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polyethylene composites
Polymer 52 (2011) 1124e1132
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Polymer journal homepage: www.elsevier.com/locate/polymer
Structureeproperty relationships on uniaxially oriented carbon nanotube/polyethylene composites Giuliana Gorrasi a, *, Roberta Di Lieto a, Giovanni Patimo b, Salvatore De Pasquale b, Andrea Sorrentino a, * a b
Department of Chemical and Food Engineering, University of Salerno, via Ponte don Melillo, 84084 Fisciano, Salerno, Italy Department of Physics E.R. Caianiello, University of Salerno, via Ponte don Melillo, 84084 Fisciano, Salerno, Italy
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 November 2010 Received in revised form 1 January 2011 Accepted 5 January 2011 Available online 11 January 2011
Multi walled carbon nanotubes have been incorporated into a linear low density polyethylene matrix through high energy ball milling technique at room temperature, without any chemical modification or physical treatment of the nanotubes. Highly oriented samples, with different draw ratios, were obtained by drawing at 80 C the composite films. SEM and FTIR results on the drawn PE films demonstrate that the molecular chains in both crystalline and amorphous phases are well oriented along the drawing direction. The effect of different weight percent loadings of nanotubes and draw ratio on the morphology, thermal, mechanical and electrical properties of the composite fibers have been investigated. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Carbon nanotube/polyethylene composites High energy ball milling Highly oriented polymeric films
1. Introduction The technological significance of the highly oriented polymer materials derives from their properties that may exceed those of the isotropic species by orders of magnitude. For example, the stiffness and strength of highly oriented crystalline polymers [1,2] can increase to a factor of 100 compared with their non-oriented counterparts. For flexible regular chain polymers, macromolecules tend to fold upon crystallization and different routes have been developed to transform chain-folded crystals into extended structures [3,4]. Carbon nanotubes (CNT) have attracted interest as reinforcing fillers because of their excellent mechanical and thermal properties [5,6], but they are also regarded as the ultimate fillers for several advanced applications [7,8]. Carbon nanotubes possess an extremely high elastic modulus z1 TPa, comparable to that of diamond (1.2 TPa) and report strengths of 10e100 times that of the strongest steel [7,8]. In addition, they exhibit electrical conductivity as high as 105e107 S/m [9] and can transform an insulating polymer into a conducting composite at very low loading because of their extremely high aspect ratio [10]. The introduction of nanofillers like CNTs in polymer fibers can lead to multifunctional high-performance materials that combine high strength with electrical conductivity
* Corresponding authors. Fax: þ39 089 964057. E-mail addresses: [email protected] (G. Gorrasi), [email protected] (A. Sorrentino). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.01.008
[11]. The keys to achieve maximum performances from nanocomposite are to obtain a uniform distribution of nanoparticles within the polymer matrix, and to have the best nanoparticlee polymer adhesion, which is critical for load transfer from the matrix to the particle. Nanotubes can be dispersed in the polymer matrices using different techniques such as melt mixing, solution processing, or in situ polymerization. An alternative and innovative strategy relies on solid-state mixing at near room temperature, which ought to involve an efficient mixing of two or more species by mechanical milling, avoiding high temperatures and solvents. High Energy Ball Milling (HEBM) is an effective unconventional technique currently used in material synthesis and processing [12]. It consists of repeated events of energy transfer, promoted by the milling device, from the milling tools (generally balls) to the milled powder. Recently it has been proved that HEBM on polymeric materials can help obtaining materials with new characteristics, which can be barely achieved through other conventional processes [13e15]. Uniaxial alignment of CNT using magnetic [16], shear fields [17], casting-drawing and gel spun processes [18,19] has been the focus of several recent studies. Linear Low Density Polyethylene (LLDPE) is a commodity polymer, cheap, versatile, with well-known applications and commercial uses in a variety of forms, including fibers. Thus, enhancing of LLDPE properties through the dispersal and alignment of CNT should be of significant interest in order to expand its application fields. Very few papers report the preparation and characterization of Polyethylene/Carbon Nanotubes fibers. In particular, only few papers present a detailed
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
investigation about the arrangement of both the PE chains and the carbon nanotube in melt drawn oriented composite films [9e11,17]. In the present work we performed Multi walled Carbon Nanotubes dispersion an LLDPE matrix through HEBM at room temperature in the dry state with no chemical or physical treatment of the nanotubes. Fibers with different draw ratios were obtained by drawing at 80 C the composite films. The effect of different weight percentage of nanotubes and draw ratio was analyzed as function of the morphology, thermal, mechanical and electrical properties. 2. Experimental 2.1. Materials Multi walled carbon nanotubes (CNT) were purchased by Nanocyl (Belgium) (NC 3100). They were synthesized by catalytic carbon vapor deposition (ccvd) process. By thermogravimetric analysis (TGA) the carbon purity was found higher than 95% whereas the metal oxide impurity was less than 5%. The specific surface area determined with the BET method was around 250e300 m2/g. Linear Low Density Polyethylene (FlexireneÒ CM50) was supplied from Polimeri Europa (Italy) in ultrafine powder, form with a melt flow index MFI ¼ 4.1 g/10 min (at 190 C, 2.16 kg).
1125
stirring 1.0 g of potassium permanganate in a solution of 95 mL of sulfuric acid (9597 vol%) and 48 mL of orthophosphoric acid (85 vol%). Subsequent to the etching treatment, a first washing was done using a cold mixture of 2 parts by volume of concentrated sulfuric acid and 7 parts of distilled water. A second washing was performed with 30 vol % aqueous hydrogen peroxide to remove any manganese dioxide and, finally, a 3 times washing using distilled water. Before the SEM analysis the sample was kept under vacuum for 2 days at ambient temperature. The samples were finally covered with a 250 Å thick gold film using a sputter coater (Agar mod. 108 A) and observed with SEM. IR spectra were collected by means of an FTIR Spectrometer M2000 FTIR (by Midac Co.). For each sample, the average of 32 scans was used, working at a resolution of 2 cm1. The scan wave-number range of the collected spectrum was 4000e400 cm1. In order to evaluate the molecular orientation of the samples, the IR beam was polarized with the polarization axis parallel and orthogonal to draw direction [20]. Polarization of the beam was done by a zinc selenide wire grid polarizer. The average Hermans orientation factor fav can be determined by measuring the dichroic ratio D ¼ (Ap/Ao) where Ap is the absorbance for the plane parallel to the draw direction and Ao the absorbance for the polarization plane orthogonal to the draw direction. The Hermans factor, assuming uniaxial orientation, is related to dichroic ratio at wavenumber n by [21]:
D1 Dþ2 n
2.2. Composite preparation (HEBM process)
fi ¼ Kn
CNT powders and Polyethylene, as received, were milled in the solid state in a Retsch (Germany) centrifugal ball mill (model PM100). The milling process was carried out in a cylindrical steel jar of 50 cm3 with 5 steel balls of 10 mm of diameter. The rotation speed used was 650 r.p.m. and the milling time was fixed to 60 min. HEBM of powders constituted by organic polymers and fillers has been proved to be an alternative and efficient technique to produce novel composites with high performances [13e15]. During the milling the carbon nanotube bundles crack, and “intimate mixing” are promoted [15]. In these experimental conditions, five series of composites LLDPE/CNT with 1, 2, 3, 5, 10 wt/wt% of carbon nanotubes were prepared. An additional LLDPE sample to be taken as a reference was also milled in absence of filler.
where Kn is an auxiliary variable that in our case was assumed to be 1. The absorption bands at 730 and 719 cm1 were employed to evaluate the orientations of the crystal a-axis and b-axis, respectively [20]. The IR technique gives fa and fb, from which fc can be evaluated by means of equation (2):
2.3. Samples preparation The LLDPE/CNT mixtures and the pure milled LLDPE were molded in a hot press (Carver Inc.) at 170 C forming 250 50 mm thick films, which were rapidly quenched in a watereice bath (0 C). From the obtained films were obtained strips large 1 cm. The strips were put between the clamps of an Instron Dynamometer (Mod 4301), equipped with a temperature chamber. The temperature was fixed to 80 C. As soon as the temperature was stabilized, the upper traverse of the dynamometric apparatus was moved at a speed rate of drawing of 10 mm/min. Films were drawn at different draw ratios l ¼ l/lo, where l is the final length and lo the initial length. The selected values of the draw ratios (followed by a display) for all the composites were l ¼ 6; 8; 10. In the following, the resulting samples will be coded as follows: PEXlY where X ¼ 1, 2, 3, 5, 10 is the weight percent of CNT and Y is the draw ratio of the samples. 2.4. Methods of investigation Scanning electron microscopy (SEM) analysis was performed with a LEO 1525 microscope. The etching mixture was prepared by
fa þ fb þ fc ¼ 0
(1)
(2)
Differential scanning calorimetry (DSC) analysis was carried out on samples with a mass ranging between 5 and 7 mg. The tests were carried out by means of a DTA Mettler Toledo (DSC 30) under nitrogen atmosphere. The samples were heated from 25 C to 200 C at 10 C/min. To ensure reliability of the data obtained, heat flow and temperature were calibrated with standard materials, indium, and zinc. The percent crystallinity was calculated by taking the specific heat of fusion of perfectly crystalline PE to be 290 J/g [22]. Thermogravimetric analysis (TGA) was carried out with a Mettler TC-10 thermobalance. Polymer composites samples were heated from 25 to 800 C at 10 C/min heating rate under air flow. The weight loss was recorded as function of temperature. The elastic moduli were evaluated using an Instron Dynamometer (Mod 4301), and derived from the initial part of the stressestrain curves, giving to the sample a deformation of 0.1%. The experiments were conducted at room temperature with the deformation rate of 2 mm/min. The initial length of the samples was 10 mm. The data were averaged on five samples. The electrical conductivity was measured with a Keithley 6517A source measurement unit in a two-probe resistance measurement configuration. The sample thicknesses were carefully measured by a micrometer, whereas the length and the width were 3 and 50 mm, respectively. The electrical conductivity was measured in the voltage range 10 þ 10 V. All the samples showed a linear behavior of the current (I) versus the applied voltage (V). The electrical conductivity was calculated by using the equation (3):
s¼
L 1
sW R
¼
L Imeasured
sW Vapplied
(3)
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where R, s, W, L are the resistance, the thickness, the width and the length of the specimens respectively. 3. Results and discussion 3.1. Morphological analysis Fig. 1 reports the SEM images of some etched samples. The undeformed sample (Fig. 1a) shows the classical spherulitic morphology. A uniform morphology with structure aligned randomly along the sample surface can be observed. Under high uniaxial drawing in the solid state, the original morphology is transformed into highly oriented microfibril morphology (Fig. 1b). The arrow in the picture represents the drawing direction of the film during stretching. The deformation processes produce the shear of the stressed spherulitic lamellae into crystal blocks via chain slip. These blocks rotate and decrease in width until microfibrils of alternating crystal and amorphous regions are formed. Ward [21,23] proposed a morphological model, which considers that the drawn morphology consists of stacks of short lamellar-type crystallites linked by intercrystalline bridges. Since the crystalline phase is essentially continuous in the draw direction, the modulus of the drawn materials is expected to be quite high, whereas the degree of crystallinity is expected to increase with draw ratio. The presence of CNTs in the polymer matrix seems to enhance this deformation process (Fig. 1c). In fact, the oriented composite samples show a clear fibrillar morphology with numerous microcrazes. The CNTs are confined between the fibril and clearly show a high degree of alignment in the drawing direction. Such morphology has a strong influence on the thermal, mechanical and electrical properties of the resulting samples. 3.2. Detection of molecular orientation The orientation of the structural units in a solid polymer, as a result of various forming processes such as drawing and extrusion, may have a profound influence on the macroscopic physical properties of the material. Several experimental methods have been employed to characterize the orientation of crystalline and semicrystalline polymers. Birefringence, Nuclear Magnetic Resonance, X-ray diffraction, Polarized Raman spectroscopy and Infrared dichroism are among the most popular methods used to characterize different aspects of molecular order [21]. Each of these techniques has inherent advantages and disadvantages when compared with one another. Infrared spectroscopy (FTIR) and birefringence measurements often cannot be used for thick polymer specimens. In contrast, Raman spectroscopy provides an effective means for the study of samples that do not efficiently transmit radiation, provided care is taken to minimize polarization scrambling by the sample [24]. A second advantage of Raman spectroscopy is the ability to determine both the second and fourth coefficients of the orientation distribution function, whereas infrared dichroism and birefringence measurements are sensitive only to the second moment of the expansion [25,26]. Unfortunately, Raman spectroscopy is more complex from both an experimental and theoretical point of view when compared to infrared dichroism and birefringence. Not surprisingly, in spite of the usefulness of the method, its practical applications have been limited to a few specific materials, for which the molecular structures are relatively simple and which polarizability tensors were already known [27]. Due to the presence of not well characterized interactions of the carbon nanotubes with the scattering nature of Raman spectroscopy, in this paper, only the
Fig. 1. SEM microphotographs showing the microstructure of the samples: (a) PE0l0; (b) PE0l8; (c) PE3l8. The arrows show the draw direction.
FTIR technique was employed to determine the molecular orientation of samples. Detailed analyses are available in literature on PE regarding the band assignments, using FTIR spectroscopy combined with
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
polarized infrared beam [20,28,29]. In particular, the band doublet observed in the spectrum of polyethylene in the range 710e 740 cm1 has been analyzed in order to characterize the orientation of the crystal unit cell as a function of the deformation [29]. This band doublet has been assigned to in- and out-of-phase CH2rocking vibrations, respectively, of the crystal phase with the 730 cm1 band polarized along the crystallographic a-axis and the 719 cm1 band polarized along the b-axis [30]. Fig. 2 shows the Hermans factor obtained with equation (1)e(2) as function of the CNT content for the LLDPE composites analyzed in this study. The undeformed samples were also reported as reference in the same Fig. 2. By changing the CNT content, the spectra corresponding to the undeformed samples exhibit a similar dicroic ratio because they have random molecular orientation. As strain was applied to the sample and the strain increased up to l ¼ 6 the changes in the orientation function indicate that there is an appreciable modification in the mean orientation of the crystalline phase. In Fig. 2 the relative molecular orientation is evidenced to increase with increasing CNT content for all the composites. This result indicates that the orientation of the crystallites changes with the extension ratio and the CNT content. In the range 5e10% of CNT all the curves of fc level off, reaching values close to unity. This behavior reflects a typical transformation from spherulitic to fibrous structure during elongation. The transitions in the orientation function curves may arise from yielding and perhaps other morphological changes occurring during straining [31]. It is well known that during the stretching of LLDPE between the glass transition and melting temperature (which is the case here) the amorphous phase will orient itself as the crystalline regions rotate and orient themselves in the direction of stretching, as a result of this process the macromolecules can be oriented by a rotation of lamellae or by unfolding [21].
-0.20 -0.25
=0 =6 =8 = 10
-0.30
fa-0.35 -0.40 -0.45 -0.50
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] 0.3 0.2 0.1
=0 =6 =8 = 10
0.0
fb -0.1 -0.2 -0.3 -0.4 -0.5
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] 1.0
0.8
3.3. DSC analysis The DSC curves obtained for some selected samples are shown in Fig. 3. The unoriented samples display a single broad peak. It is clearly the results of two endothermic events of different intensity. Multiple melting features are characteristic of commercial LLDPEs, due to the presence of a broad distribution of crystal sizes [32] that extends over a broad range of temperatures. Upon stretching, the sample shows (Fig. 3a) a narrower melting peak and at higher temperature compared to that of the isotropic undeformed samples. This phenomenon is associated with the molecular orientation of the structure into the straining direction. For the sample with the higher draw ratio (l ¼ 10), the melting feature tends to move toward lower temperature and the maximum of melting shifts to lower temperature. This behavior can be attributed to a decrease in crystals size, or to an improvement in their perfection drawing inducted, or to the existence of a polymorphic transition occurring during heating [33]. This structural effect is not visible for lower draw ratio, probably because it is overridden by the molecular orientation effect and the disruption of the crystalline phase. Similar behavior in the melting process was observed also in presence of CNTs (Fig. 3b): an increase in draw ratio produces an increase in the melting peak temperature, followed by a decrease at higher draw ratio. However, this second stage seems to be anticipated with respect to the pure polymer. At constant draw ratio, also the CNT content shows a significant effect on the melting behavior of LLDPE. Fig. 3c shows that for the lower CNT content there is an increase in melting temperature with increasing draw ratio, but for higher CNT content the melting temperatures decrease slightly.
1127
0.6
fc
= 10 =8 =6 =0
0.4
0.2
0.0
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] Fig. 2. Orientation function as function of CNT content and draw ratio, obtained with IR analysis of all investigated samples.
It is worth noting that solid-state drawing at 80 C for LLDPE is a condition for additional crystallization. The heats of fusion of unoriented and oriented samples are reported in Table 1. They are higher for the oriented samples and increase with increasing draw ratio. At relative small elongations, the crystallinity was increased due to strain induced crystallization. As the draw ratio increases, the crystallinity decreases as consequence of the fragmentation of crystalline lamellae, increasing the draw ratio the crystallinity is expected to increase again [34]. This behavior could be associated to partial melting of smallest or less perfect crystallites followed by recrystallization. Upon addition of CNT, the heats of fusion decrease slightly probably due to the hindering phenomena from carbon nanotubes.
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PE λ = 10 PE = 8 PE = 6 PE
a
100
110
120
130
140
150
160
170
180
Temperature [°C]
PE3 PE3 PE3 PE3
b
100
110
120
130
140
150
160
= 10 =8 =6 =0
170
180
Temperature [°C] PE5 PE3 PE1 PE
c
=6 =6 =6 =6
Table 1 Melting peak temperature and heat of fusion of all the analyzed samples. Sample
Tm [ C]
Dh [J/g]
Xc
PE PEl6 PEl8 PEl10 PE1 PE1l6 PE1l8 PE1l10 PE3 PE3l6 PE3l8 PE3l10 PE5 PE5l6 PE5l8 PE5l10 PE10 PE10l6 PE10l8 PE10l10
136.2 141.2 140.3 137.6 136.0 140.7 140.3 135.8 131.5 140.6 140.7 139.8 130.0 137.9 139.2 136.7 132.1 137.6 138.4 133.4
168 179 188 190 162 177 180 183 141 174 178 179 179 157 179 179 179 147 147 147
58% 62% 65% 66% 56% 61% 62% 63% 48% 60% 61% 62% 49% 54% 56% 57% 45% 51% 53% 55%
(0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2)
(3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0)
respectively, are to be ascribed to a two steps process of which one controlled by oxygen diffusion. In terms of solid-state oxidation of PE, the oxygen-controlled step may be related to the oxidation of the bulk, while the first stage to the oxidation of the polymer surface and the less to the ordered phase degradation. The last residue (w8%) oxidizes in a last step extended up to 550 C. The thermal stability of the composite fibers results greatly improved compared to the undeformed PE. Table 2 reports the maximum degradation temperature, Tdmax ( C), corresponding to about 50% of weight loss. Either the incorporation of carbon nanotubes into the polymer matrix or the drawing process greatly improves the thermal stability of all the samples, such improvement increases with nanotube increasing and draw ratio. The improvement in thermal stability can be attributed to several synergistic factors. During the thermal degradation, the carbon nanotubes form a shielding layer. This layer reduces the heat transmitted to the core of the sample, prevent the recombination the polymer peroxyl radicals, and form a barrier to the oxygen diffusion [36,37]. 3.5. Mechanical properties
100
110
120
130
140
150
160
170
180
Temperature [°C] Fig. 3. DSC melting endotherms of the reported samples.
3.4. Thermogravimetric analysis Fig. 4 reports the thermogravimetric curves, evaluated in air flow, for all the samples in the range 200 Ce600 C. Pure polyethylene shows an initial degradation temperature at about 320 C. Above this temperature free radicals are generated leading to sequential degradation and break-down of the main chain due to the thermal decomposition of the covalent CeC bond [35]. The two main stages of PE oxidation, centered at 405 C and 454 C
The enhancement of the mechanical properties of composites requires a high degree of load transfer between the matrix and the nanotubes. If the interfacial adhesion between the phases is weak, the nanotubes behave as holes or nanostructured flaws, introducing local stress concentrations, and the benefits of the CNT properties are lost [38]. The nanotubes must be well dispersed. In case of poor dispersion, they will fail by separation of the bundle rather than by failure of the nanotube itself, resulting in significantly reduced strength [8,9]. On the other hand, solid-state drawing of nanocomposites is largely dominated by regions of lower nanotube volume fraction and this could certainly alter both distribution and orientation of CNTs along the length of the fiber. Strength in composites is a complex issue involving load transfer, stress concentrations and defect distribution, especially in the case of fibers. Fig. 5 shows the elastic moduli, as function of CNT percentages, for all the samples at different draw ratios. The drawing process already increased the elastic modulus of pure PE with increasing the draw ratio. The elastic modulus evaluated for the undeformed PE was about 450 MPa, it results about three times increased after drawing at l ¼ 6, about five times increased at l ¼ 8, and about seven times increased after drawing at l ¼ 10. The
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
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5000
Elastic Modulus [MPa]
4500 4000 3500 3000
λ = 10 λ=8 λ=6 λ=0
2500 2000 1500 1000 500 0
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] Fig. 5. Elastic moduli, E (MPa), as function of CNT content.
significant property enhancement, at low CNT percentages (wt/wt), could be attributed to highly dispersed and well aligned nanotubes and/or good adhesion between nanotubes and polymer matrix. Fig. 6 shows the elastic moduli as function of the samples draw ratios. The higher the draw ratio the lower is the improvement in the strength due to the CNT addition. It seems to suggest that the high orientation of the polymer chains enforce high alignment of the nanotubes also at low concentrations. 3.6. Electrical properties Fig. 7 reports the conductivity values, s (S/m), of the fibers, as function of CNT percentages. Percolation theory deals with the effect of varying, in a random system, the number of interconnections present. In this case the interconnections are the highly conductive nanotubes. In literature [39] was proposed an analytical model, based on the Fermi-Dirac distribution, which describes the critical insulator to conductor transition:
Fig. 4. Thermogravimetric analysis (TGA) for all the samples in the range 200 Ce600 C.
composite fibers show significantly increased strength with increasing the draw ratio, in all the investigated composition range. The enhancement is more pronounced at low nanotubes loading (i.e. 1e3% wt/wt), a further increasing in modulus is observed for 5% wt/wt CNT (more evident for samples drawn at l ¼ 6 and l ¼ 8) and a plateau value for 10% wt/wt of filler. The data suggest that the
sc log sn
s log spn ¼ ð1 þ expðtðX Xc ÞÞÞ
(4)
where sc, sn and sp are the composite, filler, and polymer conductivities, respectively, X is the CNT mass fraction, and t is an empirical parameter that leads to the change in conductivity at the percolation threshold Xc. By assuming a constant value for sn and sp, from equation (4) were obtained the best fitted values of Xc and t for the four sets of samples. The data are reported in Table 3. The percolation threshold decreases from the undeformed samples to the fibers, going from 2 to 4 with increasing the draw
Table 2 Maximum degradation temperature, Tdmax ( C), of undeformed and drawn polyethylene and all the composite fibers, evaluated by TGA. Sample Undeformed PE PE0l6 PE1l6 PE2l6 PE3l6 PE5l6 PE10l6
Tdmax ( C) 405 424 450 450 453 463 467
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
Sample
C) C) C) C) C) C) C)
Undeformed PE PE0l8 PE1l8 PE2l8 PE3l8 PE5l8 PE10l8
Tdmax ( C) 405 422 445 454 467 468 468
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
C) C) C) C) C) C) C)
Sample
Tdmax ( C)
Undeformed PE PE0l10 PE1l10 PE2l10 PE3l10 PE5l10 PE10l10
405 417 446 457 456 457 473
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
C) C) C) C) C) C) C)
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G. Gorrasi et al. / Polymer 52 (2011) 1124e1132 Table 3 Percolation threshold, Xc (wt%), electrical conductivity of filler, sn (S/m), electrical conductivity of matrix, sp (S/m), and t parameters as evaluated from Equation (1), as function of draw ratio (l).
Elastic Modulus [MPa]
10000
1000
X = 10 X= 5 X= 3 X= 2 X= 1 X= 0
100 0
2
4
6
8
10
Draw ratio Fig. 6. Elastic moduli, E (MPa), as function of samples draw ratio.
ratio from 0 to 10. Such a phenomenon, already found for nanocomposites based on oriented thermoplastic polymers-carbon nanofillers [40e42], could be justified as a break-down of the conductive network, inducted by the solid-state drawing. The drawing process of LLDPE/CNT composites decreases the number of conduction pathways present above the percolation threshold with a resulting decrease in conductivity. However, we point out that the electrical conductivity of LLDPE, soon after percolation, results increased of about nine orders of magnitude at low nanotube loadings at all draw ratios. Fig. 8 illustrates the samples conductivity as function of the draw ratio. At high draw ratio the platelets are aligned parallel each other along the stretching direction and only at higher loading they form a conductive path. At higher loadings, where there are more tubeetube connections, greater anisotropy is required to break the contacts and disrupt percolation.
4. Discussion
Xc (wt%)
sn (S/m)
sp (S/m)
t
2.0 3.0 3.5 4.0
1.00E-01 1.00E-01 1.00E-01 1.00E-01
2.24E-11 2.24E-11 2.24E-11 2.24E-11
2.0 1.6 1.6 1.6
sliding or shear of the folded molecules. The intermolecular interactions (i.e. van der Waals forces) existing between polymer chains contribute to the sample morphology [43,44]. Based on the SEM experiments, it was proposed that the microstructure of LLDPE is an interconnected network of anisotropically shaped particles, connected by array of fibrils. Considering the FTIR data presented in Fig. 2 and the proposed mechanisms of deformations, the transition at 6
0.1
0.1
1E-3
1E-3
Conductivity [S/m]
Conductivity [S/m]
The initial stages of deformation involved a displacement of the spherulitic lamellae within the sample with a non-homogeneous
l 0 6 8 10
1E-5
=0 =6 =8 = 10
1E-7 1E-9
1E-11
0
1
2
3
4
5
6
7
8
9
CNT content [wt/wt %] Fig. 7. Electrical conductivity (S/m) as function of CNT content (wt/wt%)
10
X = 10% X = 5% X = 3% X = 2% X = 1% X = 0%
1E-5 1E-7 1E-9 1E-11
0
2
4
6
8
Draw Ratio Fig. 8. Electrical Conductivity (S/m) as function of draw ratio.
10
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
1131
Scheme 1. Model illustrating the structural changes during tensile deformation.
direction of loading. In that case, the effect of the CNT seems to be less important with respect to this improvement produced by the lamellae orientation.
5. Conclusions Multi walled carbon nanotubes dispersion was achieved into a linear low density polyethylene matrix through high energy ball milling technique using 1%; 2%; 3%; 5% and 10% (wt/wt) of CNT loadings. Films from composite powders were obtained on a laboratory scale and drawn at 80 C at different draw ratios l ¼ 6; 8; 10. Morphological analysis showed that under high uniaxial drawing in the solid state, the original morphology is transformed into highly oriented microfibril morphology. FTIR demonstrated that molecular orientation increases with increasing draw ratio and CNT content for all the composites. This behavior reflects a typical transformation from spherulitic to fibrous structure during elongation. DSC analysis showed that for the lower CNT content there is an increase in melting temperature with the increasing draw ratio, but for higher CNT content the melting temperatures decrease slightly. The analysis of thermal degradation in air flow showed a high improvement of thermal stability for the composite fibers, proportionally to CNT content and draw ratio. Mechanical properties showed a significantly increased strength with increasing the draw ratio, in all the investigated composition range. The enhancement resulted more pronounced at low nanotubes loading (i.e. 1e3% wt/wt) and a plateau value was reached for 10% wt/wt of filler. Electrical properties showed that the electrical conductivity of insulating LLDPE matrix, results increased of about nine orders of magnitude, with quite low nanotube loadings at all draw ratios, and the percolation threshold tends to decrease with increasing the draw ratio for a break-down of the conductive network, inducted by the solid state drawing.
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[5] Yu MF, Files BS, Arepalli S, Ruoff RS. Tensile loading of ropes of single wall carbon nanotubes and their mechanical properties. Phys Rev Lett 2000;84: 5552e5. [6] Thostenson ET, Chou TW. On the elastic properties of carbon nanotube-based composites: modeling and characterization. J Phys D Appl Phys 2003;36:573e82. [7] Robertson J. Realistic applications of CNTs. Mater Today 2004;7:46e52. [8] Guadagno L, Vertuccio L, Sorrentino A, Raimondo M, Naddeo C, Vittoria V, et al. Mechanical and barrier properties of epoxy resin filled with multi-walled carbon nanotubes. Carbon 2009;47:2419e30. [9] Thostenson ET, Ren Z, Chou TW. Advances in the science and technology of carbon nanotubes and their composites: a review. Compos Sci Technol 2001;61:1899e912. [10] Sandler JKW, Kirk JE, Kinloch IA, Shaffer MSP, Windle AH. Ultra-low electrical percolation threshold in carbon-nanotube-epoxy composites. Polymer 2003;44:5893e9. [11] Bauhofer W, Kovacs JZ. A review and analysis of electrical percolation in carbon nanotube polymer composites. Compos Sci Technol 2009;69:1486e98. [12] Suryanarayana C. Mechanical alloying and milling. Prog Mater Sci 2001;46:1e184. [13] Shaw WJD. Current understanding of mechanically alloyed polymers. Mater Sci Forum 1998;269:19e29. [14] Sorrentino A, Gorrasi G, Tortora M, Vittoria V, Costantino U, Marmottini F, et al. Incorporation of MgeAl hydrotalcite into a biodegradable Poly (3-caprolactone) by high energy ball milling. Polymer 2005;46:1601e8. [15] Gorrasi G, Sarno M, Di Bartolomeo A, Sannino D, Ciambelli P, Vittoria V. Incorporation of carbon nanotubes into polyethylene by high energy ball milling: morphology and physical properties. J Polym Sci Part B Polym Phys 2007;45:597e606. [16] Kuriger RJ, Alam MK, Anderson DP, Jacobsen RL. Processing and characterization of aligned vapor grown carbon fiber reinforced polypropylene. Compos Part A: Appl Sci Manufacturing 2002;33:53e62. [17] Haggenmueller R, Gommans H, Rinzler AG, Fischer JE, Winey KI. Aligned single-wall carbon nanotubes in composites by melt processing methods. Chem Phys Lett 2000;330:219e25. [18] Bin Y, Kitanaka M, Zhu D, Matsuo M. Development of highly oriented polyethylene filled with aligned carbon nanotubes by gelation/crystallization from solutions. Macromolecules 2003;36:6213e9. [19] Ruan S, Gao P, Yu TX. Ultra strong gel spun ultra-high molecular weight polyethylene fiber reinforced using multiwalled carbon nanotubes. Polymer 2006;47:1604e11. [20] Read BE, Stein RS. Macromolecules 1968;1(2):116. [21] Ward IM, Coates PD, Dumoulin MM, editors. Characterization of orientation in solid phase processing of polymers. Munich: Carl Hanser Verlag; 2000. [22] Lagaron JM, Lopez-Quintana S, Rodriguez-Cabello JC, Merino JC, Pastor JM. Comparative study of the crystalline morphology present in isotropic and uniaxially stretched “conventional” and metallocene polyethylenes. Polymer 2000;41:2999e3010. [23] Clements J, Jakeways R, Ward IM. Polymer 1978;639:19. [24] Voyiatzis GA, Andrikopoulos KS. Fast monitoring of the molecular orientation in drawn polymers using Micro-Raman spectroscopy. Appl Spectrosc 2002; 56:528e35. [25] Everall N, Chalmers J, Mills P. Use of polarized resonance Raman spectroscopy of a polyene probe, and FT-IR dichroism, to probe amorphous-phase orientation in uniaxially drawn poly(ethylene). Appl Spectrosc 1996;50:1229e34. [26] Balabin RM. Conformational equilibrium in glycine: experimental jet-cooled Raman spectrum. J Phys Chem Lett 2010;1:20e3. doi:10.1021/jz900068n. [27] Tanaka M, Young RJ. Polarised Raman spectroscopy for the study of molecular orientation distributions in polymers. J Mater Sci 2006;41:963e91. [28] Nielsen JR, Woollett AH. Vibrational spectra of polyethylenes and related substances. J Chem Phys 1957;26(6):1391e400. [29] Kissin YV. J Polym Sci Polym Phys Ed 1992;30:1165. [30] Agosti E, Zerbi G, Ward IM. Structure of the skin and core of ultra-drawn polyethylene films by vibrational spectroscopy. Polymer 1992;33(20):4219e29. [31] Bhattacharyya AR, Sreekumar TV, Liu T, Kumar S, Ericson LM, Hauge RH, et al. Crystallization and orientation studies in polypropylene/single wall carbon nanotube composite. Polymer 2003;44:2373e7. [32] Wunderlich B. Macromolecular Physics, vol. 3. Academic Press; 1980.
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Polymer journal homepage: www.elsevier.com/locate/polymer
Structureeproperty relationships on uniaxially oriented carbon nanotube/polyethylene composites Giuliana Gorrasi a, *, Roberta Di Lieto a, Giovanni Patimo b, Salvatore De Pasquale b, Andrea Sorrentino a, * a b
Department of Chemical and Food Engineering, University of Salerno, via Ponte don Melillo, 84084 Fisciano, Salerno, Italy Department of Physics E.R. Caianiello, University of Salerno, via Ponte don Melillo, 84084 Fisciano, Salerno, Italy
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 November 2010 Received in revised form 1 January 2011 Accepted 5 January 2011 Available online 11 January 2011
Multi walled carbon nanotubes have been incorporated into a linear low density polyethylene matrix through high energy ball milling technique at room temperature, without any chemical modification or physical treatment of the nanotubes. Highly oriented samples, with different draw ratios, were obtained by drawing at 80 C the composite films. SEM and FTIR results on the drawn PE films demonstrate that the molecular chains in both crystalline and amorphous phases are well oriented along the drawing direction. The effect of different weight percent loadings of nanotubes and draw ratio on the morphology, thermal, mechanical and electrical properties of the composite fibers have been investigated. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Carbon nanotube/polyethylene composites High energy ball milling Highly oriented polymeric films
1. Introduction The technological significance of the highly oriented polymer materials derives from their properties that may exceed those of the isotropic species by orders of magnitude. For example, the stiffness and strength of highly oriented crystalline polymers [1,2] can increase to a factor of 100 compared with their non-oriented counterparts. For flexible regular chain polymers, macromolecules tend to fold upon crystallization and different routes have been developed to transform chain-folded crystals into extended structures [3,4]. Carbon nanotubes (CNT) have attracted interest as reinforcing fillers because of their excellent mechanical and thermal properties [5,6], but they are also regarded as the ultimate fillers for several advanced applications [7,8]. Carbon nanotubes possess an extremely high elastic modulus z1 TPa, comparable to that of diamond (1.2 TPa) and report strengths of 10e100 times that of the strongest steel [7,8]. In addition, they exhibit electrical conductivity as high as 105e107 S/m [9] and can transform an insulating polymer into a conducting composite at very low loading because of their extremely high aspect ratio [10]. The introduction of nanofillers like CNTs in polymer fibers can lead to multifunctional high-performance materials that combine high strength with electrical conductivity
* Corresponding authors. Fax: þ39 089 964057. E-mail addresses: [email protected] (G. Gorrasi), [email protected] (A. Sorrentino). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.01.008
[11]. The keys to achieve maximum performances from nanocomposite are to obtain a uniform distribution of nanoparticles within the polymer matrix, and to have the best nanoparticlee polymer adhesion, which is critical for load transfer from the matrix to the particle. Nanotubes can be dispersed in the polymer matrices using different techniques such as melt mixing, solution processing, or in situ polymerization. An alternative and innovative strategy relies on solid-state mixing at near room temperature, which ought to involve an efficient mixing of two or more species by mechanical milling, avoiding high temperatures and solvents. High Energy Ball Milling (HEBM) is an effective unconventional technique currently used in material synthesis and processing [12]. It consists of repeated events of energy transfer, promoted by the milling device, from the milling tools (generally balls) to the milled powder. Recently it has been proved that HEBM on polymeric materials can help obtaining materials with new characteristics, which can be barely achieved through other conventional processes [13e15]. Uniaxial alignment of CNT using magnetic [16], shear fields [17], casting-drawing and gel spun processes [18,19] has been the focus of several recent studies. Linear Low Density Polyethylene (LLDPE) is a commodity polymer, cheap, versatile, with well-known applications and commercial uses in a variety of forms, including fibers. Thus, enhancing of LLDPE properties through the dispersal and alignment of CNT should be of significant interest in order to expand its application fields. Very few papers report the preparation and characterization of Polyethylene/Carbon Nanotubes fibers. In particular, only few papers present a detailed
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
investigation about the arrangement of both the PE chains and the carbon nanotube in melt drawn oriented composite films [9e11,17]. In the present work we performed Multi walled Carbon Nanotubes dispersion an LLDPE matrix through HEBM at room temperature in the dry state with no chemical or physical treatment of the nanotubes. Fibers with different draw ratios were obtained by drawing at 80 C the composite films. The effect of different weight percentage of nanotubes and draw ratio was analyzed as function of the morphology, thermal, mechanical and electrical properties. 2. Experimental 2.1. Materials Multi walled carbon nanotubes (CNT) were purchased by Nanocyl (Belgium) (NC 3100). They were synthesized by catalytic carbon vapor deposition (ccvd) process. By thermogravimetric analysis (TGA) the carbon purity was found higher than 95% whereas the metal oxide impurity was less than 5%. The specific surface area determined with the BET method was around 250e300 m2/g. Linear Low Density Polyethylene (FlexireneÒ CM50) was supplied from Polimeri Europa (Italy) in ultrafine powder, form with a melt flow index MFI ¼ 4.1 g/10 min (at 190 C, 2.16 kg).
1125
stirring 1.0 g of potassium permanganate in a solution of 95 mL of sulfuric acid (9597 vol%) and 48 mL of orthophosphoric acid (85 vol%). Subsequent to the etching treatment, a first washing was done using a cold mixture of 2 parts by volume of concentrated sulfuric acid and 7 parts of distilled water. A second washing was performed with 30 vol % aqueous hydrogen peroxide to remove any manganese dioxide and, finally, a 3 times washing using distilled water. Before the SEM analysis the sample was kept under vacuum for 2 days at ambient temperature. The samples were finally covered with a 250 Å thick gold film using a sputter coater (Agar mod. 108 A) and observed with SEM. IR spectra were collected by means of an FTIR Spectrometer M2000 FTIR (by Midac Co.). For each sample, the average of 32 scans was used, working at a resolution of 2 cm1. The scan wave-number range of the collected spectrum was 4000e400 cm1. In order to evaluate the molecular orientation of the samples, the IR beam was polarized with the polarization axis parallel and orthogonal to draw direction [20]. Polarization of the beam was done by a zinc selenide wire grid polarizer. The average Hermans orientation factor fav can be determined by measuring the dichroic ratio D ¼ (Ap/Ao) where Ap is the absorbance for the plane parallel to the draw direction and Ao the absorbance for the polarization plane orthogonal to the draw direction. The Hermans factor, assuming uniaxial orientation, is related to dichroic ratio at wavenumber n by [21]:
D1 Dþ2 n
2.2. Composite preparation (HEBM process)
fi ¼ Kn
CNT powders and Polyethylene, as received, were milled in the solid state in a Retsch (Germany) centrifugal ball mill (model PM100). The milling process was carried out in a cylindrical steel jar of 50 cm3 with 5 steel balls of 10 mm of diameter. The rotation speed used was 650 r.p.m. and the milling time was fixed to 60 min. HEBM of powders constituted by organic polymers and fillers has been proved to be an alternative and efficient technique to produce novel composites with high performances [13e15]. During the milling the carbon nanotube bundles crack, and “intimate mixing” are promoted [15]. In these experimental conditions, five series of composites LLDPE/CNT with 1, 2, 3, 5, 10 wt/wt% of carbon nanotubes were prepared. An additional LLDPE sample to be taken as a reference was also milled in absence of filler.
where Kn is an auxiliary variable that in our case was assumed to be 1. The absorption bands at 730 and 719 cm1 were employed to evaluate the orientations of the crystal a-axis and b-axis, respectively [20]. The IR technique gives fa and fb, from which fc can be evaluated by means of equation (2):
2.3. Samples preparation The LLDPE/CNT mixtures and the pure milled LLDPE were molded in a hot press (Carver Inc.) at 170 C forming 250 50 mm thick films, which were rapidly quenched in a watereice bath (0 C). From the obtained films were obtained strips large 1 cm. The strips were put between the clamps of an Instron Dynamometer (Mod 4301), equipped with a temperature chamber. The temperature was fixed to 80 C. As soon as the temperature was stabilized, the upper traverse of the dynamometric apparatus was moved at a speed rate of drawing of 10 mm/min. Films were drawn at different draw ratios l ¼ l/lo, where l is the final length and lo the initial length. The selected values of the draw ratios (followed by a display) for all the composites were l ¼ 6; 8; 10. In the following, the resulting samples will be coded as follows: PEXlY where X ¼ 1, 2, 3, 5, 10 is the weight percent of CNT and Y is the draw ratio of the samples. 2.4. Methods of investigation Scanning electron microscopy (SEM) analysis was performed with a LEO 1525 microscope. The etching mixture was prepared by
fa þ fb þ fc ¼ 0
(1)
(2)
Differential scanning calorimetry (DSC) analysis was carried out on samples with a mass ranging between 5 and 7 mg. The tests were carried out by means of a DTA Mettler Toledo (DSC 30) under nitrogen atmosphere. The samples were heated from 25 C to 200 C at 10 C/min. To ensure reliability of the data obtained, heat flow and temperature were calibrated with standard materials, indium, and zinc. The percent crystallinity was calculated by taking the specific heat of fusion of perfectly crystalline PE to be 290 J/g [22]. Thermogravimetric analysis (TGA) was carried out with a Mettler TC-10 thermobalance. Polymer composites samples were heated from 25 to 800 C at 10 C/min heating rate under air flow. The weight loss was recorded as function of temperature. The elastic moduli were evaluated using an Instron Dynamometer (Mod 4301), and derived from the initial part of the stressestrain curves, giving to the sample a deformation of 0.1%. The experiments were conducted at room temperature with the deformation rate of 2 mm/min. The initial length of the samples was 10 mm. The data were averaged on five samples. The electrical conductivity was measured with a Keithley 6517A source measurement unit in a two-probe resistance measurement configuration. The sample thicknesses were carefully measured by a micrometer, whereas the length and the width were 3 and 50 mm, respectively. The electrical conductivity was measured in the voltage range 10 þ 10 V. All the samples showed a linear behavior of the current (I) versus the applied voltage (V). The electrical conductivity was calculated by using the equation (3):
s¼
L 1
sW R
¼
L Imeasured
sW Vapplied
(3)
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where R, s, W, L are the resistance, the thickness, the width and the length of the specimens respectively. 3. Results and discussion 3.1. Morphological analysis Fig. 1 reports the SEM images of some etched samples. The undeformed sample (Fig. 1a) shows the classical spherulitic morphology. A uniform morphology with structure aligned randomly along the sample surface can be observed. Under high uniaxial drawing in the solid state, the original morphology is transformed into highly oriented microfibril morphology (Fig. 1b). The arrow in the picture represents the drawing direction of the film during stretching. The deformation processes produce the shear of the stressed spherulitic lamellae into crystal blocks via chain slip. These blocks rotate and decrease in width until microfibrils of alternating crystal and amorphous regions are formed. Ward [21,23] proposed a morphological model, which considers that the drawn morphology consists of stacks of short lamellar-type crystallites linked by intercrystalline bridges. Since the crystalline phase is essentially continuous in the draw direction, the modulus of the drawn materials is expected to be quite high, whereas the degree of crystallinity is expected to increase with draw ratio. The presence of CNTs in the polymer matrix seems to enhance this deformation process (Fig. 1c). In fact, the oriented composite samples show a clear fibrillar morphology with numerous microcrazes. The CNTs are confined between the fibril and clearly show a high degree of alignment in the drawing direction. Such morphology has a strong influence on the thermal, mechanical and electrical properties of the resulting samples. 3.2. Detection of molecular orientation The orientation of the structural units in a solid polymer, as a result of various forming processes such as drawing and extrusion, may have a profound influence on the macroscopic physical properties of the material. Several experimental methods have been employed to characterize the orientation of crystalline and semicrystalline polymers. Birefringence, Nuclear Magnetic Resonance, X-ray diffraction, Polarized Raman spectroscopy and Infrared dichroism are among the most popular methods used to characterize different aspects of molecular order [21]. Each of these techniques has inherent advantages and disadvantages when compared with one another. Infrared spectroscopy (FTIR) and birefringence measurements often cannot be used for thick polymer specimens. In contrast, Raman spectroscopy provides an effective means for the study of samples that do not efficiently transmit radiation, provided care is taken to minimize polarization scrambling by the sample [24]. A second advantage of Raman spectroscopy is the ability to determine both the second and fourth coefficients of the orientation distribution function, whereas infrared dichroism and birefringence measurements are sensitive only to the second moment of the expansion [25,26]. Unfortunately, Raman spectroscopy is more complex from both an experimental and theoretical point of view when compared to infrared dichroism and birefringence. Not surprisingly, in spite of the usefulness of the method, its practical applications have been limited to a few specific materials, for which the molecular structures are relatively simple and which polarizability tensors were already known [27]. Due to the presence of not well characterized interactions of the carbon nanotubes with the scattering nature of Raman spectroscopy, in this paper, only the
Fig. 1. SEM microphotographs showing the microstructure of the samples: (a) PE0l0; (b) PE0l8; (c) PE3l8. The arrows show the draw direction.
FTIR technique was employed to determine the molecular orientation of samples. Detailed analyses are available in literature on PE regarding the band assignments, using FTIR spectroscopy combined with
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
polarized infrared beam [20,28,29]. In particular, the band doublet observed in the spectrum of polyethylene in the range 710e 740 cm1 has been analyzed in order to characterize the orientation of the crystal unit cell as a function of the deformation [29]. This band doublet has been assigned to in- and out-of-phase CH2rocking vibrations, respectively, of the crystal phase with the 730 cm1 band polarized along the crystallographic a-axis and the 719 cm1 band polarized along the b-axis [30]. Fig. 2 shows the Hermans factor obtained with equation (1)e(2) as function of the CNT content for the LLDPE composites analyzed in this study. The undeformed samples were also reported as reference in the same Fig. 2. By changing the CNT content, the spectra corresponding to the undeformed samples exhibit a similar dicroic ratio because they have random molecular orientation. As strain was applied to the sample and the strain increased up to l ¼ 6 the changes in the orientation function indicate that there is an appreciable modification in the mean orientation of the crystalline phase. In Fig. 2 the relative molecular orientation is evidenced to increase with increasing CNT content for all the composites. This result indicates that the orientation of the crystallites changes with the extension ratio and the CNT content. In the range 5e10% of CNT all the curves of fc level off, reaching values close to unity. This behavior reflects a typical transformation from spherulitic to fibrous structure during elongation. The transitions in the orientation function curves may arise from yielding and perhaps other morphological changes occurring during straining [31]. It is well known that during the stretching of LLDPE between the glass transition and melting temperature (which is the case here) the amorphous phase will orient itself as the crystalline regions rotate and orient themselves in the direction of stretching, as a result of this process the macromolecules can be oriented by a rotation of lamellae or by unfolding [21].
-0.20 -0.25
=0 =6 =8 = 10
-0.30
fa-0.35 -0.40 -0.45 -0.50
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] 0.3 0.2 0.1
=0 =6 =8 = 10
0.0
fb -0.1 -0.2 -0.3 -0.4 -0.5
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] 1.0
0.8
3.3. DSC analysis The DSC curves obtained for some selected samples are shown in Fig. 3. The unoriented samples display a single broad peak. It is clearly the results of two endothermic events of different intensity. Multiple melting features are characteristic of commercial LLDPEs, due to the presence of a broad distribution of crystal sizes [32] that extends over a broad range of temperatures. Upon stretching, the sample shows (Fig. 3a) a narrower melting peak and at higher temperature compared to that of the isotropic undeformed samples. This phenomenon is associated with the molecular orientation of the structure into the straining direction. For the sample with the higher draw ratio (l ¼ 10), the melting feature tends to move toward lower temperature and the maximum of melting shifts to lower temperature. This behavior can be attributed to a decrease in crystals size, or to an improvement in their perfection drawing inducted, or to the existence of a polymorphic transition occurring during heating [33]. This structural effect is not visible for lower draw ratio, probably because it is overridden by the molecular orientation effect and the disruption of the crystalline phase. Similar behavior in the melting process was observed also in presence of CNTs (Fig. 3b): an increase in draw ratio produces an increase in the melting peak temperature, followed by a decrease at higher draw ratio. However, this second stage seems to be anticipated with respect to the pure polymer. At constant draw ratio, also the CNT content shows a significant effect on the melting behavior of LLDPE. Fig. 3c shows that for the lower CNT content there is an increase in melting temperature with increasing draw ratio, but for higher CNT content the melting temperatures decrease slightly.
1127
0.6
fc
= 10 =8 =6 =0
0.4
0.2
0.0
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] Fig. 2. Orientation function as function of CNT content and draw ratio, obtained with IR analysis of all investigated samples.
It is worth noting that solid-state drawing at 80 C for LLDPE is a condition for additional crystallization. The heats of fusion of unoriented and oriented samples are reported in Table 1. They are higher for the oriented samples and increase with increasing draw ratio. At relative small elongations, the crystallinity was increased due to strain induced crystallization. As the draw ratio increases, the crystallinity decreases as consequence of the fragmentation of crystalline lamellae, increasing the draw ratio the crystallinity is expected to increase again [34]. This behavior could be associated to partial melting of smallest or less perfect crystallites followed by recrystallization. Upon addition of CNT, the heats of fusion decrease slightly probably due to the hindering phenomena from carbon nanotubes.
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PE λ = 10 PE = 8 PE = 6 PE
a
100
110
120
130
140
150
160
170
180
Temperature [°C]
PE3 PE3 PE3 PE3
b
100
110
120
130
140
150
160
= 10 =8 =6 =0
170
180
Temperature [°C] PE5 PE3 PE1 PE
c
=6 =6 =6 =6
Table 1 Melting peak temperature and heat of fusion of all the analyzed samples. Sample
Tm [ C]
Dh [J/g]
Xc
PE PEl6 PEl8 PEl10 PE1 PE1l6 PE1l8 PE1l10 PE3 PE3l6 PE3l8 PE3l10 PE5 PE5l6 PE5l8 PE5l10 PE10 PE10l6 PE10l8 PE10l10
136.2 141.2 140.3 137.6 136.0 140.7 140.3 135.8 131.5 140.6 140.7 139.8 130.0 137.9 139.2 136.7 132.1 137.6 138.4 133.4
168 179 188 190 162 177 180 183 141 174 178 179 179 157 179 179 179 147 147 147
58% 62% 65% 66% 56% 61% 62% 63% 48% 60% 61% 62% 49% 54% 56% 57% 45% 51% 53% 55%
(0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2) (0.2)
(3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0) (3.0)
respectively, are to be ascribed to a two steps process of which one controlled by oxygen diffusion. In terms of solid-state oxidation of PE, the oxygen-controlled step may be related to the oxidation of the bulk, while the first stage to the oxidation of the polymer surface and the less to the ordered phase degradation. The last residue (w8%) oxidizes in a last step extended up to 550 C. The thermal stability of the composite fibers results greatly improved compared to the undeformed PE. Table 2 reports the maximum degradation temperature, Tdmax ( C), corresponding to about 50% of weight loss. Either the incorporation of carbon nanotubes into the polymer matrix or the drawing process greatly improves the thermal stability of all the samples, such improvement increases with nanotube increasing and draw ratio. The improvement in thermal stability can be attributed to several synergistic factors. During the thermal degradation, the carbon nanotubes form a shielding layer. This layer reduces the heat transmitted to the core of the sample, prevent the recombination the polymer peroxyl radicals, and form a barrier to the oxygen diffusion [36,37]. 3.5. Mechanical properties
100
110
120
130
140
150
160
170
180
Temperature [°C] Fig. 3. DSC melting endotherms of the reported samples.
3.4. Thermogravimetric analysis Fig. 4 reports the thermogravimetric curves, evaluated in air flow, for all the samples in the range 200 Ce600 C. Pure polyethylene shows an initial degradation temperature at about 320 C. Above this temperature free radicals are generated leading to sequential degradation and break-down of the main chain due to the thermal decomposition of the covalent CeC bond [35]. The two main stages of PE oxidation, centered at 405 C and 454 C
The enhancement of the mechanical properties of composites requires a high degree of load transfer between the matrix and the nanotubes. If the interfacial adhesion between the phases is weak, the nanotubes behave as holes or nanostructured flaws, introducing local stress concentrations, and the benefits of the CNT properties are lost [38]. The nanotubes must be well dispersed. In case of poor dispersion, they will fail by separation of the bundle rather than by failure of the nanotube itself, resulting in significantly reduced strength [8,9]. On the other hand, solid-state drawing of nanocomposites is largely dominated by regions of lower nanotube volume fraction and this could certainly alter both distribution and orientation of CNTs along the length of the fiber. Strength in composites is a complex issue involving load transfer, stress concentrations and defect distribution, especially in the case of fibers. Fig. 5 shows the elastic moduli, as function of CNT percentages, for all the samples at different draw ratios. The drawing process already increased the elastic modulus of pure PE with increasing the draw ratio. The elastic modulus evaluated for the undeformed PE was about 450 MPa, it results about three times increased after drawing at l ¼ 6, about five times increased at l ¼ 8, and about seven times increased after drawing at l ¼ 10. The
G. Gorrasi et al. / Polymer 52 (2011) 1124e1132
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5000
Elastic Modulus [MPa]
4500 4000 3500 3000
λ = 10 λ=8 λ=6 λ=0
2500 2000 1500 1000 500 0
0
1
2
3
4
5
6
7
8
9
10
CNT content [%] Fig. 5. Elastic moduli, E (MPa), as function of CNT content.
significant property enhancement, at low CNT percentages (wt/wt), could be attributed to highly dispersed and well aligned nanotubes and/or good adhesion between nanotubes and polymer matrix. Fig. 6 shows the elastic moduli as function of the samples draw ratios. The higher the draw ratio the lower is the improvement in the strength due to the CNT addition. It seems to suggest that the high orientation of the polymer chains enforce high alignment of the nanotubes also at low concentrations. 3.6. Electrical properties Fig. 7 reports the conductivity values, s (S/m), of the fibers, as function of CNT percentages. Percolation theory deals with the effect of varying, in a random system, the number of interconnections present. In this case the interconnections are the highly conductive nanotubes. In literature [39] was proposed an analytical model, based on the Fermi-Dirac distribution, which describes the critical insulator to conductor transition:
Fig. 4. Thermogravimetric analysis (TGA) for all the samples in the range 200 Ce600 C.
composite fibers show significantly increased strength with increasing the draw ratio, in all the investigated composition range. The enhancement is more pronounced at low nanotubes loading (i.e. 1e3% wt/wt), a further increasing in modulus is observed for 5% wt/wt CNT (more evident for samples drawn at l ¼ 6 and l ¼ 8) and a plateau value for 10% wt/wt of filler. The data suggest that the
sc log sn
s log spn ¼ ð1 þ expðtðX Xc ÞÞÞ
(4)
where sc, sn and sp are the composite, filler, and polymer conductivities, respectively, X is the CNT mass fraction, and t is an empirical parameter that leads to the change in conductivity at the percolation threshold Xc. By assuming a constant value for sn and sp, from equation (4) were obtained the best fitted values of Xc and t for the four sets of samples. The data are reported in Table 3. The percolation threshold decreases from the undeformed samples to the fibers, going from 2 to 4 with increasing the draw
Table 2 Maximum degradation temperature, Tdmax ( C), of undeformed and drawn polyethylene and all the composite fibers, evaluated by TGA. Sample Undeformed PE PE0l6 PE1l6 PE2l6 PE3l6 PE5l6 PE10l6
Tdmax ( C) 405 424 450 450 453 463 467
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
Sample
C) C) C) C) C) C) C)
Undeformed PE PE0l8 PE1l8 PE2l8 PE3l8 PE5l8 PE10l8
Tdmax ( C) 405 422 445 454 467 468 468
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
C) C) C) C) C) C) C)
Sample
Tdmax ( C)
Undeformed PE PE0l10 PE1l10 PE2l10 PE3l10 PE5l10 PE10l10
405 417 446 457 456 457 473
C C C C C C C
(3 (3 (3 (3 (3 (3 (3
C) C) C) C) C) C) C)
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G. Gorrasi et al. / Polymer 52 (2011) 1124e1132 Table 3 Percolation threshold, Xc (wt%), electrical conductivity of filler, sn (S/m), electrical conductivity of matrix, sp (S/m), and t parameters as evaluated from Equation (1), as function of draw ratio (l).
Elastic Modulus [MPa]
10000
1000
X = 10 X= 5 X= 3 X= 2 X= 1 X= 0
100 0
2
4
6
8
10
Draw ratio Fig. 6. Elastic moduli, E (MPa), as function of samples draw ratio.
ratio from 0 to 10. Such a phenomenon, already found for nanocomposites based on oriented thermoplastic polymers-carbon nanofillers [40e42], could be justified as a break-down of the conductive network, inducted by the solid-state drawing. The drawing process of LLDPE/CNT composites decreases the number of conduction pathways present above the percolation threshold with a resulting decrease in conductivity. However, we point out that the electrical conductivity of LLDPE, soon after percolation, results increased of about nine orders of magnitude at low nanotube loadings at all draw ratios. Fig. 8 illustrates the samples conductivity as function of the draw ratio. At high draw ratio the platelets are aligned parallel each other along the stretching direction and only at higher loading they form a conductive path. At higher loadings, where there are more tubeetube connections, greater anisotropy is required to break the contacts and disrupt percolation.
4. Discussion
Xc (wt%)
sn (S/m)
sp (S/m)
t
2.0 3.0 3.5 4.0
1.00E-01 1.00E-01 1.00E-01 1.00E-01
2.24E-11 2.24E-11 2.24E-11 2.24E-11
2.0 1.6 1.6 1.6
sliding or shear of the folded molecules. The intermolecular interactions (i.e. van der Waals forces) existing between polymer chains contribute to the sample morphology [43,44]. Based on the SEM experiments, it was proposed that the microstructure of LLDPE is an interconnected network of anisotropically shaped particles, connected by array of fibrils. Considering the FTIR data presented in Fig. 2 and the proposed mechanisms of deformations, the transition at 6
0.1
0.1
1E-3
1E-3
Conductivity [S/m]
Conductivity [S/m]
The initial stages of deformation involved a displacement of the spherulitic lamellae within the sample with a non-homogeneous
l 0 6 8 10
1E-5
=0 =6 =8 = 10
1E-7 1E-9
1E-11
0
1
2
3
4
5
6
7
8
9
CNT content [wt/wt %] Fig. 7. Electrical conductivity (S/m) as function of CNT content (wt/wt%)
10
X = 10% X = 5% X = 3% X = 2% X = 1% X = 0%
1E-5 1E-7 1E-9 1E-11
0
2
4
6
8
Draw Ratio Fig. 8. Electrical Conductivity (S/m) as function of draw ratio.
10
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Scheme 1. Model illustrating the structural changes during tensile deformation.
direction of loading. In that case, the effect of the CNT seems to be less important with respect to this improvement produced by the lamellae orientation.
5. Conclusions Multi walled carbon nanotubes dispersion was achieved into a linear low density polyethylene matrix through high energy ball milling technique using 1%; 2%; 3%; 5% and 10% (wt/wt) of CNT loadings. Films from composite powders were obtained on a laboratory scale and drawn at 80 C at different draw ratios l ¼ 6; 8; 10. Morphological analysis showed that under high uniaxial drawing in the solid state, the original morphology is transformed into highly oriented microfibril morphology. FTIR demonstrated that molecular orientation increases with increasing draw ratio and CNT content for all the composites. This behavior reflects a typical transformation from spherulitic to fibrous structure during elongation. DSC analysis showed that for the lower CNT content there is an increase in melting temperature with the increasing draw ratio, but for higher CNT content the melting temperatures decrease slightly. The analysis of thermal degradation in air flow showed a high improvement of thermal stability for the composite fibers, proportionally to CNT content and draw ratio. Mechanical properties showed a significantly increased strength with increasing the draw ratio, in all the investigated composition range. The enhancement resulted more pronounced at low nanotubes loading (i.e. 1e3% wt/wt) and a plateau value was reached for 10% wt/wt of filler. Electrical properties showed that the electrical conductivity of insulating LLDPE matrix, results increased of about nine orders of magnitude, with quite low nanotube loadings at all draw ratios, and the percolation threshold tends to decrease with increasing the draw ratio for a break-down of the conductive network, inducted by the solid state drawing.
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