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Electrical conductivity and thermal stability of polypropylene containing well-dispersed multi-walled carbon nanotubes disentangled with exfoliated nanoplatelets Chien-Chia Chu, Kevin L. White, Peng Liu, Xi Zhang, Hung-Jue Sue
*
Polymer Technology Center, Department of Mechanical Engineering, Texas A&M University, College Station, TX 77843, United States
A R T I C L E I N F O
A B S T R A C T
Article history:
The morphology, crystallization behavior, electrical conductivity, and thermal stability of
Received 21 February 2012
polypropylene (PP) modified with disentangled multi-walled carbon nanotubes (MWCNTs)
Accepted 29 May 2012
is reported. Slightly oxidized MWCNT clusters were disentangled in solution by mild son-
Available online 8 June 2012
ication in the presence of exfoliated a-zirconium phosphate nanoplatelets. The disentangled MWCNTs were isolated using acid-induced coagulation to precipitate the nanoplatelets, and were subsequently reacted with octadecylamine. The recovered functionalized MWCNTs (F-MWCNTs) are disentangled and easily dispersed in a commercial PP matrix, and serve as more efficient nucleating agents than the untreated MWCNTs. The PP/F-MWCNT composites exhibit an extremely low percolation-like transition in electrical conductivity, which is attributed to the preservation of a random dispersion of disentangled F-MWCNTs upon cooling from the melt. The thermal stability of PP in air is also substantially enhanced at loadings below the percolation threshold due to the tremendous interfacial area between the polymer chains and the free radical scavenging F-MWCNTs. The present approach provides an efficient and potentially scalable route for commercial production of conductive semi-crystalline thermoplastics. The method may be adapted to uniformly disperse MWCNTs in other polymer matrices by appropriate selection of surface functionality. 2012 Elsevier Ltd. All rights reserved.
1.
Introduction
The uniform dispersion of multi-functional nanoparticles such as carbon nanotubes (CNTs) into semi-crystalline polymers is anticipated to substantially enhance a broad range of physical and mechanical properties at low enough filler concentration to enable processing with conventional highvolume methods. However, the magnitude of property improvement at a given loading has so far been limited by the poor dispersion of CNTs within the host polymer matrix, particularly for non-polar polymers such as polyethylene (PE) and polypropylene (PP). The inability to disperse CNTs without significant modification to the nanotube structure or the host
polymer has also prevented clear understanding of the mechanisms by which CNTs individually and collectively influence the bulk response of polymers [1]. Efficient integration of CNTs into a polymer matrix requires that mesoscopic aggregates be broken down prior to mixing, and that sufficient polymer/CNT compatibility exists to maintain a stable morphology. It is difficult to disperse highly anisotropic nanoparticles such as CNTs because interparticle potential scales with the particle surface area. Single-walled carbon nanotubes (SWCNTs) are well known to assemble into bundles following synthesis and exhibit cohesive energies on the order of 105 eV/lm [2]. Multi-walled carbon nanotubes (MWCNTs) tend to form highly entangled meshes that are
* Corresponding author. Fax: +1 979 8453081. E-mail address:
[email protected] (H.-J. Sue). 0008-6223/$ - see front matter 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.carbon.2012.05.063
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mechanically interwoven and may grow as large as 25 lm in the polymer matrix [3]. An additional challenge is that, even in the absence of aggregation, the loss in conformational freedom of the host polymer chain at the CNT interface will introduce significant entropic penalties that vary inversely with particle diameter [4]. To minimize this entropic penalty, it is necessary to either use larger diameter particles to limit the conformational restriction, or to blend low molecular weight polymer chains with the host polymer to diminish the magnitude of entropic loss [5]. However, these approaches inherently limit the potential performance of the nanostructured composite material. The enthalpic contribution to mixing may be improved by selecting a matrix with pre-existing structural or dielectric interactions with CNTs, or by modifying the interaction potential between the CNTs and the polymer. Pristine CNTs possess a small negative electrostatic potential that is insufficient to prevent mesoscale aggregation in common solvents [6], but adequate to promote interaction with polymer chains that possess a permanent dipole moment. Chemical oxidation further increases the electrostatic potential of the CNTs and reduces aggregate size, but may significantly damage mechanical and electrical performance if harsh treatment conditions are used [7]. Non-polar polymers that exhibit structural interactions with CNTs, particularly those containing pendant phenyl rings such as polystyrene, have also shown well-dispersed morphology [8]. In order to achieve a stable dispersion in polymers that lack a permanent dipole moment or suitable mechanism to favorably interact with CNTs, such as PE and PP, physical or chemical modification of either the CNTs or the polymer matrix is necessary. A number of approaches have been reported to enhance the enthalpic contribution associated with mixing CNTs into a PP matrix. Surfactants have been widely used to stabilize the CNTs in solution and improve interaction with various host polymers [6]. The PP matrix itself may be modified by reactive grafting, such as with maleic anhydride (MA), and has been shown to result in excellent dispersion due to the polarity and packing characteristics of the MA groups [9]. Similarly, the MA groups may be directly grafted to the MWCNT sidewalls [10]. These approaches rely on establishing weak physical interaction between the polymer and CNTs or altering the properties of the PP matrix, and generally lack a suitable mechanism to disentangle or exfoliate the CNT aggregates. Introducing an electrostatic surface charge by covalent functionalization is a widely used, effective approach to physically separate CNT aggregates, and enables the compatibility of the CNTs to be tailored by proper selection of functional group. Chen et al. [11] showed that SWCNTs could be solubilized in organic solution by an amidation reaction of a long-chain amine with carboxylic acid sites on the surface of oxidized SWCNTs. This type of approach typically provides a stable macro-scale dispersion, but has not been shown to be effective on the individual nanotube level. In this work, slightly oxidized MWCNT (O-MWCNT) clusters were disentangled by mild sonication in the presence of exfoliated a-zirconium phosphate (ZrP) nanoplatelets. The disentangled MWCNTs were isolated by the addition of an acid to coagulate and precipitate the nanoplatelets [12,13], and were subsequently reacted with octadecylamine (ODA) to yield functionalized MWCNTs (F-MWCNTs). The
F-MWCNTs are disentangled and uniformly dispersed in PP after compression molding, and exhibit stable microstructure after subsequent injection molding, which will be reported separately. The morphology, crystallization kinetics, electrical conductivity, and thermal stability of a commercial PP modified with disentangled MWCNTs are reported. The results are compared to literature findings to provide insight into the role of individual CNTs in the bulk response of polymer composites. Potential implications of the present research for polymer composite preparation and design are also discussed.
2.
Experimental
2.1.
Materials
A commercial grade PP, Novatec MA3, was supplied by Japan Polypropylene (JPP) Corporation (Mw = 260 kg/mol, PDI = 4.4). JPP reported that the PP contains 0.15 wt.% standard antioxidant compounds to provide thermo-oxidative stability during melt mixing. The isotactic pentad fraction of the PP was reported to be 93.7%. Pristine MWCNTs (P-MWCNTs) (purity >90%, average outer diameter 10 nm, length 0.1–10 lm) were purchased from Sigma–Aldrich. Commercially available ODA (CH3 (CH2)17NH2, Sigma–Aldrich, 97%) was used as received. Fully exfoliated ZrP nanoplatelets with aspect ratio of 100 were used to disentangle and disperse the MWCNTs in aqueous solution. The synthesis and exfoliation of ZrP nanoplatelets have been reported previously [14–16]. Tetra-n-butylammonium hydroxide (TBA, Sigma–Aldrich, 1 mol L1 in methanol) at molar ratio of ZrP:TBA = 1:0.8 was used to exfoliate the ZrP nanoplatelets in water.
2.2.
MWCNT functionalization
A schematic illustrating the functionalization and dispersion of disentangled MWCNTs in PP is provided in Fig. 1. As-received P-MWCNTs were oxidized according to a procedure described in our previous work [12,13,17,18]. ZrP nanoplatelets were added to the aqueous solution of O-MWCNTs at a weight ratio of MWCNT:ZrP = 1:5 by room temperature ultrasonication (Branson 2510) for 30 min. This weight fraction was selected to achieve full disentanglement and dispersion of the MWCNTs [17]. The disentangled MWCNTs were isolated from the ZrP in solution by acid-induced coagulation and precipitation of the ZrP nanoplatelets (Fig. 1a), which is reported in detail elsewhere [12,13]. Aqueous suspensions with MWCNT concentration up to 500 ppm were successfully prepared with this method. ODA powder was added to the solution of disentangled MWCNTs at weight ratio of MWCNT:ODA = 1:5 and stirred continuously at 85–90 C for 1 h. After completion of the reaction, the MWCNTs are hydrophobic and precipitate from solution. The precipitated F-MWCNTs were recovered and dried in an oven at 80 C overnight.
2.3.
Preparation of PP/MWCNT composites
Varying amounts of F-MWCNTs were added to 15 g xylene and sonicated for 1 h to obtain individual F-MWCNTs in non-polar solvent (Fig. 1b). One gram of pelletized PP was
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Fig. 1 – Schematic illustrating the functionalization and dispersion process used in this work: (a) Disentangled MWCNTs are obtained after isolation from ZrP nanoplatelets. (b) After reacting with ODA, the MWCNTs are hydrophobic and precipitate from solution. The recovered F-MWCNTs are easily dispersed in non-polar solvents, such as xylene. (c) PP pellets directly mixed with xylene/F-MWCNT solution. (d) PP/F-MWCNT powder recovered by precipitating from xylene solution with ethanol. dissolved in the xylene/F-MWCNT solution by stirring for 1 h at 125 C (Fig. 1c). The initial loading of F-MWCNTs was selected to vary mass fraction between 0.1 and 2.0 wt.% F-MWCNT in the final composite material. Ethanol was added to precipitate the PP/F-MWCNT from solution (Fig. 1d). The recovered PP/F-MWCNT powder was washed several times to remove residual xylene and dried under vacuum at 80 C for 12 h. For comparison purposes, PP/P-MWCNT composites were prepared by melt mixing P-MWCNTs with neat PP pellets via a Haake mixer (System 40) at 60 rpm and 180 C for 2 min. Similar low shear conditions were reported by Tjong et al. [19] to result in a highly entangled, poorly dispersed microstructure. Plaques with dimensions of 2 cm · 2 cm · 0.2 cm were prepared in all cases by hot-pressing (Dake 44-225) the powder at 180 C for 1 min.
2.4.
Characterization
Morphological characterization of the treated MWCNTs and PP/MWCNT composites was performed using a JEOL 2010 high-resolution transmission electron microscope (TEM) operating at 200 kV. Solution samples of the treated MWCNTs were drop-coated on copper grids containing a thin carbon coating and dried at room temperature. Bulk composite samples were thin-sectioned to thickness of about 80 nm using a Reichert-Jung Ultracut-E microtome. The chemical structure of the modified MWCNTs was characterized by measuring absorbance between 1000 and 4000 cm1, with a resolution of 4 cm1, using a Thermo 380 Fourier-Transform Infrared (FTIR) spectrometer (Nicolet). The melting and crystallization behaviors of the PP composites were studied with a TA Q20 differential scanning
calorimeter (DSC) with nitrogen flow rate of 50 mL/min. All tests were performed with about 10 mg of material hermetically sealed in aluminum pans. Prior to all DSC measurements, samples were held at 200 C for 5 min to erase any previous thermal history. Non-isothermal crystallization studies were performed at a range of cooling rates from 5 to 40 C/min. Melting behavior was monitored from the second heat at a rate of 10 C/min after cooling samples at 10 C/ min. The degree of crystallinity, Xc, was calculated from the ratio of the enthalpy of fusion of the composite samples, determined from the area under the melting endotherm, and the total enthalpy of fusion of 100% crystalline PP (DHo 209 J/g) [20]. Electrical conductivity was measured through the thickness of 3-mm thick compression molded plaques using a two-point probe (Keithley model 2400 digital source meter). To ensure minimal contact resistance, samples were coated with 100 nm thick gold electrodes (99% purity gold shots purchased from Kurt J. Lesker) prior to characterization. Thermogravimetric analysis (TGA) was conducted using a TA Instruments Q500 instrument heated from room temperature to 800 C at rate of 20 C/min in air and N2 environments. Samples with mass of about 10 mg were tested in an open ceramic pan.
3.
Results
3.1.
Dispersion and exfoliation of F-MWCNTs
Surface treatment with strong acid is a common method to purify and exfoliate CNTs, but generally requires a long exposure at elevated temperature to achieve suitable dispersion
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[7,21–23]. In this work, oxidation was carried out at room temperature for a short time period in order to introduce a slight negative charge to the MWCNT surface without causing any noticeable physical or structural damage. The mild oxidative treatment alone is insufficient to disentangle the dense MWCNT mesh (Fig. 2a), but introduces adequate negative charge to the MWCNTs to promote electrostatic tethering with the positively charged ZrP nanoplatelet surfaces [12,13,17]. The ZrP nanoplatelets aid in the exfoliation of CNTs to the individual level by concentrating sonication energy within the CNT aggregates. We previously observed no change in the CNT length distribution due to the exfoliation process [17], and recently conducted spectroscopic analysis that verified the reversible electronic coupling between the CNTs and ZrP nanoplatelets is accomplished without detectable change in electronic structure [13]. Nanoplatelet-assisted dispersion was used here to disentangle the MWCNTs prior to mixing with PP and yields individually dispersed, disentangled MWCNTs in aqueous solution without visible damage or reduction in length (Fig. 2b). The preservation of MWCNT electronic state was further confirmed in this report by measuring the electrical conductivity of thin buckypaper films prepared by solvent evaporation. The electrical conductivity of the disentangled MWCNT buckypaper was 1100 S m1, which is in close agreement with experimental works on untreated MWCNTs of similar dimension and quality [3,24]. These observations are consistent with our previous findings [12,13,17], and strongly indicate that the nanoplatelet-assisted dispersion method effectively reduces aggregates to the individual level while preserving the structural integrity and electrical properties of the MWCNTs. The ZrP nanoplatelets were removed from solution using a previously reported acid-induced coagulation method [12,13]. The disentangled MWCNTs were compatibilized with ODA to provide stability upon mixing with the non-polar PP matrix. The reaction between the ODA and MWCNTs was confirmed using FTIR, which is shown in Fig. 3. The strong absorption bands at 2855 and 2925 cm1 are assigned to the symmetric and asymmetric stretching of C–H bonds, respectively, along the backbone of the ODA molecule [6]. The peak at 1460 cm1 is assigned to the C–H bending mode. The contribution of the amine functionality is detected at 1375 cm1, which is assigned to C–N stretching. The absorbance peaks of the F-MWNTs are characteristic of an
Fig. 3 – FTIR spectra of O-MWCNTs and F-MWCNTs; inset shows detail of region from 1000 to 2000 cm1.
amidation reaction between carboxylic acid functionalities and a long-chain alkyl amine [6,11], which indicates that the ODA successfully reacts with the disentangled MWCNTs by a covalent reaction. The F-MWCNTs were easily dispersed in xylene and remain highly disentangled and stable due to the organophilicity of the ODA functionalities on the F-MWCNT surface, as shown in Fig. 4a for F-MWCNT loading of 100 ppm. The texture around the individual F-MWCNTs is attributed to the presence of excess ODA when dried directly from xylene solution. Unlike previous reports, the MWCNTs prepared in this work are disentangled and compatibilized prior to mixing with the PP matrix. PP pellets were dissolved in the xylene/ F-MWCNT solution by stirring at elevated temperature. Fig. 4b shows that the F-MWCNTs remain individually dispersed and disentangled in the PP matrix after solvent evaporation. Nanostructured polymer composite plaques at concentrations up to 2 wt.% F-MWCNT were prepared by compression molding and are shown in Fig. 5. The micrographs clearly indicate that the disentangled morphology exhibited both in solution (Fig. 4a) and upon mixing with PP (Fig. 4b), is preserved after compression molding. We have similarly
Fig. 2 – TEM micrographs of MWCNTs before and after disentanglement process: (a) O-MWCNTs do not contain significant residual catalyst particles or other impurities, but remain significantly entangled; (b) Nanoplatelet-assisted dispersion physically breaks down aggregated structures and results in stable dispersion of disentangled MWCNTs.
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Fig. 4 – TEM micrographs of F-MWCNTs prepared by covalent reaction with ODA: (a) Drop-dried specimen after re-dispersion in xylene at loading of 100 ppm; (b) PP/F-MWCNT (0.5 wt.%) composites after solvent removal.
observed excellent dispersion after dilution of masterbatch samples, as well as after injection molding from pelletized extrudate, which will be reported separately. In all cases, the samples exhibit a highly disentangled morphology consistent with that shown in Fig. 5. We are not aware of previous work that has demonstrated comparable preservation of CNT dispersion state, and will report more extensive characterization of the phase stability of the F-MWCNT in PP separately. The observed microstructural stability indicates that the
magnitude of PP/F-MWCNT interaction exceeds both interparticle attraction between F-MWCNTs and the entropic penalty of mixing in the PP matrix.
3.2.
Crystallization kinetics
The DSC thermograms of the neat PP and PP composites are shown in Fig. 6 and summarized in Table 1. The degree of crystallinity, Xc, is nearly constant for all samples, and the
Fig. 5 – TEM micrographs of PP/F-MWCNT plaques prepared by hot pressing recovered powder after solvent evaporation. F-MWCNT concentration is (a, b) 0.6 wt.%, (c, d) 1 wt.%, and (e, f) 2 wt.%.
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crystallization onset temperature, Toc, and peak crystallization temperature, Tc, shift to higher values in the presence of F-MWCNTs. The PP/P-MWCNT (0.1 wt.%) composites exhibits negligible change in Toc and Tc from the neat PP, which is in contrast with other works and attributed to the low loading and poor dispersion reported here. The significant shift in Toc and Tc observed at only 0.1 wt.% F-MWCNTs is attributed to the increased number and uniform dispersion of active nucleation sites compared to the poorly dispersed P-MWCNTs. Increasing F-MWCNT loading further increases Toc and Tc. These findings are consistent with previous results on the crystallization behavior of PP/CNT composites [9,25– 30], and confirm that the F-MWCNTs are better dispersed than the P-MWCNTs, and that the dispersion is maintained with increasing concentration. The observed trends in crystallization and melting behaviors at rate of 10 C/min are consistent over a range of rates from 5 to 40 C/min (Supplemental Fig. 1). The PP/ODA (0.5 wt.%) has lower Toc and Tc than the neat PP, which suggests that the small molecule ODA is solubilized in the PP matrix throughout the crystallization process and has no inherent nucleating ability. The decreased melting and crystallization temperatures are consistent with a wellknown plasticization effect by ODA [31], and are likely responsible for the small decrease in melting temperature, Tm, for the PP/F-MWCNT systems. These findings indicate that functionalization with ODA does not significantly influence the crystallization behavior of the MWCNTs in PP, but may cause some plasticization and disorder at the PP/F-MWCNT interface. More extensive characterization of the morphology and crystal microstructure will be reported separately.
3.3.
Electrical conductivity
The development of electrical conductivity in polymer composites occurs due to the rapid evolution of microstructural connectivity above a critical loading, commonly referred to as the percolation threshold, /*. This type of continuous
phase transition provides a powerful indirect method to infer nanoscale morphological characteristics based on the measurement of bulk transport properties, provided several limiting assumptions are fulfilled. If the conductive fillers behave as self-similar geometric objects that randomly overlap throughout the transition region, the electrical conductivity of the bulk material should follow the simple scaling relation: r(/) |//*|t, for / > /* [32], where / is defined here as volume fraction. This relation suggests that conductivity, r, depends only on the filler concentration above the percolation threshold, (//*), and the critical scaling exponent t. In an ideal system, the scaling exponent is independent of the specific details of the physical system and relates only to the dimension of the percolation network, which is given by t 2 in a 3-dimensional system [32]. The direct application of percolation theory to provide quantitative interpretation of network formation in polymer composites is somewhat questionable because (1) the model requires an infinitely conductive, randomly oriented, and uniformly dispersed phase be added to an infinitely resistive matrix, and (2) geometrical percolation does not account for excluded volume interactions. In other words, percolation theory is inadequate to describe any local orientation or phase segregation. However, there is some generality inherent in the model because it is based on the random overlap of geometric objects and is therefore applicable to any type of individual particle or clustered aggregate of a characteristic size. Agreement between experimental results and universal scaling behavior merely indicates that the conductive units, which may be individual particles or mesostructural clusters, exhibit self-similar morphology throughout the transition region [33]. Numerous factors influence the specific conduction mode and result in deviations from ideal behavior that complicate direct physical interpretation [12,34–36]. The electrical conductivity measured through the thickness of compression molded composite plaques at room temperature is shown in Fig. 7 as a function of volume fraction, /. The melt state conductivity data of PP/MWCNT composite melts previously reported by Kharchenko et al. [37] and Obrzut et al. [38] are included for comparison. Experimental data for the PP/F-MWCNT composite solid presented in this work and the PP/P-MWCNT melt reported in [37] were fit to the simple percolation relation: rð/Þ ¼ ro ð/ / Þt
Fig. 6 – DSC thermograms of the PP composite systems obtained during first cooling cycle at 10 C/min. (1) Neat PP; (2) PP/F-MWCNT (0.1 wt.%); (3) PP/F-MWCNT (0.6 wt.%); (4) PP/P-MWCNT (0.1 wt.%); (5) PP/ODA (0.5 wt.%).
ð1Þ
where ro is the effective conductivity of the filler material, and are shown as lines in Fig. 7. To observe if the scaling behavior is consistent with a purely geometric transition, the data and fitted curves are also shown as functions of concentration above the critical volume loading, (//*), in Fig. 7b. This approach is used to identify deviations from classical scaling behavior rather than to provide meaningful fitting factors. The percolation parameters were therefore determined under the assumption that both networks are ideal and follow universal scaling with t = 2. The effective conductivities that provided the best fit to the experimental data, with scaling exponent fixed at t = 2, were ro = 70 S m1 for the PP/FMWCNT composite presented here, and 220 S m1 from the melt-state measurements reported by Kharchenko et al. [37]. Based on the previous discussion that the electronic state
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Table 1 – Non-isothermal crystallization and melting parameters for PP composites.
Neat PP PP/F-MWCNT (0.1 wt.%) PP/F-MWCNT (0.6 wt.%) PP/P-MWCNT (0.1 wt.%) PP/ODA (0.5 wt.%)
Xc (%)
Toc (C)
Tc (C)
44 ± 0.5 44 ± 0.7 45 ± 0.2 44 ± 0.7 44 ± 0.5
122.6 123.6 125.7 122.5 116.9
118.5 120.8 122.6 118.0 114.2
Tm (C) 165.5 ± 0.1 164.5 ± 0.6 164.7 ± 0.5 166.9 ± 1.0 161.2 ± 0.3
Xc is degree of crystallinity, calculated from enthalpy of melting of second melt, Toc and Tc are crystallization onset temperature and peak crystallization temperature, respectively; Tm is peak melting temperature obtained from second melt. All measurements performed at rate of 10 C/min.
Fig. 7 – Characterization of connectivity transition for electrical conductivity of PP/MWCNT composites. (a) Electrical conductivity as a function of MWCNT volume loading for PP/MWCNT composites from this work (solid symbols) and melt state values from Kharchenko et al. [37] and Obrzut et al. [38] (open symbols); (b) Scaling in the percolation regime shown with conductivity plotted as function of loading above critical concentration, (//*).Lines are fit to Eq. (1) with slope fixed to universal scaling value in 3-dimensions (t = 2). of the F-MWCNTs is not significant decreased, the difference in ro is attributed to the higher conductivity of the PP amorphous phase [39]. The critical loadings were determined to be /* 0.24 vol.% (0.56 wt.%) in both systems. Volume fractions for the solid-state measurements in this work were calculated assuming the density of PP and MWCNTs are 0.90 and 2.1 g cm3, respectively. The volume fractions for the melt state systems were provided in their respective articles. The poorly exfoliated P-MWCNTs do not form a conductive network below 1 wt.%, which is in agreement with most literature reports [19,30,40,41]. The network of disentangled F-MWCNTs experiences a percolation-like transition at substantially lower concentration of 0.4 wt.% < / < 0.6 wt.%. Based on the empirical curve fitting of random multi-phase anisotropic materials in 3-dimensions provided by Garboczi et al. [42], the observed percolation threshold of 0.40.6 wt.% corresponds to an average aspect ratio of 240– 360. Given the waviness of the MWCNTs observed in the TEM images, the true aspect ratio should be greater than this projected value, which provides an indirect confirmation that the length distribution is not significantly affected by the disentanglement process. From Fig. 7b, it is evident that the concentration-dependent conductivity of the PP/F-MWCNT composite in the transition regime is consistent with expectations from geometric percolation theory for an isotropic network of high aspect ratio conductive particles. The network structure at the onset of percolation, i.e. at /* 0.24 vol.%, may therefore be regarded as self-similar throughout the range of concentration studied, which indicates that no significant aggregation occurs at
loadings of up to 2 wt.% F-MWCNT in PP. It is interesting to note that the results in this work are remarkably consistent with those of Kharchenko et al. [37], despite being done in completely different relaxation states. This agreement suggests that the well-dispersed system obtained in the meltstate, i.e. in the absence of entropic penalties to mixing, is preserved upon cooling due to the disentanglement and compatibilization of the F-MWCNTs.
3.4.
Thermal degradation
The thermal degradation of the PP/MWCNT composites was investigated using TGA. The residual mass as a function of temperature and the corresponding first derivative (DTG) curve are reported in Fig. 8a and b, respectively, for degradation in air. The influence of MWCNTs on the thermal stability of PP in both air and nitrogen environments is summarized in Table 2 using the temperature associated with 10% mass loss, T10%, and the temperature at maximum rate of mass loss, Tp. PP is highly susceptible to oxidative dehydrogenation at tertiary carbon atoms along its repeat unit structure and generally requires the incorporation of anti-oxidants or other stabilizers to minimize thermo-oxidative degradation during processing in air. The PP used in this work is a commercial grade that contains 0.15 wt.% anti-oxidant. This matrix was selected as a control to ensure that any change in physical properties is not due to the increased stability of PP in the presence of CNTs, which has been previously reported at high loading [43–46].
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Fig. 8 – Thermogravimetric degradation curves of PP/MWCNT composites in air at heating rate of 20 C/min; (a) Residual mass as function of temperature for PP/P-MWCNT and PP/F-MWCNT composites, (b) Corresponding DTG curve for mass loss rate as function of temperature for PP/F-MWCNT composites. Percentages refer to mass fraction.
Table 2 – Summary of thermal degradation parameters of neat PP and PP composites. Air
Neat PP 0.1 wt.% 0.5 wt.% 0.1 wt.% 1.0 wt.%
F-MWCNT F-MWCNT P-MWCNT P-MWCNT
N2
T10% (C)
Tp (C)
T10% (C)
Tp (C)
370 400 415 385 392
464 473 481 471 464
436 444 465
483 493 501
T10%, decomposition temperature at 10% mass loss; Tp, temperature at peak mass loss rate.
The thermal degradation of the unfilled PP in air occurs in a single-step process that begins at 320 C, accelerates with temperature until maximum decomposition rate is reached at Tp = 444 C, and terminates with no significant residual mass around 475 C. The addition of 0.1 wt.% F-MWCNT substantially increases the onset temperature of thermal degradation and shifts T10% about 30 C higher than the neat PP. Increasing loading to 0.5 wt.% F-MWCNT enhances both measurements of thermal stability. Control specimens prepared with P-MWCNTs in PP exhibit much smaller increase in stability, with T10% only 20 C higher than the neat PP at 1 wt.% loading. It should be noted that although MWCNTs and other carbon-based materials with sp2 bond architecture are known to effectively quench free radicals due to their high electron affinity [47,48], improvement in thermal stability with P-MWCNTs might be partially due to the presence of residual catalyst particles, such as paramagnetic iron [45,49]. The presence of the iron catalyst was verified by the presence of red, magnetic particle residues after testing. The MWCNT disentanglement procedure used in this work removes the majority of impurities and catalyst residues, which was confirmed by a lack of residue after thermal decomposition above 600 C in air. This indicates that the mechanism responsible for the observed thermal stability of PP/F-MWCNT is distinct from secondary effects associated with P-MWCNTs. In order to assess the mechanism(s) responsible for the thermal stability of the nanostructured composites, it is useful to compare the first derivative TGA curves (DTG) shown in Fig. 8b. The curve shape of the PP/F-MWCNT composites are similar to the neat PP, but exhibit substantially lower rate of mass loss during the initial degradation process near 300 C,
and an increased temperature and magnitude of peak mass loss rate. Increasing the concentration of F-MWCNTs enhances the effect, which suggests that the behavior is related to the total surface area between the polymer matrix and the MWCNTs. These findings indicate that although the primary mechanism of thermo-oxidative chain degradation is not altered due to the presence of the F-MWCNTs, the kinetics of the process are significantly reduced. Improved thermal stability of PP in the presence of nanofillers is typically attributed to suppressed diffusion of degraded products to the bulk polymer surface, which thereby inhibits further mass loss [50,51]. The low MWCNT loading reported here is below /*, which implies that there is no physical network to interact with the volatile component at the initial stage of mass loss. Furthermore, as the PP degrades, the relative concentration of MWCNTs should increase and enhance the barrier effect, leading to greater stability improvement with increasing temperature, which is not observed here. The apparent correlation between nanoparticle surface area and thermal stability suggests the primary mechanism for improved stability in air is either physical adsorption of the PP chains at the MWCNT surface, or quenching of free radicals generated by the degrading polymer. In the case of physical adsorption, significant improvement in stability would be anticipated for either air or N2 environments. However, control experiments performed with N2 purge revealed that although there is some improvement in stability due to the presence of F-MWCNTs (Table 2), the effect is much smaller than in air. Furthermore, previous reports have shown that the dynamics of PP chains are not significantly affected by the MWCNTs [37], which suggests that the electron rich
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sp2-architecture of the CNTs is likely the primary cause for the improvement in thermo-oxidative stability. Fenoglio et al. [52] reported that MWCNTs exhibit a remarkable scavenging ability in the presence of hydroxyl or superoxide radicals. The disentanglement and dispersion of the F-MWCNTs significantly increases surface area, which enhances the radical scavenging effect results in improved degradation stability at much lower MWCNT concentration than previously reported. Poor compatibility likely results in segregation of the nanotubes to interspherulitic boundaries and amorphous domains, which diminishes the fraction of PP matrix affected by the CNTs. Rheological characterization is being pursued to provide further insight into the degree of physical interaction in the system.
4.
Discussion
Injection moldable conductive thermoplastics are in great demand as housing materials to protect electronic components from electrostatic discharge (ESD). Conventional spherical conductive particles, such as carbon black (CB), typically compromise material performance and processability due to the high volume fraction (/ 15%) needed to introduce acceptable levels of conductivity. CNTs are among the most attractive materials for ESD applications because they possess both large aspect ratio and exceptional electrical conductivity. In this work, we have provided extensive evidence to show that nanoplatelet-assisted dispersion can be used to uniformly disentangle MWCNTs without significant structural or electronic degradation. The disentangled MWCNTs are stable in aqueous solution and remain homogeneously dispersed in non-polar solvent after reacting with ODA. The compatibilization of the MWCNTs with an organophilic functionalizing agent provides morphological stability upon mixing and does not significantly influence crystallization behavior. The approach presented in this work provides a high-yield and potentially scalable route for the commercial production of CNT-filled conductive thermoplastics [13], and is sufficiently stable for direct melt mixing, which will be reported separately. The focus in this discussion will be to provide insight into the role of individual MWCNTs on the electrical conductivity and thermal stability of the bulk nanostructured polymer composites. Sharp divergence in the transport properties of a CNTfilled polymer occurs if the average distance between CNTs is small enough to allow electrons to tunnel or hop between conductive particles that are physically separated by the insulating matrix [53]. The critical loading where this transition occurs is determined by the dispersion and alignment of MWCNTs, as well as the potential barrier of the matrix, which controls the effective tunneling radius [54]. The lowest /* in an insulating PP matrix that we are aware of was reported in the melt state by Obrzut et al. [38], which is reproduced in Fig. 7a for comparison. In their work, /* was found to depend on shear rate and reduced a minimum value of /* 0.12 vol.% MWCNT as the shear rate approach 0 s1. This shear-induced ‘‘reverse-percolation’’ transition was attributed to the randomization of MWCNT alignment of the shear plane as the shear rate decreased below the reciprocal of the structural relaxation time of the melt matrix, s 1 s [37,38]. The
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clear disagreement between their finding and our claim of uniform, random dispersion can be rectified by comparison with similar shear-induced transitions in PC/MWCNT and epoxy/MWCNT systems [55,56]. In these works, at a given MWCNT loading below /* of a well-dispersed network, conductivity increases only in the presence of large, visibly aggregated MWCNT clusters that form at low shear rates. The deviation between melt state measurements reported by Kharchenko et al. [37] and Obrzut et al. [38] may be attributed to shear-induced aggregation rather than any degree of improved dispersion or randomized alignment. This claim is supported by the evident continuity between data sets in Fig. 7a. It is therefore reasonable to claim that the behavior reported by Kharchenko et al. [37] reflects nearly ideal behavior in the melt state, and that the similarity between their measurements and ours suggests that the ideal structure is preserved upon cooling to the solid state. In this case, our observed percolation threshold should be nearly equal to the geometric transition threshold, which is the lowest achievable value for a conductive network of random, individual particles with a given aspect ratio distribution. In a given polymer matrix, a network of monodisperse, randomly oriented, individual cylinders, /* is a geometric value that is single-valued and depends only on aspect ratio [42].The threshold will shift to higher concentration due to aggregation, phase separation, and alignment, or reduce if extended, ramified structures are formed. Lowering /* by shearinduced aggregated network formation is frequently reported in epoxies and several other amorphous matrices [55–57], and may be responsible for the unique behavior reported by Obrzut et al. in PP [38] (Fig. 7a). However, this mechanism does not occur in solid-state PP [39,57], which suggests that the crystallization of matrix determines the final morphology of the network. At low concentration, entangled and poorly dispersed CNTs are not effectively incorporated within the crystalline structure, which inhibits the formation of an interconnected network. At higher concentration, the rejection of clustered CNTs to interspherulitic and amorphous regions favors the formation of much higher conductivity networks than the random, well-dispersed system presented in this work. The disentanglement and uniform dispersion of MWCNTs appears to be most effective at enhancing properties related to the interfacial area between the filler materials and the host matrix, which is demonstrated in this work through the effectiveness of F-MWCNTs in nucleating PP crystallite formation, quenching free radicals during thermal degradation in air, and establishing a conductive network at low concentration. The conductivity at loadings above the percolation threshold is lower than previous reports [19,30,37,38,40,41], which indicates that the physical isolation of the individual MWCNTs may not be desirable for applications requiring high electrical conductivity. One approach to improve conductivity at higher loadings is to confine dense cellular networks of CNTs to boundaries between discrete polymer domains [52]. However, in thermoplastics, such an approach will involve a compromise in mechanical properties and cannot be easily reprocessed without altering the network structure. Another approach is to select a polymer matrix with higher conductivity, thereby decreasing the tunneling resistance between CNTs at a given interparticle separation (c.f. [54]), and enabling low
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5 0 ( 2 0 1 2 ) 4 7 1 1 –4 7 2 1
percolation threshold and low tunneling resistance without significant phase separation [58]. This approach is compelling to provide further electrical conductivity improvement in ESDgrade conductive thermoplastics, but has limited applicability for highly insulating polymers, such as PP. The approach in this work does not rely on the introduction of low molecular weight or chemically modified polymeric species, and is not limited to solution-based processing approaches. This method should therefore be suitable for the high-volume production of conductive thermoplastics without compromising mechanical properties, and may potentially be adapted to other polymer matrices by proper MWCNT surface functionalization.
5.
Conclusions
The morphology, crystallization, electrical conductivity, and thermal stability of PP composites containing well-dispersed, disentangled F-MWCNTs were investigated in the work. The disentanglement process enhances the crystallization behavior at a given loading due to an increase in the number of active nucleating sites. The process also reduces the electrical percolation concentration to near theoretical minimum predictions for a randomly dispersed network of high aspect ratio conductive rods. The observation of universal scaling supports morphological evidence that dispersion and exfoliation are preserved up to 2 wt.%, and indicates that the F-MWCNTs maintain an isotropic conductive network in the solid state. The maximum achievable electrical conductivity is limited by the physical separation of individual MWCNTs within the insulating PP matrix. The uniform dispersion and disentanglement result in a substantial increase in thermal stability at extremely low loading due to the large interface between the PP chains and the free radical scavenging MWCNTs. This method is suitable for the production of conductive thermoplastics with high volume processing techniques. With proper selection of functionalizing surfactants, this approach may be modified to develop low cost, high performance polymer composites for engineering applications with a wide variety of polymer matrices.
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Acknowledgements The authors would like to acknowledge Japan Polypropylene Corporation and KANEKA Corporation for financial support throughout the duration of this research.
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Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.carbon. 2012.05.063.
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