nanocomposites were evaluated through dynamic mechanical analysis (DMA) and tensile testing. Moreover, the functional properties such as thermal and electrical properties were also discussed. The dynamic rheological properties of the TPU/MWNT nanocomposites were studied to evaluate the processing characteristics.

2. Experimental 2.1. Materials Commercial biomedical grade aliphatic, polyether based thermoplastic polyurethane (Tecoflex® EG 80A injection grade) used as the matrix material for this work was supplied by Lubrizol Advanced Materials Inc., ThermedicsTM Polymer Products, Ohio, USA. Tecoflex EG 80A (around 35% of hard segments) having Shore

A hardness = 72, density = 1.04 g/cm3 and its constituent formulation contains methylene bis(cyclohexyl) diisocyanate (HMDI) hard segment, polytetramethylene glycol (PTMG) soft segment (molecular weight = 1000 g/mol) and chain extender 1,4 butane diol (BD).Carboxyl functionalized multi-walled carbon nanotube (COOH–MWNT) of 95 weight percent (wt%) purity contains 3.86 wt% COOH groups used for the research work was supplied by Cheap Tubes Inc., Vermont, USA. COOH–MWNT having bulk density: 0.27 g/cm3 , true density: ∼2.1 g/cm3 , outer diameter: <8 nm, inner diameter: 2–5 nm, length: 10–30 ␮m, ash content: <1.5 wt%, specific surface area (SSA): 500 m2 /g and electrical conductivity (EC): >102 S/cm. From the material data sheet, the energy dispersive X-ray spectroscopy (EDXS) analysis shows that COOH–MWNT contains carbon (C), aluminium (Al), chlorine (Cl), cobalt (Co) and sulfur (S) elements of 97.44%, 0.19%, 1.03%, 1.10% and 0.24% contents, respectively. The Raman spectra of MWNT shows two modes of vibrations i.e. D-band (defect/disorder induced mode) and G-band (tangential mode) at around 1360 and 1580 cm−1 , respectively [14]. All laboratory grade chemical reagents used (nitric acid (HNO3 ), sulfuric acids (H2 SO4 ) and dichloromethane (CH2 Cl2 )) are supplied by Merck Specialities Private Limited, Mumbai, India. 2.2. Purification of MWNTs The purification of MWNTs is a necessary step to remove the transition metal catalyst, which is utilized in the process for production of the CNTs all together with amorphous carbons. The HNO3 is a most effective reagent for purification and functionalization of CNTs and it has traditionally constituted as the first step in many purification schemes. Even though HNO3 is valuable for separation of the catalyst from CNT, it produces adverse effect through destruction of CNT to produce amorphous carbons [15]. The purification and functionalization scheme reported by Lozano et al. [16] was followed to carry out the same. The purification was conducted by refluxing CNTs in CH2 Cl2 for 5 days at 35 ◦ C followed by thorough washings using deionized water and refluxing further for a period of 24 h at 90 ◦ C. The CNTs were rinsed again with deionized water and vacuum filtered by a filter paper (Whatman quantitative filter paper, cat. no.: 1440090, grade no.: 40, ashless (0.007% ash)) for 24 h and dried in a circulating hot air oven at 120 ◦ C for 48 h. Furthermore, the purification of MWNTs was conducted by refluxing MWNTs in the mixture of 60:40 ratios of 65% solution or concentrated HNO3 and 65% solution of H2 SO4 to separate the impurities (catalyst and carbons) from the CNTs and to activate functional groups present on the CNTs surface. The residues were repeatedly washed by deionized water and vacuum filtered followed by drying in a circulating hot air oven at 120 ◦ C for 48 h. 2.3. Preparation of TPU/MWNT nanocomposites TPU/CNT nanocomposites of various formulations were prepared by melt mixing method through mixing of the MWNTs with the TPU matrix in an internal batch mixer (Thermo Scientific Haake PolyLab OS Rheomix, Thermo Electron Corporation, Massachusetts, USA) having a mixing chamber volume of 85 cm3 with a fill factor of 0.7. The mixing of samples was carried out at temperature of 185 ◦ C with a rotor speed of 100 rpm and mixing time of 8 min. The mixing for more time in an internal batch mixture may lead to breakdown of the CNT length as well as the degradation and chain scissoring of the TPU matrix, which drastically deteriorate the material properties of TPU/MWNT nanocomposites. TPU pellets and MWNTs were dried before processing to remove the water content if any in the supplied materials at 80 and 120 ◦ C in a preheated vacuum oven for 6 and 24 h, respectively. Different com-

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

1437

positions by wt% were prepared by taking 0.5, 2.5 and 5.0 wt% of MWNT in 100 phr (parts per hundred) of TPU matrix. This work was carried out at low wt% MWNT loading because it reduces the cost and at the same time substantially enhances the material properties. The compounds were passed once through a cold two-roll mill (Farrel Bridge Limited, Lancashire, England) immediately after batch mixing to achieve thick sheets and then cut into small pieces. 2 mm thick sheet specimens for all compositions were prepared using compression molding machine (Moore Press, GE Moore and Son, Birmingham, UK) at 185 ◦ C for 3 min with a pressure of 5 MPa. The compression-molded sheets were cooled to room temperature under adequate pressure and samples for mechanical testing were punched from the compression-molded sheets. The virgin thermoplastic polyurethane sample is coded as ‘TPU’ and TPU nanocomposite samples are designated as ‘TPUXMC’, where ‘X’ stands for wt% of CNT in TPU matrix and ‘MC’ stands for MWNT–COOH carbon nanotubes.

2.4. Measurements and characterization Fourier transform infrared spectrometer (model Nexus-870, Thermo Nicolet Corporation, Wisconsin, USA) equipped with an attenuated total reflectance (ATR) probe attachment was used to evaluate the interfacial interaction between CNT and TPU matrix. Wide angle X-ray diffraction analysis was performed with a high resolution X-ray diffractometer (model PW 3040/60, X’Pert PRO, Philips PANalytical B.V., Almelo, Netherlands) with Cu-K␣ ˚ radiation source. Field emission scanning electron ( = 1.54 A) microscope (model Supra 40, Carl Zeiss SMT AG, Oberkochen, Germany) study was carried out using the Zeiss Everhart-Thornley secondary electron detector for evaluation of the nanostructured morphology of the cryofractured surface of the TPU/CNT nanocomposites. The sample surfaces were gold (Au) coated by means of manually operated sputter coater (model SC7620, Polaron Brand, Quorum Technologies Ltd, East Sussex, UK) machine. The distribution of nanofiber into the TPU matrix and interfacial region of nanofiber and matrix were studied using a high-resolution transmission electron microscope (model JEM 2100, JEOL Limited, Tokyo, Japan). The images of the nanocomposites were captured by means of charge couple device multi-scan camera (model 794, Gatan, Inc., California, USA). The samples for analysis were prepared using an ultramicrotomy with a Leica Ultracut UCT (Leica Microsystems GmbH, Vienna, Austria). Thermal stability and composition of TPU and TPU/CNT nanocomposites were measured by thermogravimetric analysis (Q50 V6.1 series, TA Instruments, Delaware, USA). The Tg , melting point and crystallization temperature of nanocomposite samples were evaluated by means of differential scanning calorimetry (Q100 V8.1 series, TA Instruments). Dynamic mechanical thermal analyzer (2980 V1.7B series, TA Instruments) was used to measure the dynamic mechanical properties of nanocomposites. The melt rheological properties of the samples were carried out using rubber process analyzer (RPA 2000, Alpha Technologies, Ohio, USA). The alternating current (AC) conductivity ( ac ), dielectric constant (εr ) and dissipation factor (DF) of neat TPU and TPU/CNT nanocomposites were measured by a LCR meter (7600 precision LCR Meter, model B, QuadTech Inc., Massachusetts, USA). The tensile properties of the samples were determined according to the test procedure of ASTM D-412-98 using dumbbell shaped specimens with a universal testing machine (Model 4468, Instron Corporation, Massachusetts, USA). The tear tests of the samples were carried out using an unnicked 90◦ -angle test piece by Instron universal testing machine according to ASTM D-624-81 specification. The hardness of the 6 mm samples were measured by Shore A durometer hardness tester (model SHR-A-SUPER-Y2 K, TRSE Testing Machines, Blue

Fig. 1. FTIR spectra of the (a) pristine MWNT and (b) neat TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% MWNT loadings.

steel Engineers Pvt. Ltd., Mumbai, India) as per ASTM D-2240-05 standard test procedure. 3. Results and discussion 3.1. Fourier transform infrared spectroscopy (FTIR) The FTIR spectra of the carboxylic (–COOH) group grafted MWNT is shown in Fig. 1(a). The absorbance peak at 1720 cm−1 is assigned to the stretching vibrations of acid carbonyl (>C O) group of the carboxylic acid groups present on the MWNT surface [17]. The broad strong intense band centered at around 3440 cm−1 is attributed to the O–H stretching mode in carboxylic acid group [18] and also inferred to the presence of hydroxyl (–OH) group on the MWNT surface. The results indicate that carboxylic acid group is attached to the surface of MWNT. The absorption peaks featured at 1580 cm−1 is attributed to the C–C bond stretching vibration of carbon skeleton of the MWNT bulk [19]. The peak at 1020 cm−1 is referred to the vibration of C–O bond in primary alcohol present on the MWNT surface. The strong absorption peak at around 1101 cm−1 is identified as the C–O stretching mode of the characteristic of ether linkage (C–O–C) [20]. The peaks at 2920 and 2850 cm−1 are correspond to the stretching vibration of –CH3 > CH2 and >CH– groups of the MWNT skeleton. The small peak 1630 cm−1 is assigned to the stretching vibration of C C bond in aromatic ring of MWNT strand [21].

1438

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

Scheme 1. Represents the most probable interaction exist via hydrogen bonding between MWNT and TPU matrix.

The FTIR spectra of neat TPU and TPU/MWNT nanocomposites are shown in Fig. 1(b), which help to investigate the interface interaction between –COOH modified MWNT and TPU matrix. The broad band at 3050–3750 cm−1 is the overlapping of the free and hydrogen bonded amine (–NH) group stretching of urethane linkage [22] and –OH stretching of MWNT–COOH. The free and hydrogen bonded –NH groups vibrational stretching are appeared at around 3456 and 3328 cm−1 , respectively. The IR stretching modes at about 1712 and 1695 cm−1 are attributed to the free and hydrogen bonded >C O groups present in urethane linkage (–HN–COO–) of the TPU matrix, respectively. The area under these absorption peaks significantly increases with increase in wt% of MWNT loading, which concluded that the physical adhesion due to the hydrogen bond between both the –NH and >C O group of the TPU matrix and –COOH group of MWNTs. Furthermore, the magnitude of hydrogen bond increases with increase in MWNT concentration. It may be interpreted that the –COOH functional groups reside on the surface of the MWNT possible contributed to the enhancement of the interfacial interaction with the TPU matrix through hydrogen bond formation. Scheme 1 represents the most probable interaction exist via hydrogen bonding between MWNT and TPU matrix. The hydrogen bonding index (HBI) was determined by considering free and hydrogen bonded carbonyl groups through Gaussian curve fitting method. The degree of phase mixing (DPM) was evaluated from value of HBI [23]. The DPM of pure TPU is 49.26%, whereas it increased to 62.16, 63.41 and 64.10 for 0.5, 2.5 and 5.0 wt% MWNT loading, respectively. This result concluded that the interfacial interaction between TPU matrix and MWNT noticeably increases with increase in wt% of CNT loading. The hydrogen bonding between >C O groups of urethane linkage of TPU hard segments and hydrogen atom of the –COOH groups of MWNTs is prominently favorable as evident from the more shifting of the >C O stretching peak position than the –NH groups. The existence of the interfacial interaction between TPU and MWCNT significantly enhanced the material properties of TPU/MWNT nanocomposites. 3.2. Field emission scanning electron microscope (FESEM) Fig. 2 shows the FESEM images of pristine MWNTs. The CNTs show random curly or entangled structures and possess high aspect ratio even after purification with outer diameter about 10 nm. The FESEM microphotographs at ×50,000 and ×200,000 magnifications for the MWNT reinforced TPU nanocomposites are shown in Fig. 3(a–c) from that the microstructure characteristics and

Fig. 2. FESEM microphotograph of the neat MWNT at (a) ×50,000 and (b) ×200,000 magnifications.

reinforcement mechanism can be observed. The CNTs are homogeneously dispersed and distributed at all nanotube loading, which is shown in low-resolution pictures. The bright white dots visualized in the images are attributed to the uniformly dispersed CNT stubs or broken nanotube tips exposed from the TPU matrix in the plane perpendicular to the compression molding direction and grayish white color of the micrograph surface referred to the wetting of the CNT by the TPU matrix [4]. This morphology indicates that the most of the CNTs are broken into two parts rather than being pulled out from the TPU matrix in the course of cryo-fracturing process, which suggest a strong TPU–MWNT interfacial adhesion established within the nanocomposites [24]. The –COOH groups significantly stabilized the MWNT dispersion due to very good adhesions with the TPU matrix, which is attributed to the increase in polarity of the MWNTs by the functional groups and the interaction of the –COOH group of CNTs with urethane (–NHCO–) group of the TPU matrix. The volume information of nanocomposites is ambiguously detected due to the good contrast between the conductive nanotubes and the insulating TPU matrix that shows the distribution of nanotubes below the surface. It is observed from the images that the nanotubes are distinctly indentified and excellent well-dispersion state of CNT within TPU matrix is achieved with no sign of CNT aggregations up to 2.5 wt% CNT loading while noticeable CNT bundles/clusters are visualized at 5 wt% CNT loading. The FESEM microphotographs of the nanocomposites revealed that the MWNTs are well adhered to the TPU matrix, which may be due to the high shear melt mixing condition. Moreover, the randomly dispersed MWNTs are seemed to be coated or wrapped and also satisfactorily wetted by means of the TPU matrix phase. It is

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

1439

Fig. 3. FESEM microphotographs of TPU/MWNT nanocomposites with (a) 0.5 wt%, (b) 2.5 wt% and (c) 5.0 wt% MWNT loadings at ×50,000 (left) and ×200,000 (right) magnifications.

clearly observed from the FESEM pictures that the bottom of the modified CNTs end are covered up by the TPU matrix because the functionalized CNTs have a good compatibility with the polymer matrix. The FESEM images for higher CNT loading samples show that the randomly horizontal distribution of CNTs were observed on the cross-section parallel to the direction of the pressure during the compression molding process of the nanocomposites [18]. Furthermore, the CNTs having both the ends embedded in the TPU matrix are horizontally distributed on the cryo-fractured surface and become straight along or slightly curly in nature. It is visible from the FESEM pictures that the MWNT forms a highly entangled or interconnected network structure within the TPU matrix at 5.0 wt% CNT loading. However, some small CNT bundles are vis-

ible in the high-resolution picture at high nanotube content that protruding out from the elastomeric TPU matrix. Furthermore, the belt like CNTs are visualized at 5.0 wt% CNT loading that may be an individual nanotube or bundle of MWNTs interconnected between two TPU lumps present in the fracture surface. 3.3. High-resolution transmission electron microscopy (HRTEM) The HRTEM microphotographs of different wt% of MWNT reinforced TPU nanocomposites are shown in Fig. 4(a–c) at ×25,000 and ×50,000 magnifications to find out the state of dispersion of CNTs in TPU matrix through melt mixing. The low magnification micrographs show that the CNTs are uniformly distributed in the entire

1440

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

Fig. 4. HRTEM microphotographs of TPU/MWNT nanocomposites with (a) 0.5 wt%, (b) 2.5 wt% and (c) 5.0 wt% MWNT loadings at ×25,000 (left) and ×50,000 (right) magnifications.

surface of the TPU matrix and also individual nanotubes are present without any significant agglomeration or clusters. At high magnification, the individual nanotubes were clearly observed at low nanotube loading and dispersed well in the TPU matrix, whereas highly entangled (interconnected) network structure of CNT in the TPU matrix is visualized at high CNT content. The TEM micrograph

of 5 wt% MWNT based TPU nanocomposites show that the CNTs are significantly dispersed in the TPU matrix, whereas prominent black regions are visualized that indicates interlocking structure of bundles of pristine MWNTs. The high viscosity developed in high shear internal melt mixing condition hinders the dispersion of CNT as a result CNT aggregates are formed within the TPU matrix [18].

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

Fig. 5. WAXD patterns of neat MWNT, virgin TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% MWNT loadings.

3.4. Wide angle X-ray diffraction (WAXD) Fig. 5 shows the X-ray diffraction (XRD) patterns of pristine TPU matrix, neat MWNT and TPU/MWNT nanocomposites. The virgin TPU matrix shows a very strong broad amorphous peak at about ˚ 2 = 20◦ of (1 1 0) reflection plane with interchain spacing of 4.44 A, which is relevant to the existence of short range regular ordered structure of the both hard and soft domains along with disordered structure of amorphous phase of the TPU matrix [21]. The peak become broadens and the intensity is reduced with the addition of MWNT, which implied that the MWNT significantly affect the well short-range microstructural phases of the both soft and hard segments of the TPU matrix. It is assumed that the aggregation/agglomeration dynamics of both soft and hard segments of TPU matrix was significantly disordered, which may be due to the presence of strong interfacial interactions between MWNT and TPU matrix. Furthermore, the resulted steric hindrance effect of the individual or bundle nanotubes also influences the wellorganized accumulation of the soft and hard phases of TPU matrix [12]. The XRD pattern of MWNT exhibits a sharp peak at around 2 = 26◦ and a broad peak centered at 2 = 43◦ corresponding to the (0 0 2) and (1 0 0) Bragg reflection planes having interlayer spac˚ respectively. The peaks are attributed to ing of 3.43 and 2.10 A, the ordered regular arrangement of the concentric cylinders of the graphitic carbon atoms in nanotubes [25]. The XRD patterns of the TPU/MWNT nanocomposites displayed both the peaks assigned to the MWNT and virgin TPU, respectively. The diffraction peak (0 0 2) corresponds to the MWNT is completely disappeared for the XRD patterns of the TPU/MWNT nanocomposites, which may be due to the homogeneous dispersion the CNTs within TPU matrix [26]. The shear force developed during melt mixing succeeded to rupture and pull apart the MWNTs that result in the destruction of the short-range order of the concentric cylindrical geometry of MWNTs. The peak assigned to the (1 0 0) plane of the MWNTs in the TPU nanocomposites became broader and the intensity of the peak increased with increase in MWNT loading. 3.5. Thermogravimetric analysis (TGA) Fig. 6 represents the thermogravimetric and corresponding derivative (DTG) curves for neat TPU and TPU/MWNT nanocomposites. The thermal degradations of TPU are taken place at two stages due the presence of thermodynamically incompatible hard and soft segments within the TPU matrix. The first and second stages of thermal decomposition are corresponding to the hard and soft segments, respectively. The temperature at the maximum mass loss

1441

Fig. 6. TGA and DTG thermograms of neat TPU and its nanocomposites containing 0.5, 2.5 and 5.0 wt% of MWNT loadings in nitrogen atmosphere.

rate in the TGA curve is popularly known as the decomposition temperature of the nanocomposites. The TGA curves for TPU/MWNT nanocomposites shift towards higher temperature as compared to the neat TPU matrix, which indicates that the inclusion of CNT raises the temperature of maximum thermal degradation. The peak temperatures of the DTG curves or the temperature at maximal degradation rates for hard and soft domains of the neat TPU matrix are around 321 and 408 ◦ C that were improved to approximately 347 and 418 ◦ C, respectively at 2.5 wt% COOH–MWNT filled TPU nanocomposites. This observation indicates that the thermal stability of the nanocomposites is substantially enhanced by the incorporation of the MWNT that is attributed to the excellent thermal stability of CNTs and also interfacial interactions present in between the TPU matrix and MWNT [27]. The formation of chemical bond between TPU matrix and COOH–MWNT significantly reduced the TPU–CNT thermal boundary resistance, which is responsible for the smooth transfer of heat from the TPU matrix to the CNTs. Furthermore, this facilitates the uniform heat distribution throughout the nanocomposites without assembling of excessive heat on the surface as a result the thermal decomposition temperature is increased [28]. It is observed from Fig. 6 that the first and second onset temperatures of the thermal degradation of hard and soft segments are significantly improved with the addition of CNT. This result may be accredit to the homogeneous dispersion of the MWNTs, rise in thermal conductivity of the TPU matrix by the addition of CNTs [29] and the formation and stabilization of COOH–MWNT bonded macro radicals generated during the degradation process [30]. Fig. 7 represents the thermogravimetric and corresponding derivative curves for thermo-oxidative degradation of neat TPU and TPU/MWNT nanocomposites. The thermo-oxidative degradation of polyurethane generally occurs in two stages. The first stage is associated with the degradation of hard segments, which involves the chemical dissociation of the urethane linkage (–HN–COO–) to the original polyol (OH–(CH2 )n –OH) and isocyanate (O N C–C6 H10 –CH2 –C6 H10 –C N O). The polyol and isocyanate further undergo cleavage of the chemical bonds to form small molecules (primary amine, alkene, aldehyde, ketone, carbon dioxide, water). The second stage of thermal decomposition is assigned with the soft segment of the TPU matrix, which proceeds by the depolycondensation of metastable products (isocyanurate, carbodiimide, substituted urea and stable isocyanurate) formed in the course of degradation and polyol (OH–(CH2 )4 –OH) degradation mechanisms [31]. The thermo-oxidative degradation of both the soft and hard segments of TPU matrix are successfully delayed by the addition of MWNT, which indicates that the CNT uniformly interact with both the soft and hard segments within TPU molec-

1442

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

Fig. 7. TGA and DTG thermograms of neat TPU and its nanocomposites containing 0.5, 2.5 and 5.0 wt% MWNT loadings in oxygen atmosphere.

ular structure without any preferential association to one domain. The homogeneous dispersion of the CNTs, large heat transfer due to higher thermal conductivity of TPU/CNT nanocomposites and formation and stabilization of the CNT bonded macroradicals are main factors for enhancement of the thermal stability of TPU matrix [3]. 3.6. Differential scanning calorimetry (DSC) Fig. 8(a) represents the DSC first heating curves of neat TPU and TPU/MWNT nanocomposites. The curves exhibit two small and broad melting endothermic peaks in the first heating cycle. The lower temperature melting peak (Tm1(hard) ) at around 90–95 ◦ C is most probable due to the small fraction of short range or imperfect crystalline packing of hard segments that developed during the quenching of the hot samples to the normal temperature at the time of mixing and molding of specimens. The second melting peak temperature (Tm2(hard) ) at about 240 ◦ C may be referred to the melting point of the originally crystallized parts of the hard segments present in TPU matrix. The Tm1 and Tm2 of the TPU/MWNT nanocomposites initially decreased at 0.5 and 2.5 wt% CNT loading and then increased at 5.0 wt% CNT loading. The values of melting enthalpies (Hm1(hard) and Hm2(hard) ) associated with the melting temperatures (Tm1(hard) and Tm2(hard) ) are quantitatively very low, which confirms that the small fractions of the hard segments present in the TPU matrix are participated in the regular crystalline arrangements. The peak temperature of Tm1 and the broad extension of the Tm2 endotherms depend on the previous thermal history present within the TPU matrix. However, it shifted to the higher temperature range and finally merged into a single endothermic peak by removing the thermal history through annealing process [32]. Fig. 8(b) shows the second heating thermogram of virgin TPU and TPU/MWNT nanocomposites. The glass transition temperature (Tg ) of the soft segment phase (Tg(soft) ) of pure TPU is found at about −71 ◦ C. The values of Tg of the 0.5, 2.5 and 5.0 wt% MWNT filled TPU nanocomposites are −73, −72 and −73 ◦ C, respectively. The Tg of the TPU nanocomposites filled with MWNT is slightly lower than the Tg of pure TPU. The heat capacity change (Cp ) at the glass transition point remains very low up to 5 wt% MWNT content, which suggests that the restriction effect imposed on the molecular motion of the TPU chain is drastically immobilized by the incorporation of MWNT through interfacial interaction. The free volume of the nanocomposites effectively reduced due to the strong affinity of the acid modified CNT for the TPU matrix via formation of hydrogen bonding with the >C O groups of urethane linkage (–HNCOO–) present on the backbone chain of the TPU molecules [33].

Fig. 8. DSC thermograms for (a) first heating cycle, (b) second heating cycle, and (c) first cooling cycle of neat TPU and its nanocomposites containing 0.5, 2.5 and 5.0 wt% MWNT loadings.

From Fig. 8(b), it is evident from the second heating cycle that the hard segment of the TPU matrix exhibits a small melting endotherm (Tm2 ) centered at around 250 ◦ C, which may be attributed to the disruption of predominantly hard segment structures and length distributions [31]. The enthalpy of melting associated with the endotherm is of very small value because the TPU matrix containing a few quantity of small order hard segments crystalline structures (total hard segment content in TPU matrix is about 35%) [34]. All the samples show endothermic peak at about 22 ◦ C, which is attributed to the crystallization melting of the soft segment phase within TPU matrix. From Fig. 8(b), it is shown that

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

1443

the melting point tended to increase but the heat of fusion prominently increased with the addition of COOH–MWNTs than that of the pristine counterpart. However, the increase in melting point of the TPU matrix is not substantially high by the incorporation of functionalized MWNTs, which is probably due to the formation of interconnected network between MWNT and TPU. It is assumed that the incorporation of CNT significantly enhanced the directionally aligned orientation of the TPU molecular chains, which developed a well-defined long-range crystalline structure within TPU matrix. Furthermore, the result indicates that the addition of MWNTs accelerated the microphase separation structures within the TPU matrix and altered the molecular weights of the segments or clusters of the hard and soft segments of TPU matrix [35]. The crystallinity of the TPU/MWNT nanocomposites is enormously increased by the addition of CNT. It is noteworthy to mention that the CNT act as a positive nucleating agent, which is responsible for the enhancement of the melt crystallization process [36]. The nucleating effect of the CNT is contributed more effectively to the crystallization phenomenon with increasing in COOH–MWNTs content that further enhances the enthalpy of melting. Deka et al. [17] reported that the improvement of soft segment crystallinity through addition of acid functionalized MWNT is attributed to the significant enhancement of the shape memory characteristic of the TPU matrix. Fig. 8(c) represents the DSC first cooling cycle of the neat TPU and TPU/MWNT nanocomposites. The neat TPU did not show any well-defined crystallization peak but a small shoulder appeared at around −32.5 ◦ C indicates the slow crystallization process of the PTMG soft segments of TPU matrix. The soft segment crystallization temperature (Tc(soft) ) for 0.5, 2.5 and 5.0 wt% TPU/MWNT nanocomposites are −32.0, −23.5 and −29.5 ◦ C and the enthalpy of crystallization (Hc(soft) ) are 3.6, 8.4 and 2.24 J/g, respectively. It is concluded that both Tc(soft) and Hc(soft) of soft segment significantly increases with incorporation of MWNT–COOH up to 2.5 wt% because the functionalized MWNTs significantly affect the crystallization process of the soft segment of the TPU matrix. The observation suggested that the MWNT–COOH acts as a positive heterogeneous nucleating agent that appreciably improved the intrinsic crystallization tendency of the PTMG molecules present in the soft segments of the TPU matrix [37]. The Tc(soft) and Hc(soft) are drastically reduced for TPU/MWNT nanocomposite with 5 wt% nanotube concentration, which is caused by the controlled nucleating proficiency of the functionalized MWNT at higher wt% loading as a result the crystallinity and crystallization phenomenon is considerably slowed down [38]. 3.7. Dynamic mechanical analysis The storage modulus (E ) of neat TPU and TPU/MWNT nanocomposites as a function of temperature is shown in Fig. 9(a). The E value decreases with increase in temperature throughout the scanned temperature range because of the thermal expansion of the TPU matrix. The TPU matrix becomes rigid below Tg point (glassy state) as a result the magnitude of the E is higher, whereas the value of E drops drastically after Tg point (rubbery state) because the TPU chain becomes more flexible. The E values of the TPU nanocomposites between the temperature regions of −25 ◦ C to +25 ◦ C are reduced as compare to the neat TPU matrix because the presence of MWNT reduced the short-range regular arrangement of the both soft and hard segment of the TPU matrix that overshadows the reinforcing effect of the CNTs [39]. The storage modulus increases significantly with the addition of MWNT–COOH except for 0.5 wt% filled system. The enhancement of E of TPU nanocomposites is imparted through the stiffening/reinforcing effect of the CNTs. The E of the TPU nanocomposites increases with incorporation of the nanofillers because the mobility of the soft and hard segments of

Fig. 9. DMA thermograms represent the dependency of (a) E , (b) tan ı, (c)   /Econtrol of neat TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% MWNT Ecomp loadings on temperature at a given frequency and strain amplitude of 1 Hz and 10 ␮m, respectively.

the TPU backbone chains was restricted by the interfacial interactions between the MWNT–COOH nanofillers and the TPU chains along with the stiffening/reinforcing effect imparted by the CNTs [11]. Furthermore, the higher modulus resulted in filled samples may be due to the hydrodynamic effect and the adsorption of polymer chains on the nanofiller surfaces that concurrently increases the filler effective volume [40]. The temperature dependency of tan ı for neat TPU and TPU/MWNT nanocomposites are shown in Fig. 9(b). The broad relaxation peaks in tan ı curve with a maximum intensity at around

1444

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

−48 ◦ C is due to the soft segment Tg of TPU matrix. It is clearly shown in Fig. 9(b) that the introduction of MWNT into the TPU matrix caused a shifting of Tg values to higher temperature of about −49.32 ◦ C for pure TPU to 42.80 ◦ C for the 2.5 wt% loaded TPU nanocomposites. It indicates that the molecular chain segmental mobility of the TPU matrix is significantly restricted by the incorporation of CNT via interphase adhesion among the –COOH group of MWNT and >C O group of hard segment or –O– group of soft segment of TPU matrix [41]. Furthermore, it confirmed that the MWNT is compatible with the amorphous regions of the soft segments of the TPU matrix. The tan ı peak height is slightly reduced with the addition of MWNT, which indicates the reduction of damping capacity of the TPU matrix. The reduction in peak height may be caused by the decrease in the amount of amorphous hard segment phase and also reduction in the soft domain molecular mobility at vicinity of some chains, which are responsible for the glass transition behavior. It is concluded from the above observation that the mechanical properties of the TPU matrix may enhanced by the incorporation of MWNT at lower content due to its high aspect ratio that facilitates more transfer of energy at the CNT and matrix interface [11]. The relative storage modulus (ratio of storage modulus of   nanocomposite (Ecomp ) to the control neat TPU matrix (Econtrol ) is plotted as a function of the temperature for neat TPU and TPU nanocomposites based on different wt% MWNT loadings are shown   in Fig. 9(c). The value of Ecomp /Econtrol is higher around Tg than the   temperature above or below it. The rise of Ecomp /Econtrol near Tg is attributed to the change in Poisson’s ratio of the TPU matrix from ∼0.35 (below Tg ) to ∼0.50 (around Tg ) as well as the change in the ratio of modulus of nanofiller to that of TPU matrix on passing   /Econtrol in the range through the Tg region [42]. The value of Ecomp of −25 ◦ C to +25 ◦ C is lower than the unity, which may be due to the crystallization melting point of the soft segment of the TPU matrix   as confirmed by the DSC result. The magnitude of Ecomp /Econtrol ratio is significantly improved above Tg , which has also reported in literature by Lee et al. [43] in case of CNT based polypropy  lene composites. The enhancement of Ecomp /Econtrol above Tg is contributed to the increase in reinforcing effect of the dimensionally more stable rigid CNTs consequently the TPU relaxation dynamics is significantly restricted by the physical entanglements between nanotubes and TPU matrix or formation of three dimensional interconnected crosslinked networks of CNTs at higher filler loading [44]. 3.8. Rubber processing analysis (RPA) Dynamic frequency sweep tests were used to explore network formation and microstructure of the nanocomposites. The storage modulus (G ) of neat TPU and TPU/MWNT nanocomposites measured at 145 ◦ C are logarithmically plotted as a function of angular frequency in Fig. 10(a). The G value significantly increases with increase in wt% of MWNT loading. Moreover, the increase in G is more significant particularly at low frequency region as compared to the high frequency region. A noticeable qualitative change in G versus frequency plot is observed at low frequency region for MWNT filled TPU nanocomposites. The significant plateau of the flow regime is developed at lower frequency range where G almost independent of the dynamic frequency. At low frequency region, the relaxation exponent (n) of the power law (G ∼ ωn ) drastically drops with increase in MWNT content resulting a transition of viscoelastic response from ‘liquid-like’ to ‘solid-like’ of TPU nanocomposites that beginning at 0.5 wt% MWNT content [45]. At 5.0 wt% MWNT filled TPU nanocomposites shows higher storage modulus and lower terminal region slope that indicates the higher solid-like behavior. At higher frequency region, the viscoelastic behavior of all TPU/MWNT nanocompos-

Fig. 10. RPA rheographs depict the variation of (a) G and (b) |*| of neat TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% MWNT loadings with angular frequency at temperature and shear strain of 145 ◦ C and 0.98%, respectively.

ites is almost same. The increment in value of G for 5 wt% filled MWNT based TPU nanocomposites is significant as compare to the pristine counterpart. However, the rate of increase in value of G decreases with increase in wt% of filler loading due to aggregation of the CNTs. Fig. 10(b) represents the frequency dependency of complex viscosity (|*|) for pristine TPU and TPU/MWNT nanocomposites. The magnitude of * of TPU/MWNT nanocomposites substantially increases with increase in MWNT loading because of the establishment of TPU–CNT and CNT–CNT interactions, which noticeably oppose the segmental chain molecular mobility of the TPU matrix. Furthermore, the enhancement of * with CNT loading is especially significant at lower frequency region due to the dramatic increase in the G , whereas the change in * is diminished at higher frequency region. The * value remarkably reduced with increase in frequency due to decrease in shear thinning exponent (n) of the power law (|*| ∼ ω−n ). The rate of decrease in * with rise in frequency that prominent at higher frequency region may be due to the strong shear thinning effect [46]. At higher frequency region, the dispersed MWNTs tend to align along the direction of strong shear stress, which is responsible for the demolition of the TPU chain crosslinked networks formation as a result strong shear thinning behavior is shown in CNT filled systems. The * almost decreases linearly with increase in frequency, whereas the slope of curve or reduction gradient becomes high at higher wt% MWNT loading as a consequence the * difference is significantly low at the high frequency region for TPU/MWNT nanocomposites [47].

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

1445

ites is not significantly changed with the increase of CNT loading from 0.5 to 5.0 wt%. It is observed that the magnitude of the εr for neat TPU and TPU nanocomposites is not significantly affected with the variation of applied frequency because the relaxation time (t = 1/ω = 1/2 f) for realignment of interfacial dipoles along the direction of external electric field is reduced with increase in frequency [49]. It is also concluded that the εr is approximately frequency independent at room temperature. The variation of electrical dissipation factor or loss tangent (tan ı) of virgin TPU and the TPU/MWNT nanocomposites as a function of the frequency at various wt% of CNT loading are shown in Fig. 11(b). The tan ı of the neat TPU and its nanocomposites are exponentially decayed with increase in frequency at lower frequency region (up to 105 Hz), while it rapidly increases with frequency at further higher frequency region. The dielectric loss of TPU/MWNT nanocomposites enhanced with an increase in CNT loading up to an applied frequency of 105 Hz but above which it reduced than that of the neat TPU matrix. The significant increment of the tan ı with MWNT concentration is caused by the high electrical conductivity of the CNT [50]. The values of the DF for unfilled TPU and its nanocomposites are usually observed below unity throughout the practically important frequency range, which conclude that the system is suitable for utilization in developments of capacitors. Fig. 11(c) shows the frequency dependence of the alternating current electrical conductivity ( ac ) of the neat TPU and TPU/MWNT nanocomposites at room temperature. The AC conductivity of the TPU nanocomposites noticeable increases with increase in wt% of CNT loading at low frequency range, whereas its  ac is almost equal to the virgin counterpart at high frequency range. At low frequency region, the conductivity of the nanocomposites is independent of the applied frequency, which results a direct current (DC) conductivity ( dc ) plateau region spreads up to a critical or crossover frequency (fc ) beyond that it obeys a power law behavior. It is noteworthy to mention that the formations of interconnected physical networks in the bulk of the TPU matrix due to the presence of CNT clusters are responsible for the electrical conductivity of the TPU nanocomposites [51]. The value of fc increases with increase in CNT loading. The conductivity of the TPU/MWNT nanocomposites at higher frequency range follows a single master curve. The variation of conductivity with frequency increases approximately linearly for virgin TPU is a general phenomenon of typical dielectric behavior of materials. Furthermore, its follows the same trend after the fc for different amount of MWNT loaded TPU samples, which most plausible reflects the interfacial polarization effect that plays a predominant role for the frequency dependency of the effective  ac of the nanocomposites [52]. The transition from insulation to conduction behavior of the MWNT filled TPU nanocomposites are observed in between 2.5 and 5.0 wt% CNT concentration. 3.10. Mechanical properties

Fig. 11. Variation of (a) εr , (b) tan ı and (c)  ac of neat TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% of MWNT loadings as a function of frequency.

3.9. Electrical properties Fig. 11(a) shows the effect of electrical frequency and CNT concentration on the dielectric constant or relative permittivity (εr ) behavior of the TPU/MWNT nanocomposites at room temperature. The value of εr of the nanocomposites increases with addition of CNT to the TPU matrix, which is most plausible due to the supplementary contribution of the CNTs on the polarization of TPU under the applied electric field [48]. The εr of the TPU nanocompos-

The typical stress–strain curves for the pristine TPU and TPU/MWNT nanocomposites are shown in Fig. 12(a). All the systems show a linear elastic behavior at low stress region and plastic deformation at high stress region. The tensile strength and modulus are synchronously improved in comparison with the neat TPU by the incorporation of MWNT into the TPU matrix at lower strain region. The tensile strength and modulus at break of the 0.5 wt% MWNT filled TPU nanocomposites incremented to optimum level and further addition of CNT obviously reduces the tensile strength. The initial modulus, tensile strength and elongation at break of the TPU/0.5 wt% MWNT nanocomposites were 3.43 MPa, 16.00 MPa and 1392.50%, respectively, which are enhanced by about 15.10%, 64.27% and 13.37%, respectively as compared to the pristine counterpart. The mechanical properties of the MWNT filled TPU nanocomposites strictly depends on the following critical fac-

1446

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

59.60 N/mm, respectively. It is observed that the tear strength of the TPU nanocomposites remarkably increases with increase in MWNT loading, which may be due to the large inter-twist energy of the CNTs in the TPU matrix. The tear strength of the TPU/2.5 wt% MWNT nanocomposites is 69.30 N/mm, which is incremented by about 70.86% as compared to the neat TPU. The increment in hardness is presumably due to the reinforcing effect of evenly dispersed MWNT as mentioned previously. The high aspect ratio of the CNT more efficiently restricted the molecular slippage of the TPU chains and significantly resisted the development of the crack propagation. At extremely lower strain region, the tear strength and modulus of the neat TPU matrix is higher than that of the TPU nanocomposites, whereas it significantly incremented with increase in applied strain. The tear strength appreciable increased up to 2.5 wt% filled TPU/MWNT nanocomposites but it drastically deteriorated at 5.0 wt% MWNT filled loading, because of the aggregation of some of the CNTs within the TPU matrix. The Shore hardness of the virgin TPU matrix is 72 A, whereas TPU/MWNT nanocomposites with 0.5, 2.5 and 5.0 wt% CNT loadings are 76, 79 and 83 A, respectively. It is concluded that the hardness of the TPU/MWNT nanocomposites significantly improved by the incorporation of MWNT. The hardness of the materials indicates the inherent ability of the material to restrict the localized deformation/distortion caused by the application of external stimuli. The hardness is directly associated with the interconnected crosslink networks, elasticity to plasticity, strength to modulus and porosity of the polymeric matrix [55]. The hardness of the TPU/5.0 wt% MWNT nanocomposites is improved by about 15.28% as compared to the virgin TPU matrix. The hardness of the TPU/MWNT nanocomposites increased with increase in wt% of MWNT loading may be most probably due to the reinforcing effect imparted by the dimensionally more stable MWNTs and also due to contribution from the strong interfacial interactions and good compatibility of MWNT–COOH with the TPU matrix [54].

Fig. 12. Typical stress–strain curve for (a) tensile and (b) tear test of neat TPU and its nanocomposites with 0.5, 2.5 and 5.0 wt% MWNT loadings.

tors like filler dispersion, filler geometry, aspect ratio, orientation of fillers, interfacial interaction between polymer and filler and filler concentration [20]. The interface between MWNT and TPU plays a vital role for toughening mechanism of the nanocomposites, which effectively transfer the applied force from the TPU matrix to the CNT dispersed phase [53]. Furthermore, the improvement of mechanical properties may be caused by the presence of polar groups (the acid (–COOH) group of MWNT surface and amide (–CONH–) group of TPU molecular chains) of the TPU matrix and the MWNT that introduced strong interfacial adhesion bonding between them [54]. However, both tensile strength and modulus of the 5.0 wt% MWNT filled TPU nanocomposites are not significantly enhanced as compared to the neat TPU due to the formation of CNT aggregates that result in an effective reduction of the aspect ratio (length/diameter) of the MWNTs, consequently the interfacial TPU–MWNT adhesion strength is substantially diminished. The elongation at break increases up to 2.5 wt% MWNT based TPU nanocomposites because of the well dispersion of CNTs within the TPU matrix and also most probable due to the unconstrained cubic deformation of the CNTs resulted from the unique structural feature with topology and cage [28]. The elongation at break decreases at 5.0 wt% MWNT than that of the neat TPU matrix content because of the poor dispersion of the CNTs in TPU matrix [18]. The stress–strain curves for tear properties of the neat TPU and TPU/MWNT nanocomposites are shown in Fig. 12(b). The tear strength of the neat TPU matrix is 40.56 N/mm along with 0.5, 2.5 and 5.0 wt% MWNT filled TPU matrix are 60.00, 69.30 and

4. Conclusions TPU nanocomposites with different compositions of MWNT (0.5, 2.5 and 5.0 wt%) were prepared by melt blending. The structure–property relationship of the TPU/MWNT nanocomposites was established based on the results of different analysis techniques. FTIR spectra established that strong interactions exist between the MWNT–COOH and TPU matrix. FESEM, TEM and WAXD microphotographs showed that the functionalized MWNTs effectively dispersed within the soft and hard segments of the TPU matrix through interfacial interactions between MWNT and TPU, whereas it showed aggregated or cluster structure of the CNT bundles at 5 wt% MWNT loading. TGA study showed that the incorporation of CNTs significantly improved the thermal stability caused by the high thermal conductivity of the CNTs. DSC result indicated that the melting temperature, Tg and crystallinity slightly increased with inclusion of MWNT, which suggested that the functionalized MWNT substantially influenced both crystalline and amorphous structure of the TPU matrix. DMA study revealed that the E and Tg of the nanocomposites increased with increase in wt% of CNT loading, which indicated that CNTs are compatible with the TPU matrix. RPA test determined that the G incremented with increase in MWNT content and variation of * with applied frequency showed a shear-thinning characteristic. Both different wt% of MWNT and frequency dependency of the electrical properties of the TPU/MWNT nanocomposites was studied. The εr , tan ı and  ac of the nanocomposites are comparatively higher than that of the TPU matrix. Mechanical tests showed that the addition of MWNT significantly improved the tensile properties, tear strength and hardness of the TPU matrix without sacrificing the elongation

A.K. Barick, D.K. Tripathy / Materials Science and Engineering B 176 (2011) 1435–1447

at break by incorporation of MWNTs. The homogeneous dispersion of MWNTs throughout the TPU matrix at lower MWNT loading and presence of strong interfacial adhesion between functionalized MWNTs and the TPU matrix are responsible for the significant improvement of overall material properties of the TPU/MWNT nanocomposites. The overall properties of the 5 wt% MWNT filled TPU nanocomposites are not significantly enhanced and also not reduced than that of the pristine counterpart, which indicated that the higher wt% of MWNT loading adversely affect the properties of the nanocomposites due to the poor dispersion of CNTs throughout the polymer matrix and weak adhesive bonding with the polymer matrix. Acknowledgements The authors would like to thankfully acknowledge the financial support of the Extra Mural Research Division II (EMR-II), Human Resource Development Group (HRDG), Council of Scientific and Industrial Research (CSIR), New Delhi 110012, India vide sanction no. 22(0410)/06/EMR-II, dated 13-09-2006. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]

S. Iijima, Nature 354 (7) (1991) 56–58. J.N. Coleman, U. Khan, W.J. Blau, Y.K. Gun’ko, Carbon 44 (9) (2006) 1624–1652. H. Xia, M. Song, Soft Matter 1 (5) (2005) 386–394. H. Koerner, G. Price, N.A. Pearce, M. Alexander, R.A. Vaia, Nat. Mater. 3 (2) (2004) 115–120. P. Pötschke, L. Häußler, S. Pegel, R. Steinberger, G. Scholz, Kautsch Gummi Kunstst 60 (9) (2007) 432–437. N.G. Sahoo, S. Rana, J.W. Cho, L. Li, S.H. Chan, Prog. Polym. Sci. 35 (7) (2010) 837–867. P.M. Ajayan, O. Stephan, C. Colliex, D. Trauth, Science 265 (5176) (1994) 1212–1214. Z. Spitalsky, D. Tasis, K. Papagelis, C. Galiotis, Prog. Polym. Sci. 35 (3) (2010) 357–401. W. Chen, X. Tao, Y. Liu, Compos. Sci. Technol. 66 (15) (2006) 3029–3034. M. Raja, A.M. Shanmugharaj, S.H. Ryu, Soft Mater. 6 (2) (2008) 65–74. S.A. Abdullah, A. Iqbal, L. Frormann, J. Appl. Polym. Sci. 110 (1) (2008) 196–202. F. Jiang, G. Hu, S. Wu, Y. Wei, L. Zhang, Polym. Polym. Compos. 16 (8) (2008) 471–480. R. Zhang, A. Dowden, H. Deng, M. Baxendale, T. Peijs, Compos. Sci. Technol. 69 (10) (2009) 1499–1504. http://www.cheaptubesinc.com/MWNTs.htm#multi walled nanotubesmwnts 8nm raman spectra and elemental analysis. H. Hu, B. Zhao, M.E. Itkis, R.C. Haddon, J. Phys. Chem. B 107 (50) (2003) 13838–13842. K. Lozano, J. Bonilla-Rios, E.V. Barrera, J. Appl. Polym. Sci. 80 (8) (2001) 1162–1172. H. Deka, N. Karak, R.D. Kalita, A.K. Buragohain, Carbon 48 (7) (2010) 2013–2022.

[18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53] [54] [55]

1447

J. Xiong, Z. Zheng, X. Qin, M. Li, H. Li, X. Wang, Carbon 44 (13) (2006) 2701–2707. L. Liu, Y. Qin, Z.-X. Guo, D. Zhu, Carbon 41 (2) (2003) 331–335. J. Xiong, D. Zhou, Z. Zheng, X. Yang, X. Wang, Polymer 47 (6) (2006) 1763–1766. H.-C. Kuan, Ma C-CM, W.-P. Chang, S.-M. Yuen, H.-H. Wu, T.-M. Lee, Compos. Sci. Technol. 65 (11–12) (2005) 1703–1710. A.K. Barick, D.K. Tripathy, Composites A 41 (10) (2010) 1471–1482. A.K. Barick, D.K. Tripathy, J. Appl. Polym. Sci. 117 (2) (2010) 639–654. J.-M. Yuan, Z.-F. Fan, X.-H. Chen, X.-H. Chen, Z.-J. Wu, L.-P. He, Polymer 50 (14) (2009) 3285–3291. O. Zhou, R.M. Fleming, D.W. Murphy, C.H. Chen, R.C. Haddon, A.P. Ramirez, S.H. Glarum, Science 263 (5154) (1994) 1744–1747. T. McNally, P. Pötschke, P. Halley, M. Murphy, D. Martin, S.E.J. Bell, G.P. Brennan, D. Bein, P. Lemoine, J.P. Quinn, Polymer 46 (19) (2005) 8222–8232. H. Xia, M. Song, J. Mater. Chem. 16 (19) (2006) 1843–1851. J. Xiong, Z. Zheng, W. Song, D. Zhou, X. Wang, Composites A 39 (5) (2008) 904–910. M. Moniruzzaman, K.I. Winey, Macromolecules 39 (16) (2006) 5194–5205. S. Mondal, J.L. Hu, J. Elastom. Plast. 38 (3) (2006) 261–271. A.K. Barick, D.K. Tripathy, Mater. Sci. Eng. A 527 (3) (2010) 812–823. R.W. Seymour, S.L. Cooper, J. Polym. Sci. B 9 (9) (1971) 689–694. S.-M. Yuen, C-C.M. Ma, Y.-Y. Lin, H.-C. Kuan, Compos. Sci. Technol. 67 (11–12) (2007) 2564–2573. D.J. Martin, G.F. Meijs, P.A. Gunatillake, S.J. McCarthy, G.M. Renwick, J. Appl. Polym. Sci. 64 (4) (1997) 803–817. X. Chen, J. Wang, J. Zou, X. Wu, X. Chen, F. Xue, J. Appl. Polym. Sci. 114 (6) (2009) 3407–3413. Q. Meng, J. Hu, Y. Zhu, J. Appl. Polym. Sci. 106 (2) (2007) 837–848. C.-X. Zhao, W.-D. Zhang, D.-C. Sun, Polym. Compos. 30 (5) (2009) 649–654. N.G. Sahoo, Y.C. Jung, H.J. Yoo, J.W. Cho, Macromol. Chem. Phys. 207 (19) (2006) 1773–1780. D.A. Nguyen, Y.R. Lee, A.V. Raghu, H.M. Jeong, C.M. Shin, B.K. Kim, Polym. Int. 58 (4) (2009) 412–417. A.V. Shenoy, Rheology of Filled Polymer Systems, 3rd ed., Springer-Verlag, London, 1999. T.L. Wang, C.G. Tseng, J. Appl. Polym. Sci. 105 (3) (2007) 1642–1650. L.E. Nielsen, Mechanical Properties of Polymers and Composites, vol. 2, Marcel Dekker, Inc., New York, 1974. G.-W. Lee, S. Jagannathan, H.G. Chae, M.L. Minus, S. Kumar, Polymer 49 (7) (2008) 1831–1840. I. Dubnikova, E. Kuvardina, V. Krasheninnikov, S. Lomakin, I. Tchmutin, S. Kuznetsov, J. Appl. Polym. Sci. 117 (1) (2010) 259–272. S.B. Kharchenko, J.F. Douglas, J. Obrzut, E.A. Grulke, K.B. Migler, Nat. Mater. 3 (8) (2004) 564–568. S. Huang, M. Wang, T. Liu, W.-D. Zhang, W.C. Tjiu, C. He, X. Lu, Polym. Eng. Sci. 49 (6) (2009) 1063–1068. K.Q. Xiao, L.C. Zhang, I. Zarudi, Compos. Sci. Technol. 67 (2) (2007) 177–182. R.K. Goyal, P.A. Jagadale, U.P. Mulik, J. Appl. Polym. Sci. 111 (4) (2009) 2071–2077. Y. Yang, M.C. Gupta, K.L. Dudley, R.W. Lawrence, Nanotechnology 15 (11) (2004) 1545–1548. L. Wang, Z.-M. Dang, Appl. Phys. Lett. 87 (4) (2005) 042903 (1–3). E. Logakis, Pandis Ch, V. Peoglos, P. Pissis, J. Pionteck, P. Pötschke, M. Miˇcuˇsík, M. Omastová, Polymer 50 (21) (2009) 5103–5111. S.-L. Shi, L.-Z. Zhang, J.-S. Li, J. Polym. Res. 16 (4) (2009) 395–399. G.L. Hwang, Y.-T. Shieh, K.C. Hwang, Adv. Funct. Mater. 14 (5) (2004) 487–491. J. Kwon, H. Kim, J. Polym. Sci. A 43 (17) (2005) 3973–3985. J.-Y. Kwon, H.-D. Kim, J. Appl. Polym. Sci. 96 (2) (2005) 595–604.