Influence of Chain Branching and Molecular Weight on Melt Rheology

Article pubs.acs.org/Macromolecules

Influence of Chain Branching and Molecular Weight on Melt Rheology and Crystallization of Polyethylene/Carbon Nanotube Nanocomposites Juan Francisco Vega,† Yudith da Silva,†,‡ Ernesto Vicente-Alique,† Rafael Núñez-Ramírez,† Mariselis Trujillo,‡ María Luisa Arnal,‡ Alejandro J. Müller,*,‡,§,∥ Philippe Dubois,⊥ and Javier Martínez-Salazar*,† †

Biophym, Departamento de Física Macromolecular, Instituto de Estructura de la Materia, CSIC, C/Serrano 113 bis, 28006 Madrid, Spain ‡ Grupo de Polímeros USB, Departamentos de Mecánica y Ciencia de los Materiales, Universidad Simón Bolívar, Apartado 89000, Caracas 1080-A, Venezuela § Institute for Polymer Materials (POLYMAT) and Polymer Science and Technology Department, Faculty of Chemistry, University of the Basque Country (UPV/EHU), Paseo Manuel de Lardizabal 3, 20018 Donostia-San Sebastián, Spain ∥ IKERBASQUE, Basque Foundation for Science, E-48011 Bilbao, Spain ⊥ Service des Matériaux Polymères et Composites SMPC, Center of Research and Innovation in Materials & Polymers CIRMAP, University of Mons, Place du Parc 20, B-7000 Mons, Belgium ABSTRACT: In this paper, the synergistic effects that carbon nanotubes (CNTs) produce on the basic rheological properties and crystallization of polyethylenes with different branch contents and molecular weights was investigated. Multiwalled carbon nanotubes coated with polyethylene (as produced by in situ polymerization) were blended in the melt (in a 1% wt. ratio) with three polyethylene matrices of different molecular weights and branch contents. Transmission electron micrographs demonstrated excellent carbon nanotube dispersion in all samples and the existence of a geometrical percolation network. The rheological and calorimetric properties of the nanocomposites were determined and the results compared to those obtained for neat polyethylene resins. Both Newtonian viscosity and steady-state shear recoverable compliance increased with the addition of CNTs in all cases. However, the increase was strongly dependent on the molecular weight (and dispersity index) of the matrices regardless of the branch content. A novel screening effect of the CNTs network due to the high relaxation times of the matrix with the highest molecular weight was detected. This important result demonstrates that viscoelasticity can hinder the measurement of the rheological percolation threshold of CNTs network depending on the scale of relaxation times involved. Additionally, it was found that in relative terms (comparing each nanocomposite with its neat polyethylene matrix), the Mw values also play a vital role in CNT nucleation besides chain branching content. Both nonisothermal and isothermal nucleation effects caused by CNTs increased as the Mw of the polyethylene matrix decreased in spite of the role played by short chain branches in decelerating their overall crystallization kinetics. The capability for producing more stable lamellae through successive annealing of the nanocomposites as compared to their neat matrices also followed a decreasing trend with molecular weight increases, as indicated by SSA thermal fractionation results. Nevertheless, the presence of branches played a major role, since fractionation quality improved greatly as the branch content increased in the samples, as expected on the basis of the sensitivity of thermal fractionation to the presence of defects along crystallizable sequences.

I. INTRODUCTION

key factors required for improving the properties of the nanocomposites. In general, CNTs show a strong tendency to aggregate; hence dispersion becomes a critical issue during nanocomposite

During the last few decades, nanotechnology has been guiding the design of new materials with unique physical properties. In this respect, carbon nanotubes (CNTs) are being extensively used as fillers in polymeric matrices due to their high specific surface and excellent mechanical properties.1,2 It has been widely recognized that both the quality of dispersion and interfacial interaction of CNTs with the polymeric matrix are © XXXX American Chemical Society

Received: June 19, 2014 Revised: July 21, 2014

A

dx.doi.org/10.1021/ma501269g | Macromolecules XXXX, XXX, XXX−XXX

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Table 1. Molecular and Melt Linear Viscoelastic Features (T = 140 °C) of the Neat Polyethylene Matrices and the Nanocomposites Studieda Je0 (mPa‑1)

η0 (kPa s) sample

SCB CH3/1000C

Mw kg mol

PE1 PE2 PE3

10.7 0 4.4

113 185 326

‑1

‑1

Mw/Mn

Mz kg mol

neat

CNT

neat

CNT

2.1 2.2 37.4

225 370 1990

14.0 60.0 1020

900 425 1460

0.065 0.64 2.30

3.32 4.95 2.96

SCB: Short chain branching content. Mw: weight-average molecular weight Mz: z-average molecular weight. η0: zero shear melt viscosity. Je0: linear steady-state recoverable compliance.

a

preparation. In situ ethylene polymerization on the CNT surface has been proven to overcome this difficulty in the case of polyethylene (PE)/CNT nanocomposites, as PE chains grow directly on the surface of the nanotubes, which was previously treated with a single-site catalytic system, allowing the breakup of the native CNT bundles and leading, upon further melt blending with polyolefinic matrices, to high-performance nanocomposites.3 In a well-dispersed system, the eventual differences in the properties of polymeric nanocomposites should be attributed to the nature of the interactions between the components of the system.4 Of special interest are melt viscoelasticity and processing, mechanical properties and the nanostructure and crystallization behavior of polymeric nanocomposites. The influence of CNT on the viscoelastic properties of molten polymers has been a subject of great interest in recent literature.5−12 The presence of fillers of any type in polymeric matrices leads to a well-known hydrodynamic effect of reinforcement at low concentrations.13 In this regime, no particle−particle interactions occur since the nature of the interface between the particles and the matrix is the main reason for the changes in properties. At a given concentration, i.e., the percolation threshold, the fillers form a network, which is responsible for the appearance of yield-points and solid-like behavior.14 The transition from liquid-like to solid-like behavior can be observed at lower filler concentration as the quality of the dispersion is improved, the interaction between the filler and the matrix is stronger or the filler aspect ratio and alignment is larger. Hence, different values of the percolation threshold can be found in the literature for several types of nanocomposites based on CNTs.2,7,15−19 Percolation threshold values between 0.05 wt % to 15.0 wt %, have been reported in the case of PE/CNT and polypropylene (PP)/CNT systems.1,6,16,18,20,21 Most of the reported values for the percolation threshold are within the range of 2−3 wt %. In general, all these works have described high increases in viscosity, dynamic storage shear moduli and steady shear compliance with CNT content, but little attention has been paid to the effect of the molecular architecture of the matrix. Zhang et al. have reported a decrease in viscosity and moduli for a series of ultrahigh molecular weight polyethylenes (UHMWPE) with broad molecular weight distribution (MWD) upon single wall CNT addition, in the range of compositions 0.1−1 wt %.20 This decrease was not observed for different nanocomposites based on UHMWPE with narrow MWD or conventional high density polyethylene (HDPE).7,18 In our previous work, we have found a similar behavior for bimodal PE/CNT nanocomposites with high-molecular weight tails.10 Zhang et al. explained their findings as a consequence of the selective adsorption of the longest molecules of the distribution onto the CNTs surface,20 an idea first proposed by

Maurer et al. in the 80s for polydisperse UHMWPE on silica particles.22 In the context of polymer nanocomposites with silica nanoparticles, Trievel et al. have recently found interesting differences between the observed increases of the steady-state compliance values in PS/silica nanocomposites depending on the MWD of the matrix, for nanoparticle volume fractions as low as 1 vol %.23 These authors ascribed the enhanced elastic character of a nanocomposite prepared with a very narrow MWD PS matrix to a broader relaxation time distribution, as compared to a nanocomposite with a broader MWD PS matrix. However, they also mentioned the elusive physical nature of these long relaxation times that are probably induced by the interaction between the molecules and the nanoparticles. In relation to physical properties, CNTs can induce exceptional improvements when effectively dispersed in polymeric matrices, such as elastic modulus increases and spectacular electrical conductivity enhancements.2 These improved properties are also connected with the large specific surface of the CNTs, which favors both the interaction between the polymer matrix and the CNTs and the establishment of a percolation path among them.1−3,24,25 The effects of CNTs on the nucleation, crystallization, and morphology of semicrystalline matrices has been extensively studied in the past few years as reflected in the comprehensive recent review by Laird and Li.26 The particular case of PE/ CNTs has been the subject of several studies. As regards to morphology, it has been demonstrated that nanohybrid shishkebabs or similar structures (row nucleated or bottle-brush morphologies) can be formed from solution crystallization, thin films and bulk crystallization, with a size-dependent soft epitaxy, where experiments have shown that chains orient preferentially parallel to the CNT axis.9,27−35 The preferential perpendicular lamellar growth with respect to the CNT axis has also been confirmed by computer simulation studies.36−40 CNTs can affect the nucleation of polymeric materials in four different ways:26 (a) producing nucleation effects, (b) causing antinucleation effects, (c) not affecting the crystallization, and (d) producing supernucleation effects (a phenomenon first identified by some of us).34,41−43 In the case of PE/CNT nanocomposites, the literature reports mostly cases a and d.5,7,10,18,30,31,34,35,42,44−47 Since CNTs generally induced nucleation or supernucleation effects on PE matrices, the overall crystallization rate is usually accelerated when CNT contents are kept below 10%. At higher CNT loadings, topological confinement effects may be observed that can eventually lead to reductions in nucleation and crystallization kinetics and crystallinity values.34,35,42 Another interesting effect is that CNTs can induce the formation of more stable PE lamellar crystals (i.e., thicker lamellae) when prolonged isothermal crystallization is performed.34 Recent computer simulation studies performed B


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mixed in a Haake MiniLab II Rheomex mini-extruder (model CTW5) at 180 °C and 60 rpm with 2 wt % of an in situ polymerized PE-CNT masterbatch. The mixing time was 6 min. The masterbatch composition was 43 wt % PE, 52 wt % multiwall carbon nanotubes (MWCNT) and 5 wt % oxidized catalytic residues (mainly alumina). The masterbatch is denoted PE43M52A5, where the subscripts indicate the composition in wt %. The nanocomposites prepared in this work (PEx/2 wt % PE43M52A5) had a final content of 1.04 wt % of MWCNT. The preparation and characterization of the masterbatch has been described in detail elsewhere, and the diameters of the CNT were found by TEM to be 5−27 nm.34,35 The in situ polymerization process has been designed to produce an optimum dispersion since the PE chains grow directly on the surface of the CNTs, allowing for the breakup of the native bundles leading, upon further melt blending with PE, to high-performance polyolefinic nanocomposites.3 Sheet samples for neat materials and nanocomposites for rheological measurements with 1 mm thickness were compression-molded with a Schwabenthan Polystat 200T hot press at a temperature of 160 °C and a pressure of 25−50 bar during 5 min and then cooled to room temperature. II.3. Transmission Electron Microscopy. To observe the morphology of the nanocomposites, carbon replicas of the surfaces of clean and flat films were prepared. The films were obtained by hotmelt pressing the nanocomposites at 160 °C during 5 min between flat surfaces of high melting temperature polyester films under a nominal pressure of 25 bar, which is low enough to avoid any orientation effects. For comparison purposes replicas of the neat PE3 sample were also obtained. After cleaning the surface of the films, the specimens were placed in a coating chamber fitted with a carbon evaporation source. The specimens were imprinted with an amorphous carbon film after shadowing at an incident angle of 30° to obtain a replica of the surface topography. A small amount of poly(acrylic acid)solution was deposited onto the coated specimen and dried at room temperature. After drying, the polymer sample was detached from the poly(acrylic acid)which remained glued to the carbon replica. Then, the poly(acrylic acid)layer was dissolved in water, and the obtained carbon replicas mounted on grids for transmission electron microscopy (TEM) observation in bright field mode using a JEOL JEM-2100 microscope operated at 200 kV. II.4. Torsion Rheometry. Melt rheological measurements were carried out using a stress-controlled Bohlin CVO rheometer (Malvern Instruments), in both dynamic and static modes. Dynamic oscillatory shear measurements were performed at a temperature of 140 °C. The applied shear stress amplitudes corresponded to a shear strain of 0.05, which was proven to be in the linear viscoelastic regime for both the neat materials and the nanocomposites. The properties measured were the storage and the loss moduli, G′(ω) and G″(ω), as well as the modulus of the complex viscosity, |η*(ω)|. Creep and creep−recovery experiments were also performed at 140 °C. It was tested that the range of shear stress was within the linear regime (located below τ0 = 50 Pa in all the materials studied). The time-dependent creep and recovery deformation was measured. From these results the creep, Jc, and recovery, Jr, compliances were obtained, which allowed one to extract the values of the zero-shear rate viscosity, η0, and the steadystate recoverable compliance, Je0. Because these experiments take a long time especially for the neat samples with the highest Mw and broadest MWD and also for the nanocomposites, thermal stability of the systems must be guaranteed. In order to check the thermal stability of the samples, a comparison was made between creep compliance measurements performed before and after the recovery experiments. The results were quantitatively comparable demonstrating that the samples were thermally stable at 140 °C for periods of time of more than 24 h. We have converted the time-dependent compliance data to frequency dependent moduli in order to extend the master curves from dynamic measurements to lower frequencies. The reader is referred to previous works for details of this methodology.55−57 II.5. Thermal Properties. A PerkinElmer differential scanning calorimeter (DSC) Pyris 1 instrument calibrated with indium and tin under an ultra high purity nitrogen atmosphere was employed. Samples (with a weight of 5 mg) were encapsulated in aluminum pans

by us were able to reproduce many of the above experimental facts on linear and branched PE/CNT nanocomposites.40 The simultaneous effect of chain branching and molecular weight on the rheology and crystallization behavior of PE/CNT nanocomposites has been seldom studied. Therefore, the aim of the present study is to investigate the way in which the properties of well dispersed nanocomposites are affected by the molecular architecture of the matrix. These effects will be explored in this work, by studying three well-dispersed PE/ CNT nanocomposites where the PE matrices contain different branching degrees and molecular weights.

II. EXPERIMENTAL SECTION II.1. Materials. In order to investigate the influence of the molecular features of the matrix in the properties of PE/CNT nanocomposites, materials of the same chemical composition have to be selected. Three PE samples were chosen as matrices. Two of them (PE1 and PE2) are noncommercial samples with narrow MWD synthesized employing single-site metallocene catalysts and provided by Repsol (Spain). However, PE1 is a random copolymer of ethylene and 1-hexene while PE2 is a linear polyethylene material (see Table 1). Their molecular features, physical properties and processing behavior have been deeply investigated in previous works from the experimental48−50 and computer simulations points of views.51−53 The other sample (PE3) is a commercial ethylene/1-butene copolymer with a broad MWD and a small short chain branching (SCB) content. The molecular parameters were determined by size exclusion chromatography (SEC) using IR and viscometric detectors and 13C NMR spectroscopy. Please refer to our previous works referred above for specific details. The SEC profiles of the samples are provided in Figure 1.

Figure 1. Size exclusion chromatography traces of the PE samples studied. SEC results show a variation in the weight-average molecular weight, Mw, of the samples. Special attention should be paid to sample PE3. As it can be observed in Figure 1 the MWD is bimodal, with an important contribution of very high-Mw species (Mw > 1000 kg/mol). The molecular properties of the samples are listed in Table 1. The three samples were chosen with differences in both Mw and SCB. The reason was that in the case of rheological measurements it is expected that SCB will not significantly affect the measurements in the content range and temperature explored while the Mw effects will dominate the behavior.48−53 Conversely, the Mw range explored should not produce a significant effect on the crystallization behavior since all samples are well above the entanglement limit and the maximum Mw used cannot be classified as predominantly ultra highMw.54 However, SCB differences will cause very important effects on the crystallization kinetics and melting points of the samples.54 II.2. Nanocomposites and Sample Preparation. In order to prepare the nanocomposites, the described PE samples were melt C


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and sealed. The crystalline thermal history was erased by heating the samples at 170 °C for 3 min. Nonisothermal experiments were performed by cooling and subsequent heating the samples at 10 °C/ min. Self-Nucleation Tests (SN). Self-nucleation tests (SN) were performed according to standard procedures devised originally by Fillon et al.58−60 The following steps were followed: (a) Heating to 170 °C for 3 min in order to erase crystalline history. (b) Cooling at 10 °C/min to 25 °C to create a “standard” thermal history by controlled crystallization of the sample. (c) Heating to a temperature denoted Ts, depending on which, the sample will be completely molten, just self-nucleated or selfnucleated and annealed. If Ts is too high, then the molten sample is said to be in “Domain I” or complete melting domain where only temperature resistant heterogeneities remain to cause nucleation upon cooling. If Ts is high enough to melt almost completely the sample but low enough to induce selfnuclei without any crystal annealing, the sample is said to be under “Domain II” or self-nucleation domain (for details on the controversial nature of Domain II; see ref 58). If Ts is low enough, only part of the crystal population will melt, and therefore the unmelted crystals will anneal during the 5 min at Ts (step d). Upon subsequent cooling, the molten polymer will be self-nucleated by the annealed crystal fragments. Then the sample is said to be in “Domain III” or self-nucleation and annealing domain. (d) Thermal conditioning at Ts during 5 min. (e) Cooling at 10 °C/min from Ts to 25 °C, where the effects of the thermal treatment will be reflected on the crystallization of the PE. (f) Heating at 10 °C/min from 25 to 170 °C, where the effects of the thermal treatment will be reflected on the melting of the PE. Successive Self-Nucleation and Annealing Experiments (SSA). The SSA technique was developed and implemented by Müller et al.60−65 It can be considered a thermal fractionation technique that promotes the molecular fractionation process that occurs during crystallization, while encouraging annealing of the unmolten crystals at each stage of the process, so that small effects can be magnified. The complete thermal conditioning comprises: (a) Heating to 170 °C for 3 min in order to erase crystalline history. (b) Cooling at 10 °C/min to 25 °C to create a “standard” thermal history by controlled crystallization of the sample. (c) Heating at 10 °C/min from 25 °C to Ts. The first Ts to be employed is normally the ideal self-nucleation temperature (see refs 60 and 63). In the case that a series of samples is to be compared among them with identical SSA treatments, then the highest ideal Ts must be employed. (d) Isothermal step for 5 min at Ts. (e) Cooling at 10 °C/min from Ts to 25 °C, where the effects of the thermal treatment will be reflected on the crystallization of the PE. (f) Steps c through e are repeated for the number of fractionation steps desired. (g) Final melting: The sample was heated at 10 °C/min from 0 to 160 °C and a multiple melting endotherm was obtained. In this work, a Ts temperature of 131 °C was chosen as the starting fractionation temperature or first Ts. Steps “c” to “e” of the SSA method described above were repeated at increasingly lower Ts. The fractionation window or difference in Ts values was always kept constant at 5 °C. The SSA procedure was performed for all the samples, in a Ts range of 131 to 81 °C for a total of eight selfnucleation/annealing steps. Isothermal Crystallization Kinetics. We closely followed the procedure and guidelines given by Lorenzo et al.66 in order to perform the isothermal crystallization experiments. In order to erase all crystalline thermal history, the samples were first heated to 170 °C and

kept in the melt for 3 min. Then, they were quenched under a controlled cooling rate (at 60 °C/min) to the chosen isothermal crystallization temperature, Tc. In order to choose the Tc range, tests were run by quenching the sample to a specific Tc and then immediately heating it while recording its heating scan at 10 °C/min. If any melting was observed, it indicated that the sample was able to crystallize during the previous quenching at 60 °C/min. Therefore, this Tc was not employed and the test was repeated at higher temperatures until no crystallization during quenching was obtained. This procedure was used to find the lowest Tc value to be employed, then isothermal crystallization experiments were run every half a degree from that Tc upward in temperature until the crystallization exotherm was too small to be detected by the DSC. For the temperature range evaluated (118 to 135 °C), the samples did not crystallize during the previous cooling. Finally, the samples were crystallized isothermally at the different chosen Tc values.

III. RESULTS III.1. Morphology of the Nanocomposites. Figure 2 shows the micrographs of the surface of replicas taken by TEM for PE1/CNT, PE2/CNT, and PE3/CNT samples (panel c includes an inset with the TEM image of the replica of neat PE3 surface for comparison purposes). The morphology of the nanocomposites is clearly observed in the micrographs showing large-scale dispersion of CNTs, as it is indicated by the absence of agglomerates. Moreover the images do not reveal significant differences between the matrices, and then it is guaranteed that a similar degree of dispersion was achieved in all cases. Additionally, there exists no preferential alignment of the CNTs due to the preparation procedure of the films. Figure 3 shows the case of PE3/CNT nanocomposite at a lower magnification, encompassing a wider sample surface field of the order of a few microns. In this picture, the tridimensional CNT network is clearly observed, but most interestingly the flexible nature of the CNTs is also revealed. From the TEM images (Figures 2 and 3) the dimensions of the CNTs can be evaluated. Their lengths are between 0.5 and 1.5 μm and their diameters range from a few nanometers to around 30 nm, in agreement with previous observations made in similar nanocomposites.34,35 The images also suggest that contacts or “entanglements” between individual CNTs exist in the nanocomposites, and that the percolation threshold has been exceeded in the three samples studied. III.2. Linear Viscoelastic Response in the Melt of the Materials Studied. III.2.1. Pure Materials. The shear compliance, Jc, obtained from creep measurements (Figure 4a) together with the storage modulus, G′, obtained from dynamic measurements (Figure 4b) follow the expected trends for the three neat PE samples. The results for Jc are within the linear viscoelastic region, as it is guaranteed by the identical trend obtained for each material irrespective of the shear stress applied in the range τ0 = 6.25−25 Pa (12.5−50 Pa for PE3). A higher viscosity (lower values of Jc) and elasticity (higher values of G′) are observed as the Mw and Mw/Mn of the sample increase from PE1 to PE3. Also, a broader linear viscoelastic response was obtained in the case of the PE3 sample in view of its broad MWD. For the lowest angular frequency investigated in the dynamic mode (Figure 4b), the terminal region (G′∝ ω2) is not entirely reached for the neat materials. This is a common result observed in polydisperse PE samples for which very low frequencies, sometimes unreachable, are necessary to attain the terminal region. However, the creep mode allows working in the terminal region providing the application of low enough values of shear stress for long creep times. From these D


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Figure 3. TEM micrograph of a carbon replica of the surface of the PE3/CNT nanocomposite at a lower magnification than that shown in Figure 2

Figure 2. TEM micrographs of carbon replicas of the surface of the nanocomposites studied: (a) PE1/CNT; (b) PE2/CNT; (c) PE3/ CNT. The inset in part c shows the carbon replica of the surface of neat PE3 (the scale bar in this case is also 200 nm).

Figure 4. (a) Shear creep compliance, Jc, as a function of creep time of the neat materials at T = 140 °C. Applied shear stress, τ0: (○) 6.25 Pa, (□) 12.5 Pa and (Δ) 25 Pa; (∇) 50 Pa. Dotted lines indicate the characteristic slope of the flow region. (b) Dynamic modulus, G′, versus angular frequency, ω, of the neat materials at T = 140 °C. Applied strain γ = 0.05 (3 independent measurements for each sample are plotted): (red ○) PE1, (blue △) PE2, and (□) PE3.

results, it is possible to obtain the value of η0, as η0 = lim t/J(t). t →∞

The expected increase with Mw is obtained as judged by the results listed in Table 1. Additionally, from the creep-recovery tests we have obtained Je0 for the three PE samples, as J0e = lim Jr(tc,tr). tc →∞ tr →∞

Figure 5a shows examples of the elastic recoverable compliance for PE1 and PE3. In the particular case of PE3, Je increases with tc up to a certain strain level at which the maximum of molecular orientation is achieved in the creep experiment.67 A constant steady-state value, Je0, is reached for tc values of at least 3125 s for PE2 and 12 500 s for PE3 (see Figure 5b). In contrast, the PE1 sample reaches the steady state value for a relatively short recovery time. The value of Je0

Increased values of the creep time, tc, as Mw and Mw/Mn of the samples increase were necessary to obtain reliable values of Je0. At very long recovery times, the elastic recoverable compliance, Jr, reaches a plateau, which is Je. The results indicated that in the case of the highest Mw and Mw/Mn samples, PE2 and PE3 respectively, the values of Je are strongly dependent on the creep time, tc. E


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Figure 5. (a) Recovery shear compliance, Jr, versus recovery time at T = 140 °C as a function of creep time for PE1 (solid symbols) and PE3 (open symbols). (b) Equilibrium shear recovery compliance, Je, versus creep time, tc: (red ●) PE1, (blue ▲) PE2, and (■) PE3. The inset shows the normalized Je versus tc plot.

obtained in this sample is in very good agreement with that reported for similar materials (similar Mw and MWD).68 The observed values of Je0 for PE2 and PE3 samples are of one and 2 orders of magnitude higher the obtained for PE1, respectively. This trend has been also widely observed in the literature, as Je0 values increase even due to small differences in MWD.69 The values obtained for Je0 are listed in Table 1. The inset in Figure 5b demonstrates that dependence of Je on tc is similar for the three neat materials, in agreement also with previous results from the literature.68 This means that the molecular architecture does not play a significant role in the effect of the creep time on the measured recoverable compliance. However, the differences in the values of the minimum creep time, tc0, and in the level of Je0 are very important among the three samples and should be attributed to the differences in Mw and MWD. III.2.2. Nanocomposites. For the three matrices the addition of the CNTs becomes evident by a decrease of Jc and an increase of G′ over the whole ranges of time and angular frequency, as it can be observed in Figure 6, parts a and b, respectively. The viscoelastic response obtained is clearly dependent on the polymer matrix. Especially interesting are the results obtained for PE1/CNT and PE2/CNT nanocomposites, for which the viscous flow region is delayed to very long times in the creep mode and to low frequencies in the dynamic mode. As a result of the delayed dynamics, a clear “plateau” develops for PE1/CNT nanocomposite and a shoulder appears in the case of PE2/CNT nanocomposite at long times (Figure 6a) and low frequencies (Figure 6b). However, the delayed dynamics seems not to be significant in the case of PE3/CNT sample. It is important to note that the results for Jc are also within the linear viscoelastic region in the same range of stress than in the case of the neat materials. For the values of τ0 applied (τ0 = 6.25−50 Pa) the observed Jc values in PE/CNT samples directly lead to values of the equilibrium shear rates of the order of 10−5 s−1 or lower and consequently to total values of the strain γ< 1−10. The values of Jc in Figure 6a have been converted to G′, G″, and |η*|, following the procedures described elsewhere,55−57 in order to obtain a complete linear viscoelastic fingerprint in the dynamic mode. The results obtained from the experiments can

Figure 6. (a) Shear creep compliance, Jc, as a function of creep time for the neat materials (open symbols) and the nanocomposites (lines) at a shear stress of τ0 = 6.25−50 Pa and T = 140 °C; (b) Dynamic modulus, G′, versus angular frequency, ω, of the neat materials (open symbols) and the nanocomposites (close symbols and lines) at T = 140 °C (three independent measurements for each sample are plotted). (red ●, ○) PE1, (blue ▲, △) PE2 and (■, □) PE3.

then be compared in a broad range of angular frequencies, as it is observed in Figure 7 for G′ and G″ and in Figure 8 for |η*|. The strong effect of the CNTs can be seen, mainly in the nanocomposites with lower Mw and narrower MWD matrices. The effect is again noticeable in the case of PE1/CNT and PE2/CNT nanocomposites, for which clear enhancements of the elastic character and of |η*| are noticed at low frequencies. Especially interesting is again the PE1/CNT nanocomposite for which an increase in η0 of 2 orders of magnitude with respect to neat PE1 is obtained. In this case a crossover between G′ and G″ is observed even at very low frequencies, indicating that elasticity dominates the linear viscoelastic response within this region. In the case of PE2/CNT an increase of 1 order of magnitude is measured, but in the case of PE3/CNT the increase seemed to be only marginal (see the values of η0 listed in Table 1). The enhancements of linear viscoelastic properties observed in PE1/CNT and PE2/CNT samples are quite similar to those obtained in unentangled PDMS/CNT suspensions, for CNT contents as low as 0.1 wt %.70 The transition found at low frequencies by these authors has been attributed to the appearance of a second relaxation zone in the semidilute regime of a rigid rod suspension. In our case, however, the effect of entanglements should play a role, as the Mw of the matrices are well above the entanglement molecular weight. The creep−recovery experiments for the PE/CNT nanocomposites have been performed using the same conditions of applied shear stress and temperature as for the neat materials, F


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Figure 7. Experimental storage, G′ (squares), and loss, G″ (circles), moduli of the nanocomposites studied at T = 140 °C: (a) PE1/CNT; (b) PE2/ CNT, and (c) PE3/CNT. The lines are the converted G′ and G″ values obtained from the creep compliance, Jc.

Figure 9. Shear recoverable compliances, Jr, versus recovery time for pure samples (lines) and nanocomposites (symbols) at T = 140 °C. PE1 (dotted line), PE2 (dashed line) and PE3 (solid line). The different symbols are ascribed to different creep times for the nanocomposites: PE1/CNT: (red □) 1250 s; (red ○) 6250 s; PE2/ CNT: (blue ○) 6250 s; (blue △) 12 500 s; and PE3/CNT: (△) 12 500 s; (▽) 25 000 s.

Figure 8. Comparison of the magnitude of complex viscosity of the neat materials (dashed lines) and of the nanocomposites obtained from the combination of creep (solid lines) and dynamic experiments (symbols) at T = 140 °C. (red ○) PE1/CNT, (blue △) PE2/CNT, and (□) PE3/CNT.

and they also allowed us to obtain the magnitude of Je0. The creep times applied were shifted to higher values than in the case of pure materials to guarantee a maximum orientation of the systems prior to their recovery, but below 50 000 s (total experimental time below 100 000 s) to prevent the samples from degradation. The values of Jr obtained for the nanocomposites compared to those for the neat PE matrices are observed in Figure 9. The value reached for Je0 is very similar in the three PE/CNT nanocomposites studied, and located around Je0 = 3.0−5.0 mPa−1. The values are also listed in Table 1. Again, an increase of 2 orders of magnitude for PE1/ CNT sample and of 1 order of magnitude in the case of PE2/ CNT sample, with respect to those obtained for the neat materials, can be observed in Figure 9. However, the PE3 matrix seems to be only slightly affected by the presence of CNTs, as the corresponding value of Je0 for the PE3/CNT nanocomposite remains very close to the value obtained for the pure PE3 sample. We have to take the values obtained for J0e in PE1/CNT and PE2/CNT nanocomposites with caution. In the case of the PE3/CNT nanocomposite, a constant value of J0e is obtained for tc values higher than 12,500 s, similarly as in the case of pure

PE3. However, for PE1/CNT and PE2/CNT samples, maximum tc values of 6,250 and 12,500 s have been reached, respectively, prior to the recovery experiment, in order to obtain the constant value of J0e . These tc values are located inside the “plateau” region of the Jc curve for these samples (see Figure 6a), preventing the samples to reach higher values of the imposed strain up to times of around 50 000 s or higher. A deeper study is in progress in order to test the suitability of the experimental conditions used to reach the required maximum molecular orientation for the correct estimation of J0e on these specific samples. Notwithstanding, the tc values, and consequently the strain reached in the creep experiments, are high enough to obtain a higher subsequent recovery than in pure PE1 and PE2 samples (Figure 9), which is a clear signature of a characteristic broad relaxation time distribution in PE1/CNT and PE2/CNT samples as referred to in previous lines. III.3. Thermal Properties. Non-Isothermal Crystallization. Figure 10 shows standard DSC cooling scans from the melt and subsequent heating scans for both neat PEs and the nanocomposites with CNTs. Both neat PE2 and PE3 exhibit G


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nucleation temperatures (Ts) and their subsequent heating scans for PE1 and PE1/CNT samples.

Figure 10. DSC cooling (a) and subsequent heating scans (b) for the neat PEs and the corresponding nanocomposites with CNTs.

monomodal crystallization and melting peaks at temperature values that are characteristic of HDPE. In fact, PE2 is a linear PE while PE3 is a bimodal ethylene/1-butene copolymer but with only 4.4 CH3/1000C. PE1 on the other hand is an ethylene/1-hexene copolymer with a much higher SCB content (i.e., 10.7 CH3/1000C) and displays bimodal crystallization exotherms and endotherms (which are a typical sign of a bimodal distribution of SCB) at substantially lower temperatures as compared to PE2 and PE3. The incorporation of well dispersed CNTs cause nucleation effects as indicated by the shift in the crystallization peaks to higher temperatures. In Table 2, the calorimetric data obtained from Figure 10 is presented. The most relevant effect seen in this table is the

Figure 11. Self-nucleation behavior of PE1 (a and b) and PE1/CNT (c and d) nanocomposite. Cooling scans from the indicated Ts temperatures (a and c) and subsequent heating scans (b and d).

Table 2. Calorimetric Data extracted from Figure 10 composition

Tc onset [°C]

Tc peak [°C]

Tm onset [°C]

Tm peak [°C]

Xc [%]

PE1 PE1/CNT PE2 PE2/CNT PE3 PE3/CNT

104.2 113.5 115.2 118.7 118.4 119.8

101.2 108.8 112.4 115.1 115.2 116.2

111.6 110.8 124.4 124.7 121.7 121.2

117.4 119.5 132.8 134.5 133.0 132.7

42 41 67 73 75 60

In the case of neat PE1 (Figures 11, parts a and b) the typical behavior of semicrystalline polymers is observed.60 At temperatures of 122 °C and above, the sample is in domain I where complete melting is achieved and no changes in crystallization and melting are observed since the nucleation density of temperature resistant heterogeneities remain constant. At a Ts = 121 °C the cooling scan shows a typical self-nucleation effect, where the crystallization temperature has been increased, while the melting trace does not exhibit any signs of annealing. This is the typical behavior of domain II or exclusive self-nucleation domain. At Ts = 120 °C, the sample crosses to domain III or the self-nucleation and annealing domain. Notice how the melting trace exhibits a high temperature small peak that is produced by the annealing of a small crystal population that was unmelted at 120 °C and was able to thicken during the 5 min holding step at that temperature. When the self-nucleation behavior of PE1/CNT is examined, a dramatically different behavior was encountered as shown in Figure 11, parts c and d. Figure 11c shows that the cooling scans do not change at all in the Ts range of 170−125 °C. However, Figure 11d clearly shows (also in a close up as an inset) that at Ts = 124 °C a small population of the crystals have experienced annealing. In other words, we do not observe self-nucleation but a direct transition from domain I (i.e., the fully molten domain shown for Ts values 170 and 125 °C) to a domain III where only annealing is appreciated (also denoted domain IIIA).71,72 Domain II is absent (see Table 3), and selfnucleation is absent in domain III. The lack of self-nucleation caused by self-nuclei can be attributed to the excellent

increase in crystallization temperature. The largest increase in the peak crystallization temperature was produced in PE1 (i.e., 7.6 °C), followed by PE2 (i.e., 2.7 °C) and PE3 (i.e., 1 °C). All increases in Tc are significant but the effect on PE1 is remarkable and it will be quantified in terms of nucleation efficiency below. A slight increase in the peak melting point of the nanocomposites with PE1 and PE2 matrices as compared to the corresponding neat polymers was detected while the crystallinity was not significantly affected. In the case of PE1, it is worthy of notice that an increase of the peak melting point of 2.1 °C was produced when 1 wt % CNTs were added, especially when the results obtained by SSA are considered (see below). We have studied the self-nucleation of the materials to ascertain the effect of the carbon nanotubes on the presence and location of the self-nucleation domains and to determine the nucleation efficiency of the CNTs. Figure 11 shows examples of relevant DSC cooling scans from selected selfH


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CNTs can be considered good nucleating agent for PE3 since NE in that case was 71%. It is note worthy that the nucleation efficiency of CNT unexpectedly increases as the molecular weight decreases. Isothermal Crystallization. The study of isothermal crystallization by DSC yields kinetic results that include primary nucleation and growth (i.e., overall crystallization kinetics). The experiments were performed according to the experimental guidelines given by Lorenzo et al.66 and previous tests were performed to ensure that true isothermal data was obtained for the entire crystallization process of each sample (i.e., that the sample did not start to crystallize during the cooling from the melt to Tc). The inverse of the half-crystallization time is a quantitative measure of the overall crystallization rate. This experimentally determined quantity is plotted as a function of the isothermal crystallization temperature in Figure 12. The crystallization

Table 3. Self-Nucleation Domains and Nucleation Efficiency (NE) composition

Ts I−II [°C]

Ts II−III [°C]

Ts I−III [°C]

Tc max [°C]

NE (%)

PE1 PE1/CNT PE2 PE2/CNT PE3 PE3/CNT

122 − 131 − 130 −

120 − 130 − 129 −

− 124 − 130 − 129

106.5 − 115.2 − 116.6 −

100 143 100 96 100 71

nucleation effect of CNTs, which can be even more efficient than polyethylene crystals to induce nucleation in PE1. This result will be quantified by determining the nucleation efficiency of the CNTs (see below). The absence of domain II was first reported in crystallizable blocks of confined block copolymers71,72 and was also previously reported by us in PE/ CNTs with much higher amounts of nanofiller (i.e., 7% or larger).35 It is the first time that such a disappearance of domain II is reported in a PE/CNT nanocomposite with only 1 wt % CNTs addition. For the other samples (PE2, PE2/CNT, PE3, PE3/CNT) the self-nucleation behavior was qualitatively similar in the sense that domain II was also present for PE2 and PE3 but absent for PE2/CNT and PE3/CNT and a direct transition from domain I to domain III was recorded. However, for both PE2/CNT and PE3/CNT, domain III included both selfnucleation and annealing (i.e., the classic behavior of domain III). The results can be explained by the reduced nucleating power of the CNTs in the case of PE2/CNT and PE3/CNT as compared to PE1/CNT. Table 3 summarizes the selfnucleation results obtained. The nucleation efficiency (NE) can be calculated by73 NE =

Tc , CNT − Tc Tc ,max − Tc

Figure 12. Overall crystallization rate, expressed as the experimental inverse of half-crystallization time, as a function of isothermal crystallization temperature for the indicated samples. The solid lines represent the results of fittings to the Lauritzen and Hoffman nucleation and growth theory, used here as a way to guide the extrapolation of the data.

× 100 (1)

where Tc,CNT is the peak Tc value determined in a DSC cooling run for the nanocomposite under consideration, Tc is the peak Tc value for neat PE and Tc,max is the maximum peak crystallization temperature determined after neat PE has been self-nucleated at the ideal self-nucleation temperature (i.e., at the self-nucleation temperature that produces maximum selfnucleation without any annealing). Table 3 shows the nucleating efficiency of the CNTs employed in this work on their corresponding PE matrices. Self-nuclei are ideal nucleating agents for the polymer in consideration, hence the value of 100% NE in Table 3. If a nucleating agent is capable of being even more effective than self-nuclei, the value of NE is higher than 100% and in that case, a supernucleation effect is detected.34,41−43 Müller et al. reported supernucleation for the first time for in situ polymerized PE/CNT nanocomposite masterbatches (with CNT contents of 7% and higher), and later for PEO-g-MWNT, PCL-g-MWNT and PCL/MWNT with smaller amounts of CNTs.34,41−43 Supernucleation effects are believed to be caused by excellent dispersion and/or outstanding interactions between CNTs and the matrix.26,42,43 In Table 3, CNTs are supernucleating agents (i.e., NE = 143%) only for PE1, which is the sample with the largest amount of SCB and lowest molecular weight among the PEs employed in this work (see Table 1). CNTs are also excellent nucleating agents for PE2, since their NE is 96%, almost the same as self-nuclei. Finally,

kinetics is expected to be more sensitive to branch content in this case (since they interrupt the linear crystallizable sequences) than Mn differences, since the three samples are well above their entanglement molecular weight. The order of overall crystallization rate found is proportional to the degree of SCB in the polymers employed: PE2 > PE3 ≫ PE1. In fact the difference between neat PE2 and PE3 is rather small, as is the SCB content of PE3. The nucleation effect of CNTs accelerate the overall crystallization rate and shift the 1/τ50% versus Tc curves to higher temperatures (see Figure 12), as the nucleated matrices will need less supercooling to crystallize. Therefore, the acceleration of the overall crystallization rate is completely dominated by nucleation and is fully consistent with the nonisothermal results presented above. In fact, if a horizontal line is drawn at a constant value of 0.5 min−1 in Figure 12, the difference in the curves between each neat PE and its corresponding nanocomposite is almost the same as the difference in peak crystallization temperatures recorded during nonisothermal crystallization in Figure 12 and Table 2. A quantitative comparison can be observed in Figure 13. It is interesting to note that this figure summarizes both the nonisothermal and the isothermal nucleation effects in relative terms by comparing each nanocomposite with its matrix. These I


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presence of a comonomer or stereodefects along the chains) than to MWD. The results presented in Figure 14 are a good example. The quality of the fractionation is related to the clear separation of thermal fractions in individual melting peaks, and in Figure 14, it can be observed that it follows the order: PE1 ≫ PE3 > PE2. It must be remembered that the SCB content in these polymers are 10.7, 4.4, and 0 (in CH3/1000 C) for PE1, PE3 and PE2 respectively. The fractionation profile reveals the distribution of SCB in the sample. Hence, PE1 exhibits a broad and bimodal distribution of SCB, while PE3 shows a unimodal and shorter distribution of SCB. In the case of PE2, the poor resolution of the fractionation is due to its lack of SCB, therefore, the fractionation in this case is performed by differences in chain lengths. In the case of the nanocomposites, SSA reveals interesting differences as compared to the homopolymer that can be ascribed to the carbon nanotubes effect on the polyethylene chains. By far the largest effect can be seen in Figure 14 for the PE1/CNT nanocomposite. In this case, SSA reveals a new high melting point fraction (signaled with an arrow in Figure 14 and absent in neat PE1) that is a direct consequence of the supernucleation effect of CNTs on PE1. Even in the nonisothermal DSC results presented above (see Table 2) the melting point of PE1/CNT was already significantly higher (for a branched PE) than PE1. This is a consequence of the generation of a population of crystals with enhanced thermodynamic stability (i.e., with thicker lamellae) induced by the supernucleating action (see Table 3) of the CNTs. In a previous work, we found that in HDPE/CNT nanocomposites with much larger amounts of CNTs (>7%) if the samples were isothermally crystallized at a high Tc temperature for a long period of time (days), the nanocomposites were able to develop thicker lamellae as compared to neat HDPE.34 Recently, a similar effect has been found in PCL/CNT nanocomposites with prolonged isothermal crystallization treatments.74 In the case of PE2/CNT nanocomposites, there is also an enhanced annealing capacity of PE2 when it is nucleated by the CNTs (see the arrow in Figure 14), as indicated by the melting point increase in the highest melting point fraction. A similar effect was reported in a previous work for HDPE/CNT masterbatches prepared by in situ polymerization.34 Finally, for PE3, the differences in SSA thermal fractionation are not significant in comparison with PE3/CNT. It should be noted that the annealing capacity of the nanocomposites increases as Mw of the PE matrix decreases. Of course, SCB has a very important effect, since the quality of the fractionation increases with SCB content. Therefore, in PE1 we have a combined effect of both the highest SCB content and lowest Mn values that play a synergistic role in the differences between PE1 and its corresponding nanocomposite after SSA. The effect of CNT on the PE1 matrix is striking since their presence induces the formation of a new fraction of more stable crystals (whose melting point is higher) that can form under the conditions provided by SSA (Figure 14). SSA fractionation drives molecular fractionation in order to produce maximum stability lamellae because of the successive self-nucleationcrystallization-annealing steps. Hence it can be compared with results obtained after prolonged isothermal crystallizations where also CNTs have been found to promote the nucleation and growth of thicker crystals.34,74 In fact, PE1/CNT nanocomposites exhibited higher melting points than neat PE1 even under nonisothermal conditions (Table 2). This is a

Figure 13. Comparison of nucleation effects obtained under nonisothermal and isothermal conditions, expressed as relevant crystallization temperature differences (see text).

nucleation effects are unexpectedly a function of the molecular weight of the matrix regardless of the influence of the SCB content on decelerating the overall crystallization kinetics (Figure 12). Thermal Fractionation by Successive Self-Nucleation and Annealing (SSA). The SSA technique was designed and implemented by Müller et al.60−65 as a new way to improve both the sensitivity and fractionation times of step crystallization techniques. In the case of nanocomposites, the technique can reveal the influence of nanotubes on the molecular fractionation and annealing capacity of polymer chains. Figure 14 shows final heating scans after SSA. The results show complex heating scans composed of multiple melting

Figure 14. Final DSC heating scans after SSA for the neat polymers (black thicker lines) and their corresponding nanocomposites (blue lines). The arrows indicate large differences observed in the highest melting temperature fractions.

peaks. This melting peak distribution is derived from a distribution of lamellar thickness as each thermal fraction is composed by lamellae of different thickness where different crystallizable sequence lengths have been molecularly segregated by a combination of self-nucleation, isothermal crystallization, and annealing.60 It is well-known that SSA, as well as any thermal fractionation technique,60 is particularly sensitive to SCB that interrupt methylene linear sequences that are able to crystallize. In fact, it is much more sensitive to the presence of branches (or any other defect capable of disrupting crystallization like the J


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Figure 15. Scheme showing the possible differences in the interaction between macromolecules and CNTs as a function of molecular size (based on Mw values).

characteristic overlap concentration in ideally dispersed nanocomposites as φc = d7/5lp−3/5 L−4/5. For the given CNT dimensions this equation gives a volume fraction of φc = 0.00025−0.0007. To make comparisons with our experiments (in which one measures the CNT loading by weight percent), one needs to convert the volume fraction φ into the weight fraction, w. Using the density of CNT, ρCNT = 2.1 g/cm3,2 we can estimate the percolation threshold at wc = 0.1−0.5 wt % (this value compares well with that obtained by Huang et al.84 of 0.5−1.5 wt % for CNTs with a higher dimension d = 60−100 nm). This range is certainly much lower than the amount of CNT in our nanocomposites, which means that geometrical percolation should exist in the three samples studied (as also indicated by TEM), and thus the enhancements in linear viscoelastic properties with respect to the neat polymers should have been similar for the three nanocomposites prepared.81 From our rheological results, the similar CNT network in the three PE/CNT systems does not actually give the same enhancement in the linear viscoelastic properties. In fact, it seems that the Mw of the matrix is the main factor controlling the observed results. Indeed, the higher the molecular weight the lower is the effect on η0 and J0e caused by the CNTs. More interestingly, if the matrix already exhibits very high values of η0 and Je0 (or relaxation times) due to higher Mw and/or polydispersity index, like in the case of PE3, the CNT network is screened. As suggested by Pötschke et al., the superposition of the polymer network, rather than the CNT network alone, dominates the linear viscoelastic fingerprint of the nanocomposites.82 As judged by our results, however, this effect seems to be more pronounced as the terminal relaxation time of the matrix increases. The differences obtained for the basic rheological properties in the PE/CNT samples studied here are even more pronounced than those recently observed by Trievel et al. for PS/silica nanocomposites.23 These authors also found a stronger effect of the silica nanofillers in the rheological properties of a PS sample with a narrow MWD as compared to a PS sample with a broader one, but they also stated the elusive nature of their observations. Other authors have found an effect of the Mw of the polymer in the dynamics of the corresponding nanocomposites in the case of the incorporation of spherical silica nanoparticles.82 In the case of PDMS nanocomposites below the polymer critical molecular weight for the onset of entanglements, Mc, the incorporation of nanoparticles increases η0 relative to the polymer matrix. Above Mc, η0 decreases with the volume fraction of nanoparticles. The reduction found in η0 in this case

remarkable fact for PE/CNT nanocomposites, because such large effects on the melting point and crystal stability had been only reported with much higher loadings of CNTs.34

IV. DISCUSSION The chemical structure of the three PE samples used for the preparation of the nanocomposites is virtually the same (mainly −CH2− units). Thus, we do not expect different interactions between the matrices studied and the PE-coated CNTs. It has been reported in a few recent works that the viscosity of the matrix is a key factor in the morphological state of the final dispersion of this type of nanocomposites. In all these studies the nanocomposites are prepared by melt mixing untreated CNTs with polymeric matrices of different nature (PE, PP, EPDM, PS, PMMA, and PC) in corotating twin screw extruders. The results show that a better state of dispersion (usually observed by means of microscopy techniques) and a lower percolation threshold are found for the matrices with lower viscosity.75−82 In general, the explanation given for these results is the restricted mobility of the CNT when using high viscosity matrices for the preparation of nanocomposites. As the CNT-based masterbatch employed here for the preparation of the PE/CNT nanocomposites is the same whatever the PE matrix investigated, the simplest explanation for the differences found in their basic rheological features could be a different state of dispersion of the CNTs. Taking into account those previous works in the literature, one could suspect a different quality of the CNT dispersion in the samples, i.e., a poorer quality as the viscosity of the matrix increases. However, (i) in the master batch used for the preparation of the nanocomposites the bundles of CNTs have been already previously disaggregated by means of the in situ polymerization of ethylene on their surface;10,34,35 and (ii) the micrographs obtained by TEM represent a proof that a good dispersion of the PE-coated CNTs has been achieved and also the formation of a CNT entangled network in the nanocomposites. The formation of a CNT percolated network can be theoretically predicted, if the dimensions of the CNTs are known. On the basis of the semiflexible “wormlike chains” framework,80 it is possible to estimate the characteristic overlap concentration, which theoretically signals the boundary between dilute individual tubes in solution and semidilute overlapping tube coils regimes. Given the dimensions of the CNTs: diameter d = 5−27 nm, persistence length lp = 0.5−1 μm,81 and total arc length L = 5−15 μm, we can estimate the K


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nonisothermal and isothermal nucleation of CNTs increased as the molecular weight of the PE matrix decreased in spite of the role played by SCB (see Figure 13). The annealing capability for producing more stable lamellae in the nanocomposites as compared to their neat matrices also followed a similar trend with Mw as indicated by the SSA results (Figure 14). Nevertheless, the fractionation quality and extend were dominated by SCB content.

was attributed to a dilution of the entanglement density of polymer chains and an increase in the translational motion of the macromolecules. Mackay et al. also found viscosity reductions in polymer nanocomposites of well dispersed nanoparticles, attributed to polymer confinement, which may either cause enhanced constraint release and/or introduce free volume effects.87−89 Martiń et al. have first observed by neutron spin echo measurements the first microscopic evidence of the dilution of the entanglement density of confined entangled polymer in nanopores, but also a slowing down of the local dynamics.90 Because of the different molar mass, the radii of gyration (on average) of the three polymers is different, increasing with Mw, then the occupancy rate of molecules within the CNTs for a given volume fraction (and then for a given internanotube distance), decreases (see the sketch in Figure 15). By taking Mw as a measure of the molecular size, the difference in the radius of gyration between the neat polymers can be calculated as RgPE3 = 1.3RgPE2 = 1.7RgPE1 (the differences are even higher if Mz is taken into account: RgPE3 = 2.4RgPE2 = 3RgPE1). As Mw increases a decreased number of molecules will be able to interact with the CNTs surface due to the restricted space between the CNTs. These larger molecules would see their configurational space altered and consequently their entanglement density, giving rise to less pronounced changes in η0 and J0e than those produced by smaller macromolecules with respect to the bulk materials. These results have both fundamental and practical importance concerning the rheological percolation threshold. The main implication concerns the location of the percolation threshold by rheological methods, i.e., although the electrical (geometrical) percolation is reached, it will not be possible to detect the CNT network if the relaxation time of the matrix is high enough (due to higher Mw and/or polydispersity index values) to screen it and/or to suffer constraining effects. The differences in nucleation and crystallization rate among the different neat PEs employed here are dominated by their SCB content with limited influence of the molecular weight as expected. However, the incorporation of 1% PE-coated MWCNTs had a tremendous impact on the nucleation and crystallization ability of the prepared nanocomposites, especially for the PE1/CNT case. In fact, the order found with respect to the nucleation efficiency was similar to that found in the rheological properties where PE1/CNT experienced the maximum change in nucleation followed by PE2/CNT and finally by PE3/CNT where the effect was minimum, in agreement with a slower local dynamics of the chains with higher Mw. However, it must be emphasized that branches interrupt linear crystallizable sequences and the effect of the Mw must be scaled down to the available crystallizable methylene portions of the chains in between branches. In this sense, the lower Mw of PE1 will also provide many more contacts per unit surface area of CNTs that could be used as nucleation sites by the crystallizable methylene segments. The higher the number of contact points between crystallizable methylenic segments and CNTs the higher the probability of heterogeneous nucleation at the CNT surface. Even though PE1/CNT is the nanocomposite that crystallizes and melts at lower absolute temperatures, as compared to the others (Figure 10) in view of the SCB content of PE1, it is also the only nanocomposite where supernucleation was observed. Our results indicate that unexpectedly, in relative terms (comparing each nanocomposite with its neat polyethylene matrix), the Mw values are also playing an important role. Both

V. CONCLUSIONS We have demonstrated that molecular weight effects can play a determining role in melt viscoelastic properties, while they can also influence to a certain extent nucleation, annealing and melting of PE/CNT nanocomposites. The influence of short chain branching content on the polyethylene matrices is, as expected, very important in determining their crystallization kinetics. Our findings have important implications on the determination of rheological percolation thresholds, since at high molecular weights (or very broad polydispersities), a viscoelastic screening effect of the CNTs may prevent the detection of percolation.

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AUTHOR INFORMATION

Corresponding Authors

*(A.J.M.) E-mail: [email protected]. *(J.M.-S.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The authors acknowledge funding support from the Spanish Ministerio de Economiá y Competitividad (MINECO) for the Project: MAT2012-36341. We acknowledge Repsol (Spain) for providing the noncommercial PE1 and PE2 samples.

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REFERENCES

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