Dynamic Mechanical and Dielectric Properties of Ethylene Vinyl Acetate/Carbon Nanotube Composites

Journal of Macromolecular Science R©, Part B:Physics, 53:496–512, 2014Copyright © Taylor & Francis Group, LLCISSN: 0022-2348 print / 1525-609X onlineDOI: 10.1080/00222348.2013.846814

Dynamic Mechanical and Dielectric Propertiesof Ethylene Vinyl Acetate/Carbon

Nanotube Composites

HELENA VALENTOVA,1 MARKETA ILCIKOVA,2 KLAUDIACZANIKOVA,2 ZDENKO SPITALSKY,2 MIROSLAV SLOUF,3

JAN NEDBAL,1 AND MARIA OMASTOVA2

1Faculty of Mathematics and Physics, Charles University, Prague, CzechRepublic2Polymer Institute, Slovak Academy of Sciences, Bratislava, Slovakia3Institute of Macromolecular Chemistry, Academy of Sciences of the CzechRepublic, Czech Republic

A comparative study of the dielectric and mechanical properties of ethylene vinyl acetatecopolymer (EVA) filled with various concentrations of pristine and modified carbonnanotubes is reported. The surface of the carbon nanotubes was modified with 4-(2-(cholesteryloxycarboxy)ethyl) phenyl to improve the interaction of the filler with theblock copolymer matrix. The improved interaction and the better dispersion of themodified carbon nanotubes (mMWCNTs) were demonstrated by a detailed study of theEVA molecular mobility through dynamic mechanical analysis and broadband dielectricrelaxation spectroscopy. The storage modulus of the nanocomposite with 6 wt.% ofmMWCNTs at −50◦C was enhanced by 103%, whereas for the nanocomposite withthe same amount of unmodified filler, the storage modulus was only enhanced by 76%compared to the pure elastomeric matrix. This difference is more pronounced in therubbery region in which the storage moduli were increased by 117% and 48% for thecomposite with the modified and unmodified fillers, respectively. The morphologies ofthe nanocomposites were studied with scanning and transmission electron microscopiesto demonstrate the dispersion of the mMWCNTs within the EVA matrix.

Keywords carbon nanotubes, dielectric relaxation spectroscopy, dynamic mechanicalanalysis, ethylene vinyl acetate, mechanical properties, nanocomposites

Introduction

Polymeric matrices are often reinforced with carbon black[1–3] silica[4] and other micro-scopic fillers to achieve the desired properties. Composites filled with nanofillers, such asmetallic nanoparticles, clays, and carbon nanotubes (CNTs), have been envisaged to havesuperior physical and mechanical properties compared to conventional fiber- or micro-particle-reinforced composites.[5–8] CNTs have been extensively studied over the two pre-vious decades. CNTs attract attention in materials science because of their exceptional

Received 9 April 2013; accepted 9 July 2013.Address correspondence to Maria Omastova, Polymer Institute, Slovak Academy of Sciences,

Dubravska cesta 9, 845 41 Bratislava 45, Slovakia. E-mail: [email protected]

496

Studies of EVA/MWCNTs Composites 497

mechanical, electrical, and thermal properties combined with their low density. CNTs arecategorized as single-wall carbon nanotubes (SWCNTs), few-wall CNTs, and multi-wallcarbon nanotubes (MWCNTs), based on the number of graphene sheets rolled to formone tube. MWCNTs have been widely used as fillers in different types of polymers in thedevelopment of high-performance composite materials.[9–13] Recently, another interestingproperty has also been discussed, namely photo-actuation[14]—the ability of a nanocom-posite with CNTs to change its shape in response to light. CNT-enriched compositesthat exhibit this behavior have opened a new direction for applications of the CNT-basedcomposites.

Ethylene vinyl acetate (EVA) is widely used as an insulating and sheath material forhigh voltage cables and in the footwear and toy industries because of its high flexibilityand chemical inertness.[15] Nanocomposites of EVA with MWCNTs are of special interestbecause the incorporation of a small amount of MWCNTs into the EVA matrix wouldbe expected to significantly enhance the thermal and mechanical properties of the EVAmatrix.[16] Nanocomposites based on EVA copolymers also represent interesting modelsystems because the polarity and degree of crystallinity of the copolymer matrix caneasily be changed by changes in the amount of vinyl acetate (VA).[17–19] It has beenshown that an increased content of amorphous VA blocks in EVA copolymers significantlyinfluenced the kinetics and morphology of the original crystalline structures, i.e., increasedcontent disrupted the crystalline regions formed by the polyethylene (PE) segments ofthe copolymer and enhanced the polar nature of the matrix caused by the acetoxy sidechain of VA. Miltner et al.[18] have studied interfacial interactions in EVA/CNT and EVA-clay nanocomposites. They stated that increased polarity of the matrix (increased VAcontent) resulted in stronger interaction with the nanoclay and also a progressive increasein the gallery spacing was reported for VA contents increasing up to 12 wt.%, but itremained constant upon further increasing.[20] The nanocomposites based on functionalizedCNTs allowed better dispersion throughout the matrix and an increasing of the surfaceavailable for interfacial interaction. Li et al.[21] examined the nonisothermal crystallizationof EVA/CNT nanocomposites. They found that the CNTs acted as nuclei and raised thecrystallization temperature of the EVA/CNT nanocomposite relative to that of the neat EVAcopolymer.

The dielectric properties of EVA copolymers have been studied since the early six-ties.[22–24] Several transitions were found in the EVA composites based on the VA content.For random EVA copolymers with less than 20 wt.% VA, three relaxation regions weredetected;[23,24] these regions correspond to the local motion of VA groups (β-relaxation), themain transition region, α-relaxation, which corresponds to the segmental motion (manifes-tation of Tg in amorphous region,[23]), and the high temperature region, which is consistentwith the crystallization of PE.

Recently, studies have focused on the development of new types of nanocompositeswith EVA matrix and on the investigation of the mechanical, electric, and dielectric be-haviors of these nanocomposites because, even at low concentrations, nanofillers, such asconductive CNTs, have substantial potential to influence the properties of the compos-ites.[25] George et al.[26] studied cured EVA copolymers filled with expanded graphite (EG),MWCNTs, and carbon nanofibers (CNFs). They found that the critical concentration ofnanofiller and the percolation threshold for electrical conductivity depended on the type offiller. The lowest concentration, 8.2 vol.%, was needed for CNFs, followed by MWCNTs at10 vol.%, and 10.9 vol.% of EG in order for the matrix to reach the percolation threshold.The role of CNTs in reinforcement and fracture toughness was studied by Liu et al.[27,28]

They observed improvement in the fracture toughness and a reinforcement effect with

498 H. Valentova et al.

the addition of modified MWCNTs to a PP/EVA blend. The relation between the electri-cal and mechanical properties of the polymer/nanotube nanocomposites was investigatedby Masenelli-Varlot et al.[29] The authors concluded that the mechanical and dielectricproperties did not depend on the same microstructural parameters. The study concludedthat the electrical properties were much more sensitive than the mechanical propertiesto mutual contact between CNTs, and this difference arose from the weak CNT/polymerinteraction.

EVA/CNT composites have potential in a wide range of industrial applications; how-ever, deeper study of the internal structure of the relations between the CNTs and theEVA host polymer and an examination of the effects of dispersion of the nanofillersis required. We used a special surfactant to modify the surfaces of the CNTs, 4-(2-(cholesteryloxycarboxy)ethyl) phenyl,[30] because this surfactant is an aliphatic, branchedchemical substance with nonpolar character, which can improve interactions with the EVAcopolymer. Thus, an impact on the crystallinity, melting temperature and the overall me-chanical properties can be expected. For comparison, composites were prepared with thesame amount of unmodified CNTs and the properties of these composites were compared.Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) wereused to examine the dispersion of the pristine and modified CNTs within the polymericmatrix.

Experimental

Materials

The EVA random copolymer, LEVAPREN 500, was supplied by Lanxess, Germany; thecopolymer contained 50 wt.% of VA, and its structure is shown in Fig. 1a. MWCNTs (purity> 95%) were provided by Nanostructured & Amorphous Materials, Inc. (USA). Based onthe data sheet, the diameter of the CNTs was 60 to 100 nm, the length was 1 to 25 μm, andthe aspect ratio varied from 10 to 400.

Modification of MWCNTs

In this study, for the preparation of the composites, both unmodified MWCNTs(uMWCNTs) and modified MWCNTs (mMWCNTs) were used. The synthesis of 4-(2-(cholesteryloxycarboxy)ethyl)phenyl–modified MWCNTs was performed in two steps;

a) b)

CH2 CH2 CH2 CH mn

O

COH3C

Figure 1. Structure of random EVA copolymers (a), and scheme of MWCNT modified with 4-(2-(cholesteryloxycarboxy)ethyl) phenyl, where n means that the modification was repeated along theCNT surface.


first 4-hydroxyethylphenyl-modified MWCNTs were synthesized according to Wu[31] viaa diazotation reaction. The detailed description, including reaction conditions and charac-terization of product, was already published by Brzezinski et al.[32] The hydroxyl groupspresent on the MWCNTs surface were further esterified with cholesteryl chloroformateunder conditions described by Mrlik et al.[30] The structure of the modified MWCNTs isdepicted in Fig. 1b.

Nanocomposite Preparation

The nanocomposites were prepared by melt mixing of EVA with modified and unmodifiedMWCNTs with a microcompounder, DSM Xplore TM 15 (the Netherlands) at 120◦C.Nanocomposites filled with 1 to 5 wt.% of uMWCNTs were mixed for 20 min at a speedof 100 rpm. For materials filled with 6 wt.% of uMWCNTs and with more than 2 wt.% ofmMWCNTs (up to 6 wt.%), the mixing speed was decreased to 60 rpm because the meltswith higher amounts of filler had high viscosities.

Characterization Techniques

Scanning electron microscopy (SEM) investigations were performed with a ZEISS EVO40HV (Germany) electron microscope. The samples, pure EVA matrix, and compositescontaining 2 and 6 wt.% uMWCNTs and mMWCNTs, were cryo-fractured in liquid nitro-gen. The fractured surface of all of the studied samples was sputtered with gold prior toobservation.

TEM was performed with a Tecnai G2 Spirit Twin 12 microscope (FEI, Czech Re-public) after the thin samples were prepared with an ultramicrotome (Ultracut UCT, Leica,Germany) under cryo-conditions (the sample and knife temperatures were −80 and −50◦C,respectively). The ultrathin sections were transferred to a microscopic grid covered with athin carbon layer to improve the stability of the sections under the electron beam duringthe TEM measurements. All micrographs were captured in bright field mode at 120 kV.

Dynamic mechanical measurements (DMA) were performed on a dynamic mechanicalanalyzer, Tritec 2000 (Triton Technology Ltd., UK). Rectangular samples were used inthe dual cantilever bending mode. The temperature dependence of the complex Young’smodulus E∗ = E′ + iE′′ (E′ is the storage modulus, and E′′ is the loss modulus), interms of the ratio tan δ = E′′/E′, was measured at a heating rate of 2◦C/min over thetemperature range from −80◦C to 180◦C with the multifrequency mode; this which affordedan opportunity to determine the temperature dependence of E∗ at five different frequencies:f = 0.31, 1, 3.1, 10, and 31 Hz item-by-item in one temperature sweep.

Dielectric relaxation spectroscopic (DRS) measurements were performed with a Novo-control Alpha N analyzer combined with a Quatro Cryosystem (Novocontrol, Germany).The complex dielectric permittivity ε∗ = ε′−iε′′ was measured over the frequency range of10−1 to 107 Hz and over the temperature range from −100◦C to 200◦C in steps of 10◦C. Thefrequency dependence of the complex permittivity ε∗, particularly the imaginary part ε′′,is composed of relaxation, ε∗

d , and conductivity, ε∗c , contributions, ε∗ = ε∗

d + ε∗c . The relax-

ation contribution ε∗d is well described by the empirical nonsymmetrical Havriliak–Negami

(HN) equation.[33]

ε∗d = ε∞ + �ε

[1 + i (f/fr )a

]−b, (1)


with five generally temperature-dependent parameters: the high-frequency (unrelaxed) valueof the real part of the permittivity, ε∞; a characteristic frequency fr, which corresponds tothe most probable relaxation time τ r (fr = (2πτr )−1); the frequency of maximal loss, fm;the relaxation strength, �ε (�ε = εs − ε∞); and the shape parameters a and b.

The contribution of the DC conductivity to the complex dielectric permittivity ε∗c is

most notably manifested by dielectric losses at low frequencies and can be described bythe equation

ε∗c = i

(2πf ε0

/σ)η

, (2)

where ε0 is the vacuum permittivity and the exponent η would have the value −1 for theideal case of purely time-independent DC conductivity σ . Generally, the DC conductivityis temperature dependent.

The temperature dependence of the frequency f m could be either described by theArrhenius formula

fm = A exp(�U/kT ), (3)

where �U is an activation energy, k is the Boltzmann constant, and A is a pre-exponentialfactor or by the Williams–Landel–Ferry (WLF) equation[34]

log

(fm

f0

)= −c1(T − T0)

c2 + T − T0, (4)

with the adjustable parameters c1 and c2; T0 is a reference temperature. For T0 = Tg (Tg

is the glass transition temperature), the fractional free volume fg at Tg and the thermalcoefficient of expansion αf can be calculated from the relations[32]

fg = 1/2.303c1, (5)

αf = 1/2.303c1c2. (6)

In the second, so-called standard form of the WLF equation, we assumed universal valuesfor the parameters c1 and c2 (c1 = 8.86 and c2 = 101.6◦C) at the temperatureTs = Tg+50◦C;these values correspond to a fractional free volume of fg = 0.025 and a thermal expansioncoefficient of αf = 4.8 × 10−4K−1

log

(fm

fs

)= −8.86(T − Ts)

101.6 + T − Ts

. (7)

The values for the reference temperature Ts and the corresponding maximum frequency f s

of dielectric loss can be determined from the experimental data if Eq. (7) is assumed to bevalid.

Results and Discussion

Morphology Study of EVA-CNT Composites by SEM and TEM

The cryo-fracture cross-sections of the EVA matrix and composite samples were studiedby SEM, as shown in Fig. 2. For nanocomposites containing 2 or 6 wt.% of unmodifiedMWCNTs, homogeneous dispersions of the uMWCNTs were observed (see Figs. 2a and


Figure 2. SEM pictures of EVA/MWCNT nanocomposites containing (a) 2 wt.% uMWCNT; (b)6 wt.% uMWCNT; (c) 2 wt.% mMWCNT; and (d) 6 wt.% mMWCNT.

2b). However, the average distance between individual CNTs (up to 10 μm) and the brittlefracture surface suggested that (i) at microscopic level the fracture was partially plasticdue to local overheating and, consequently, some of the nanotubes were hidden because ofmicroplastic deformations of the polymer matrix in the fracturing process and/or (ii) thesample also contained agglomerates, in which the nanoparticles were concentrated. Thepresence of agglomerates was subsequently confirmed by TEM analysis (see Fig. 3).

A significantly different morphology was detected in composites containing modifiedMWCNTs. The composite was evidently more brittle; the sharp fracture lines that indicate abrittle fracture were more evident at higher concentrations of mMWCNTs—compare Figs.2c and 2d—and higher concentrations of individual nanotubes were observed on the surfaceof the fracture. Therefore, the SEM results suggested that modification of the MWCNTsresulted in a more homogeneous dispersion of the nanotubes in the polymer matrix. Thishigher homogeneity resulted in higher brittleness of the composites with mMWCNTs,which was manifested by the more brittle fracture of the material. As a result, a higherconcentration of individual, nonagglomerated MWCNTs could be observed on the fracturesurface of the EVA/mMWCNT composites. The concentration of individual nanotubeswas not decreased by agglomeration, and the visibility of the individual nanotubes on thesurface was not limited by the plastic deformations of the polymer matrix.

TEM investigation can provide a much greater in-depth view of the dispersion ofthe MWCNTs in the polymeric matrix. As shown in Figs. 3a and 3b, EVA compositeswith uMWCNTs contained not only individual nanotubes (Fig. 3a) but also nanotubeagglomerates (Fig. 3b), whose number increased as the MWCNT concentration increased.

In EVA composites containing both 2 and 6 wt.% mMWCNTs, well-dispersed CNTswere detected, as illustrated in Figs. 3c and 3d. The size of a few agglomerates of themMWCNTs (not shown) was 3 to 6 μm for the composite of EVA/6 wt.% mMWCNTs.

In general, good dispersion was obtained for the modified MWCNTs because of thecovalent surface modification with the cholesteryl groups; this modification was performed


Figure 3. TEM pictures of EVA/MWCNT nanocomposites containing (a) 2 wt.% uMWCNT; (b)6 wt.% uMWCNT; (c) 2 wt.% mMWCNT; and (d) 6 wt.% mMWCNT.

to improve the interface interactions of the filler and the EVA matrix.[14] Compared to thearomatic structure of the surfaces of the CNTs, the cholesteryl group has an aliphatic tail;therefore, improved compatibility with the ethylene acetate chains is expected. Moreover,the cholesteryl group is large and is branched; this branching (CH3 groups) should enablemultiple interactions and thus better entanglement with the chains of polymer matrixcompared to that of the linear aliphatic chains typically used to modify CNT surfaces.[13]

The entanglement of the polymer chains of the matrix with the cholesteryl group shoulddecrease the mobilities of the chains and improve stress transfer at the interface. Thesegmental immobilization is assumed to result in the previously reported enhancement ofthe elastic modulus and the strength of the composite material.[30,35,36]

Mechanical Properties

Figure 4a shows the temperature dependence of the storage and loss Young’s moduli andtanδ measured for neat EVA at five different frequencies. The same measurements wereperformed for all of the prepared composites, and Fig. 4b presents the measurement for


105

106

107

108

109

105

106

107

108

109

-50 0 50 100

10-1

100

0.31 Hz 1 Hz 3.1 Hz 10 Hz 31 Hz

E' [

Pa] (a)

E" [

Pa]

tan

δ

T [°C]

105

106

107

108

109

105

106

107

108

109

-50 0 50 100

10-1

100

(b)

0.31 Hz 1 Hz 3.1 Hz 10 Hz 31 Hz

E' [

Pa]

E" [

Pa]

tan

δ

T [°C]

Figure 4. Temperature dependence of the storage modulus E′, and loss modulus E′′, and loss factortanδ for sample (a) EVA without filler and (b) EVA with 2 wt.% of mMWCNT measured at indicatedfrequencies.

the EVA composite containing 2 wt.% of mMWCNTs. The storage modulus of both ofthe compared materials decreased as the material passed through the glass transition. Si-multaneously, the tanδ plots showed clear peaks representing the glass transition region,characterized by the temperature Tg. As expected, the maximum of the tanδ peak wasfrequency dependent; at 0.31 Hz, the maximum of tanδ was at −20◦C for both samples,whereas for the higher frequency, 31 Hz, the maximum was at −13◦C. The Tg temperature,which can be determined from DMAs to be the peak temperature of tanδ at a frequencyof 1 Hz was −19◦C for the neat EVA matrix and −18◦C for EVA filled with 2 wt.% ofmMWCNTs. The secondary tanδ peak had a maximum at a temperature of approximately35◦C and was well developed at lower frequencies; this peak is associated with the crys-talline ethylene regions in the EVA copolymer.[22] The magnitude of this peak is higher ifthe filler was added to the EVA matrix than if neat EVA was used.

Both the uMWCNT and mMWCNT fillers had only a small impact on the temperaturedependence of the dynamic mechanical properties and glass transition temperatures, Tg, ofthe EVA nanocomposites (Fig. 5). In the glassy state, the increase in the storage moduluswas proportional to the filler content. In the rubbery state, above Tg, a small decrease of E′

and a shift of Tg to lower temperatures (a softening effect) were detected after the additionof 2 wt.% of uMWCNTs (Fig. 5a). In contrast to the softening observed at 2 wt.% additionof uMWCNTs, the expected reinforcing effect which caused an increase in E′, caused bythe addition of 6 wt.% uMWCNTs was observed in the entire temperature region comparedto neat EVA (Fig. 5b).

The addition of 6 wt.% mMWCNTs to the EVA matrix (Fig. 5b) led the storagemodulus in the glassy state at −50◦C to increase to 750 MPa (measured at a frequency of1 Hz) compared with the storage modulus of the pure EVA matrix, 370 MPa, at the sametemperature. In the rubbery region, the addition of filler had no influence on the position


105

106

107

108

109

uMWCNT 0% uMWCNT 2% uMWCNT 6%

E' [

Pa]

(a)

-50 0 50 100 150

10-1

100

T [°C]

tan

δ

20 40 60 800.40.60.81.01.21.4

105

106

107

108

109

mMWCNT 1% mMWCNT 2% mMWCNT 6%

E' [

Pa]

(b)

-50 0 50 100 150

10-1

100

T [°C] ta

nδ

E' [

MP

a]

T [°C]

Figure 5. Temperature dependence of the storage, E′, and loss factor tanδ for neat EVA matrix andsamples with indicated concentration of filler (a) uMWCNT and (b) mMWCNT. Insert representsdetails of storage modulus E′ dependencies in temperature region from 20 to 80◦C, measured atfrequency 1 Hz.

of the tanδ peak (i.e., on the glass transition temperature Tg). However, the presence of thefiller led to an increase of the elastic modulus E′ above Tg (in the rubbery region). Themodulus E′ at 50◦C increased from 0.20 MPa for neat EVA matrix to 0.68 MPa for 6 wt.%uMWCNT and 1.2 MPa for 6 wt.% mMWCNT, respectively. The increase in the storagemodulus at −50◦C for a nanocomposite containing 6 wt.% of mMWCNTs was 103%,and for a nanocomposite with the same amount of unmodified filler, the increase was only76%, compared with the value for the EVA matrix. In the rubbery region, at 100◦C, thisdifference was much higher: the calculated increases were 117% for the composite preparedwith 6 wt.% mMWCNTs and 48% for the composite containing 6 wt.% uMWCNTs. Thisresult means that the reinforcing effect was much more pronounced for samples filled withthe mMWCNTs than it was for samples filled with unmodified MWCNTs. This result iscaused by the improved interactions of the polymer matrix with the modified surfaces ofthe nanotubes and the improved distribution of CNTs in the EVA matrix (Fig. 5b, inset).

The amount of CNTs were found to have a significant influence on the E′ in thetemperature range from ∼ 20◦C to 55◦C, and this influence is depicted in the insert of Fig. 5b.In this temperature range, crystallization of the EVA matrix was also detected. E′ developeda plateau followed by a gradual decrease (a corresponding peak in tanδ was observed) inthis intermediate temperature interval (Figs. 4 and 5). The crystallization phenomenon hasbeen described earlier, and it was shown that the fibers added as reinforcing agents tosemicrystalline polymer composites act as nucleation centers for crystallization.[37] Untilnow, relatively little attention has been devoted to the crystallization of EVA copolymerscontaining CNTs. In [38], the effect of crystallization on such samples was studied by DSC.In these two references, for PP/CNT and EVA/CNT composites, large nucleation effectswere observed. Crystallization began at higher temperatures; however, the crystallizationrate for EVA/CNTs was lower than that of neat EVA. This behavior can be explained


by the high aspect ratio and by the strong interfacial interactions of the CNTs and thecopolymer matrix. In other words, the polymer chains were adsorbed to the CNTs surface,and this adsorption restricted the chain mobility. The results of the DMAs confirm theseconclusions (Fig. 5). Melting of the crystalline phase was detected by the peak in tanδ withthe insignificant maximum at approximately 35◦C. The peak in tanδ for the EVA/CNTscomposites was shifted by approximately 5◦C to higher temperatures compared to those ofneat EVA. The magnitude of the peak in tanδ decreases with increasing frequency.

Dielectric Properties

The structure–property relationships of the EVA/MWCNTs nanocomposites were studiedby DRS. The incorporation of CNTs should influence the mobility of the polymer chains,DRS being a suitable method to study the changes in the relaxation dynamics of the mainbackbone chains and to study the changes in the side chain relaxations and the interfacialinteractions. The changes in the relaxation processes can be observed in the peaks in thetemperature dependence of ε′′. In the studied samples, four relaxation processes were found(Figs. 6–8), and a contribution from DC conductivity was revealed. The four processeswere as follows: the α−process, which corresponds to segmental motion in the main (Tg)transition region; the α′-process, which corresponds to melting of the polyethylene portionsof the EVA; the β−process, which corresponds to the local motion of the acetoxy polargroups in the side chains; and the δ-process, which corresponds to the interfacial interactions

10 1

10 2

-50 0 50 100

10 -1

10 0

10 1

10 2

0.3 Hz 16 Hz 9x10 2 Hz 5x10 4 Hz 3x10 6 Hz

ε'

βα'-relaxation

α-relaxation

ε"

T [°C]

Figure 6. Temperature dependence of dielectric real, ε′, and imaginary, ε′′, part of complex permit-tivity measured at indicated frequencies for neat EVA matrix.


10-1 100 101 102 103 104 105 106

10-1

100

4

5

6

7

0°C-10°C-20°C

40°C

ε"

f [Hz]

α-relaxation

40°C

0°C-10°C-20°C-30°C

ε'

-40°C

(a)

10-1 100 101 102 103 104 105 106

10-1

100

4

5

6

7

α '-relaxation

60°C 110°C

ε"

f [Hz]

(b)110°C

60°C

ε'

Figure 7. Frequency dependence of permittivity, ε′, and dielectric loss, ε′′, for sample with 2 wt%of mMWCNT in the region of the main α-relaxation (a) and high temperature α′-relaxation measuredin the temperature range from 60 to 110◦C with, increasing step 10◦C (b). The numbers at curvesdenote temperatures, or boundary temperatures.

(Maxwell–Wagner–Sillars relaxation, MWS) and was detected at the highest temperaturesin the low frequency region.

Representations of the three well-resolved relaxation processes α, α′, and β are shownin Fig. 6 which displays the temperature dependence (isochronal) of the real, ε′, and

2,5 3,0 3,5 4,0 4,5 5,0 5,5

0

2

4

6 EVA mMWCNT 6%

log

f m [H

z]

1000/T [K-1]

150 100 50 0 -50 -100

β

'α

α

T [°C]

α

Figure 8. Logarithm of the peak frequencies fm as a function of inverse temperature (Arrheniusdiagram) for all the observed relaxations in neat polymer matrix EVA and for sample with 6 wt% ofmMWCNT. Lines are fit curves assuming Arrhenius (β, α′, and δ processes) and WLF (α process).


imaginary, ε′′, parts of the complex dielectric permittivity of the neat EVA at fixed frequen-cies from 0.3 Hz to 3 MHz.

Alpha Relaxations

Figure 7a shows the frequency dependence of the real and imaginary parts of the dielectricpermittivity for a sample containing 2 wt.% of mMWCNTs and measured at constant tem-perature between −40◦C and 40◦C in steps of 10◦C. Well-resolved ε′′ peaks correspondingto the main glass transition (α−relaxation) associated with cooperative segmental motionof the EVA backbone chain were observed. Corresponding to the WLF equation (eq. 4),the position of this peak shifted to higher frequencies as the temperature was increased.

To quantitatively describe the effect of the type and amount of filler added, the ex-perimental data were mathematically fit according to the Marquard criterion.[39] The ε′(ω)and ε′′(ω) experimental data were fit to the HN-function (eq. 1) to determine the frequencyof maximum loss, fm; the relaxation strength, �ε; the shape parameters, a and b; and theDC conductivity parameters. The results are summarized in Table 1 for the primary α

transition and in Table 2 for the α′ and δ relaxation and the DC conductivity. Based on theparameters obtained from the fit to the HN-function, the ability of the dipoles to respondto changes in the external field, expressed as the dielectric strength, �ε, the free volume,the coefficient of thermal expansion at 300 K and the glass transition temperature weredetermined.

To obtain the values of Tg or Ts and log fs, the frequencies of maximum absorptionwere fit according to the WLF equation in the standard form (eq. 4). As shown in Table 1,the values of Ts were practically independent of the composition of the samples exceptfor the sample with 2 wt.% of unmodified CNTs, where the Ts value was approximately3◦C lower than that of the other measured samples. The lowering of Ts can be explainedby the softening effect, which was also observed through mechanical measurements. Theposition of the ε′′ peak, characterized by the frequency fm, and the shape of the α−peakwere the same for all of the samples. This result corresponds to the observation from theDMA measurement that the CNTs do not interact with the amorphous part of the EVA, andtherefore, the nanotubes do not influence the Tg. The curvature of the experimental points(Fig. 8) was higher than that of the standard WLF equation; therefore, the universal form ofthe WLF equation was used to determine the free volume parameters, αf and vf (Table 1).The free volume, vf, for the unmodified MWCNTs did not depend on the MWCNTsconcentration, and for the modified MWCNTs, the vf decreased as the concentration of

Table 1Estimated WLF and HN parameters for main transition, Ts and fs—parameters of the WLFequation (4), Tg maximum of tanδ of the mechanical measurements, αf and vf are free

volume parameters, �ε, a and b HN parameters

Sample Ts [◦C] log fs Tg [◦C] αf.103 [◦C−1] vf [%] �ε a b

EVA 10 5.2 −18.9 1.19 5.0 2.1 0.57 0.62EVA/2% uMWCNT 7 5.0 −19.6 1.16 4.9 1.9 0.56 0.64EVA/6% uMWCNT 11 5.4 −18.0 1.02 5.0 1.8 0.57 0.66EVA/2% mMWCNT 10 5.2 −18.2 1.03 4.7 2.2 0.59 0.65EVA/6% mMWCNT 10 5.3 −18.3 0.96 4.6 1.9 0.59 0.65


Table 2Estimated parameters for high temperature relaxation regions α′, and δ

α′-relaxation MWS δ-relaxation

Sample �U [eV/K] �ε (300 K) a �U [eV/K] �ε (300 K) b

EVA 2.0 0.17 0.54 0.6 0.8 0.69EVA/2% uMWCNT 3.0 0.22 0.41 1.1 1.7 0.37EVA/6% uMWCNT 1.8 0.35 0.42 0.3 7.1 0.57EVA/2% mMWCNT 2.0 0.15 0.53 1.2 2.6 0.58EVA/6% mMWCNT 1.8 0.22 0.51 1.0 24.6 0.63

mMWCNTs increased, as expected. The thermal expansion coefficient, αf, did not dependon the composition of the samples.

Figure 7b, where the high temperature α′-relaxation was observed from 60◦C to 110◦C,shows that the presence of CNTs had an impact on the crystallinity of the EVA. The shapeof this absorption corresponds to the Cole–Cole distribution[40] (eq (1) with b = 1), withan Arrhenius-like temperature dependence. The magnitude of this transition was 10-foldsmaller than the main transition, as shown by the data presented in Table 2. The magnitudeof the α′-relaxation transition was affected by the filler concentrations (uMWCNT as wellas mMWCNT), and this result implies that the uMWCNTs acted as nucleation centers forcrystallization.

Beta Relaxation

The secondary β dielectric absorption was found at the lowest temperatures for all of themeasured polymers. The existence of this absorption is well known from former measure-ments;[24,41] this absorption corresponds to the local motion of the small, polar acetoxygroups of the VA side chains. The temperature dependence of the β absorption followedthe Arrhenius equation (2) with an activation energy ≈ 0.3 eV and a very low magnitude.This low-temperature absorption was not influenced by the CNT content of the measuredsamples (Fig. 8). This result again confirms that neither the uMWCNTs nor the mMWCNTsinteract with the vinyl acetate parts of the EVA copolymer.

Delta Relaxation

The MWS δ-relaxation was detected at the highest temperatures in the low-frequencyregion (Table 2). The Arrhenius equation (2) describes the temperature dependence ofthe frequency fm. The shape of this absorption was well described by the Davidson–Coledistribution function[40] [a = 1 in (1)]. The magnitude of this absorption depended onthe concentration of the CNTs and increased proportionally to the concentration of thenanotubes. The change in �ε was much more pronounced in the mMWCNTs samples thanit is in the uMWCNTs composites (�ε ≈ 24.6 for 6 wt.% of mMWCNTs, �ε ≈ 7.1 for6 wt.% of uMWCNTs). The results predicted that the interfacial area was larger in themMWCNTs composites and that this larger area may result from the improved dispersionof the modified filler over the uMWCNTs. This effect is in accordance with the physicalfundamentals of MWS polarization.


DC Conductivity

All of the prepared composites and the neat EVA matrix were characterized by a DCconductivity obtained from the low-frequency portion of the frequency dependence of thecomplex permittivity with equation (2). The conductivity of the pure EVA matrix wasapproximately 10−10 S/cm. The values of the DC conductivities of all of the preparedcomposites did not significantly depend on the concentration of the CNTs; the value at300 K was 3.3 ± 0.2 × 10−4 S/cm. The DC conductivity is temperature dependent; theactivation energy of the DC conductivity lies, within the range of the accuracy of themeasurement, at ≈ 0.68 eV for all of the measured samples. More pronounced changes canbe expected in composites with filler concentrations that exceed the percolation threshold,but these composites were not the subject of this work; the highest concentration discussedhere was 6 wt.%, which is less than half of the percolation concentration suggested byZhang,[41] who determined that the percolation thresholds were between 15 and 20 wt.%.Sohi et al.[42] did not find a percolation threshold for EVA/MWCNTs composites preparedby melt mixing, whereas for the same types of composites prepared by solution casting, thepercolation threshold was achieved, and the composites attained a maximum conductivityvalue of 2.5 × 10−4 S/cm. The dominant parameter that affects the electrical conductivity ofthe composites should be the amount of conductive filler. For nanofillers, such as MWCNTs,the effects of the mixing time, rotor speed, temperature, and compression molding time mustbe investigated for every type of MWCNTs (prepared by different preparation methods,with different aspect ratios, purities, and surface modifications) to result in the preparationof composites with the desired electrical properties.

Conclusions

Composites of EVA copolymer with modified and unmodified MWCNTs were preparedby melt mixing. Modification of the nanofiller played an important role in the disper-sion of CNTs within the matrix, as indicated by SEM micrographs of the fracture sur-faces and confirmed by TEM micrographs of ultrathin sections of all of the studiedmaterials.

For nanocomposites based on EVA copolymer filled with modified and unmodifiedMWCNTs fillers, the relaxation dynamics were studied by DRS and dynamical mechanicalanalysis. DRS and DMA showed that incorporation of the CNTs did not influence theamorphous phase of the copolymer and that the shift of the Tg to lower temperatureswas observed only for the composite containing 2 wt.% of unmodified CNTs. In all othersamples, the effect on Tg was negligible in both the DRS and DMA results. This conclusionwas supported by the fact that the dynamics of the acetoxy groups, which are present inthe amorphous phase, were not influenced by the incorporation of the filler. Comparedto unmodified MWCNTs, the modified MWCNTs reinforced the EVA over the entireconcentration range studied. The difference is attributed to the improved filler–matrixinteractions caused by the cholesteryl groups, and this improvement was also confirmed byDRS. From comparison of the MWS polarizations, it can be concluded that the modifiedCNTs exhibited stronger interfacial interactions.

Funding

This work was supported by the Ministry of Education, Youths, and Sports of the CzechRepublic (grant MSM 0021620835), European project NOMS (contract no. 228916),


ORITUPOCO (ERG contract no. 231085), and the projects VEGA 2/0119/12 and VEGA2/0149/14. M. Ilcıkova thanks the Centre of Excellence SAS for Functionalized MultiphaseMaterials (FUN-MAT) for financial support. This contribution was also the result of theproject implementation of the Centre for Materials, Layers, and Systems for Applicationsand Chemical Processes under Extreme Conditions, supported by the Research & De-velopment Operational Program funded by the ERDF. Electron microscopy at IMC wasperformed with financial support from the TACR TE01020118 grant.

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Dynamic Mechanical and Dielectric Properties of Ethylene Vinyl Acetate/Carbon Nanotube Composites

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