An exfoliation of organoclay in thermotropic liquid crystalline polyester nanocomposites

Ultimate strength and initial modulus of the TLCP hybrids increased with increasing clay content and the maximum values of the mechanical properties were obtained from the hybrid contain

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An exfoliation of organoclay in thermotropic liquid crystalline

polyester nanocomposites

Jin-Hae Changa,*, Bo-Soo Seoa, Do-Hoon Hwangb

Dedicated to Prof Jung-Il Jin of Korea University, Seoul, Korea, on the occasion of his 60th birthday Received 9 October 2001; received in revised form 7 January 2002; accepted 8 February 2002

Abstract

A thermotropic liquid crystalline polyester (TLCP) with an alkoxy side-group was synthesized from ethoxyhydroquinone and 2-bromoterephthalic acid Nanocomposites of TLCP with Cloisite 25A (C25A) as an organoclay were prepared by the melting intercalation method above the melt transition temperature (Tm) of the TLCP Liquid crystallinity, morphology, and thermo-mechanical behaviors were examined with increasing organoclay content from 0 to 6% Liquid crystallinity of the C25A/TLCP hybrids was observed when organoclay content was up to 6% Regardless of the clay content in the hybrids, the C25A in TLCP was highly dispersed in a nanometer scale The hybrids (0±6% C25A/TLCP) were processed for ®ber spinning to examine their tensile properties Ultimate strength and initial modulus of the TLCP hybrids increased with increasing clay content and the maximum values of the mechanical properties were obtained from the hybrid containing 6% of the organoclay Thermal, morphological and mechanical properties of the nanocomposites were examined by differential scanning calorimetry (DSC), thermogravimetric analyzer (TGA), polarized optical microscope, electron microscopes (SEM and TEM), and capillary rheometer q 2002 Elsevier Science Ltd All rights reserved

Keywords: Thermotropic liquid crystalline polyester; Organoclay; Nanocomposite

1 Introduction

Thermotropic liquid crystalline polymers have already

been established as high performance commercial

engineer-ing polymers This is due to their speci®c chemical

structures, high strengths, high moduli, low viscosities,

and other good mechanical properties [1±4] The

structure-property relationships of thermotropic liquid crystalline

polyesters (TLCPs) have been the subject of much research

[4±7] In spite of their inferior physical strength when

compared with lyotropic liquid crystalline polyamides,

TLCPs are attracting a great deal of interest based on their

melt processability [8,9]

Although wholly aromatic TLCPs exhibit very attractive

mechanical properties, they generally have high melting

points, thus making them dif®cult to process [10,11]

Inclu-sion of ¯exible alkyl groups in otherwise wholly aromatic

polyesters not only lowers the melting point, but also

improves solubility and increases mixing entropy Thus,

despite the predictable reduction in mechanical properties,

these polyesters possess considerable advantages in some applications and show improved interfacial adhesion between the two phases [5,12,13]

Nanocomposites possess unique properties, such as stiff-ness, strength and gas permeability, for their dispersion structure [14±18] The methods used for creating nanocom-posites include in situ polymerization, solution intercala-tion, and melting intercalation [19,20] Of them, melting intercalation can be used with the most polymers, especially thermoplastic materials, but it needs a polymer that has good process properties in the melting state In recent years much attention has been paid to layered clay/polymer nanocomposites, since these represent advanced plastic materials prepared via the melting intercalation method

In our previous paper [21], large improvements were achieved in the thermal stabilities of TLCP nanocomposites

by using organo-montmorillonite This enhancement of the thermal stabilities explains reasonably well the dispersed structure of clay in the nanocomposites caused by the formation of the large aspect ratio of the clay particles For this paper, we synthesized TLCP with an alkoxy side group base on a nematic liquid crystalline phase We also examined the correlation between the thermo-mechanical Polymer 43 (2002) 2969±2974

PII: S0032-3861(02)00125-8

www.elsevier.com/locate/polymer

* Corresponding author Tel.: 182-54-467-4292; fax: 182-42-483-6155.

E-mail address: changjinhae@hanmail.net (J.-H Chang).

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properties and the clay content in TLCP nanocomposites

with variances in the dispersed morphology of the clay

particles The general goal of this work was to use a

minimum amount of clay in the hybrids and still obtain

thermo-mechanical properties signi®cantly superior to

those of matrix polymer

2 Experimental

2.1 Monomer synthesis

All reagents were purchased from Aldrich Chemical Co

Commercially available solvents were puri®ed by

distilla-tion The compound 2-ethoxyhydroquinone was

synthe-sized via a multi-step route, and 2-bromoterephthalic acid

was purchased from Aldrich Chemicals

2.2 Polymer preparation

The TLCP was prepared by direct polycondensation of

equivalent weights of the appropriate

2-ethoxyhydroqui-none and 2-bromoterephthalic acid in the presence of

thionyl chloride and pyridine The detailed procedure was

earlier described by us [22], as well as by Lenz et al [23] The ethoxy side group and Br on the TLCP not only lowers the melting point, but also improves some applications, as mentioned in the previous section The polymer formed was thoroughly washed with methanol, with dilute HCl, and then with water prior to drying at 60 8C in a vacuum oven

Inherent viscosity of the TLCP was 0.64 dL/g which was measured at 30 8C at a concentration of 0.2 g/dL solutions in

a phenol/1,1,2,2-tetrachloroethane 50/50 (v/v) mixture Fig 1 shows the thread nematic textures for pure TLCP at both 197 and 210 8C

2.3 Preparation of C25A /TLCP nanocomposites Cloisite 25A (organically modi®ed MMT; C25A) was obtained from Southern Clay Product, Co Since the synthetic procedures for C25A/TLCP nanocomposites with different weight percent (wt%) organoclay are very similar, only a representative example for the preparation

of the C25A/TLCP (2 wt%) is given 50 g of TLCP and 1 g

of C25A were dry-mixed and melt-blended at 190 8C, within the nematic region of the polymer, for 30 min using a mechanical mixer For simplicity, the hybrids will

be referred to as 0% C25A/TLCP, 2% C25A/TLCP, 4% C25A/TLCP, and so on, in which C25A and TLCP represent the organoclay and polymer components used to prepare the hybrids, respectively, and the number denotes the organo-clay weight percent in the hybrid

2.4 Extrusion The TLCP hybrids were processed for ®ber spinning to examine their tensile properties The dried blends were pressed at 160 8C, 2500 kg/cm2for a few minutes on a hot press The ®lm-type blends were dried in a vacuum oven for

24 h prior to being extruded through the die of a capillary rheometer From the capillary rheometer, the hot extrudates were immediately drawn at constant take-up speed to form extended extrudates having the same diameters The cylinder temperature of the extruder was 190 8C and the mean residence time in the capillary rheometer was about 2±3 min

To identify chemical reactions such as transesteri®cation and thermal degradation at the processing temperature, annealing was conducted for 4% C25A/TLCP hybrid at

190 8C DSC thermograms of the heat-treated hybrids are shown in Fig 2 When heat treatment time increased from

10 to 60 min at 190 8C, there were no signi®cant changes in the DSC scans Chemical changes thus do not take place to any appreciable extent at the extrusion processing tempera-ture 190 8C It was also con®rmed by 1H- and 13C-NMR spectroscopy that no detectable transesteri®cation reaction occurred in TLCP under the processing condition

J.-H Chang et al / Polymer 43 (2002) 2969±2974 2970

Fig 1 Optical micrographs of TLCP taken at (a) 197 8C and (b) 210 8C

( £ 250).

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2.5 Characterization

The thermal and the thermogravimetric analyses of

hybrids were carried out under N2atmosphere on Du Pont

910 equipment The samples were heated and cooled at a

rate of 20 8C /min Wide-angle X-ray diffraction (XRD)

measurements were performed at room temperature on a

Rigaku (D/Max-IIIB) X-ray Diffractometer, using

Ni-®ltered Co-Ka radiation The scanning rate was 28/min

over a range of 2u 2±308 Tensile properties of the

extru-date were determined using an Instron Mechanical Tester

(Model 5564) at a crosshead speed of 2 mm/min The

speci-mens were prepared by cutting strips 5 by 70 mm long An

average of at least eight individual determinations was

obtained The experimental uncertainties in tensile strength

and modulus were ^1 MPa and ^0.05 GPa, respectively

A polarizing microscope (Leitz, Ortholux) equipped with

a Mettler FP-5 hot stage was used to examine the liquid

crystalline behavior The morphology of the fractured

surfaces of the extrusion samples was investigated using a

Hitachi S-2400 scanning electron microscope (SEM) The

fractured surfaces were sputter-coated with gold for

enhanced conductivity using an SPI Sputter Coater TEM

photographs of ultrathin section polymer/organoclay hybrid

samples were taken on an EM 912 OMEGA (CARL ZEISS)

transmission electron microscope using an acceleration

voltage of 120 kV

3 Results and discussion

3.1 Dispersibility of organoclay in TLCP

The XRD patterns of C25A, pure TLCP, and their TLCP

hybrids with 2±6% C25A were represented in the region from 2u 2±158 in Fig 3 The interlayer spacing was observed in 2u 5.648 (d 18.14 AÊ) for C25A A peak was observed in 2u 4.698 (d 21.98 AÊ) for pure TLCP When the amount of organoclay increased from 2 to 6%, C25A/TLCP hybrids showed a same peak at the same posi-tion (2u 4.698) In the TLCP hybrids with 2±6% C25A,

no obvious clay peaks appeared in their X-ray diffraction curves This indicated that these clay layers were exfoliated and dispersed homogeneously in the TLCP matrix This was also direct evidence that the C25A/TLCP hybrids formed nanocomposites

Unfortunately, XRD is unable to detect regular stacking exceeding 88 AÊ One may note that the commonly used de®nition of an exfoliated nanocompoisite is based on layer spacing larger than this value In reality, it was the electron microscopic analyses that evidenced the formation

of nanocomposites

Fractured surfaces of the ®lms were viewed under SEM

A comparative analysis of the SEM photograph for TLCP hybrids with different clay content exhibiting the ®brous and platelet orientation distribution morphology including over-all projection, as shown in Fig 4

More direct evidence of the formation of a true nanocom-posite is provided by TEM of an ultramicrotomed section TEM micrographs of TLCP with different C25A content from 2 to 6% are shown in Fig 5(a)±(c), respectively The dark lines are the intersections of the clay layer of 1-nm-thickness and the spaces between the dark lines are interlayer spaces This TEM photograph proves that most clay layers of organoclay were exfoliated and dispersed homogeneously into the TLCP matrix This is consistent with the observation of XRD studies shown in Fig 3 In conclusion, we were able to successfully synthesize TLCP nanocomposites using C25A via a melting intercalation

Fig 3 XRDs of C25A and C25A/TLCP hybrids.

Fig 2 DSC thermograms of 4% C25A in TLCP hybrid annealed at 190 8C

for different times.

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method Considering the preceding results, the existing state

of clay particles could be determined to affect the thermal

behaviors and the tensile mechanical properties for each

organoclay/polymer hybrid

3.2 Thermal behaviors

The thermal properties of TLCP hybrids with different

contents of C25A are listed in Table 1 The glass transition

temperatures (Tg) of TLCP hybrids linearly increased from

92 to 98 8C with clay loading from 0 to 4 wt% and leveled

off at the content range of more than 4 wt% of organoclay

The increase in the Tgof these hybrids could be the result of

two factors First, the effect of small amounts of dispersed

clay layers on the free volume of TLCP is signi®cant, and

does in¯uence the glass transition temperature of TLCP

hybrids The second factor is ascribed to the con®nement

of the intercalated polymer chains within the clay galleries,

which prevents segmental motions of the polymer chains

DSC traces of the pure TLCP and the hybrids are shown

in Fig 6 The endothermic peak of the pure TLCP appears at

143 8C and corresponds with the melt transition temperature

(Tm) Maximum transition peaks of the TLCP hybrids

containing different clay contents in the DSC thermograms

are slightly increased to 150 8C (see Table 1) This increase

in the thermal behavior of the hybrids may result from the heat insulation effect of the clay layer structure, as well as the strong interaction between the organoclay and TLCP molecular chains The isotropic transition temperatures (Ti) of pure TLCP was virtually unchanged regardless of organoclay loading, compared with the TLCP hybrids Fig 7 shows the thread nematic textures for 2 and 6% C25A/TLCP hybrids, respectively Regardless of the clay content in the hybrids, liquid crystallinity of the C25A/ TLCP hybrids was observed when organoclay content was

up to 6%

In addition to having a higher melting point, thermal degra-dation properties of TLCP hybrids also show improvement

Fig 5 TEM photomicrographs of (a) 2%, (b) 4%, and (c) 6% C25A in TLCP hybrids.

Table 1

Thermal behavior values of C25A/TLCP hybrids

R b(%)

Fig 4 SEM photomicrographs of (a) 0% (pure TLCP), (b) 2%, (c) 4%, and

(d) 6% C25A in TLCP hybrids.

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A comparative thermal gravimetric analysis (TGA) of pure

TLCP and three nanocomposites with 2±6% C25A is shown

in Table 1 and Fig 8 TGA curves do not show weight loss

below 100 8C, as shown in Fig 8, indicating no water

remained in the samples The weight loss due to the decom-position of TLCP and its hybrids was nearly the same until a temperature of about 300 8C After this point, the initial thermal degradation temperature (Ti

D) was in¯uenced by organoclay loading in hybrids Table 1 summarizes Ti

Dof the C25A/TLCP hybrids (at 2% weight loss) increased with the amount of organoclay Ti

Dwas observed at 352±353 8C depending on the composition of the clay from 2 to 6 wt% in the TLCP hybrids, with a maximum increase of 23 8C in the case of the 6% C25A/TLCP as compared with that of the pure TLCP Weight of the residue at 600 8C increased with clay loading from 0 to 6%, ranging from 37 to 47% This enhancement of the char formation is ascribed to the high heat resistance exerted by the clay itself

Considering the above results, it is consistently believable that the introduction of inorganic components into organic polymers can improve their thermal stability on the basis of the fact that clays have good thermal stability [24,25] 3.3 Tensile properties

The pure TLCP and the TLCP hybrids were extruded through a capillary die with draw ratio (DR) 1 to examine the tensile strength and modulus of the extrudates The DR was calculated from the ratio of the diameter of the drawn extrudate to that of the extruder die

The tensile mechanical properties of pure TLCP and its hybrid ®bers are listed in Table 2 The tensile strength and initial modulus of C25A/TLCP hybrids increased with corresponding increases in the amount of organoclay The

Fig 7 Optical micrographs of (a) 2% and (b) 6% C25A in TLCP hybrids

taken at 200 8C ( £ 250).

Fig 8 TGA thermograms of C25A and C25A/TLCP hybrids.

Table 2 Tensile properties of C25A/TLCP hybrid ®bers

Fig 6 DSC thermograms of C25A and C25A/TLCP hybrids.

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ultimate tensile strength of TLCP hybrid ®bers increased as

the organoclay contents increased When the C25A was

increased from 0 to 6% in hybrids, the strength linearly

improved from 11.03 to 17.28 MPa The ultimate strength of

6% C25A/TLCP was 1.6 times higher than that of pure TLCP

The same kind of behavior was observed for the initial

moduli For example, the initial tensile modulus of 2%

C25A was 4.03 GPa, which was about 140% higher than

the modulus of pure TLCP When the organoclay in TLCP

reaches 6%, the modulus increases about 2.0 fold

(5.76 GPa) over that of the pure TLCP

This large increase in tensile property of hybrids owing to

the presence of organoclay can be explained as follows: the

amount of the increase of tensile property by clay layers

depends on the interactions between rigid, rod-shaped

TLCP molecules and layered organoclays, as well as on

the rigid nature of the clay layers Moreover, the clay was

much more rigid than the TLCP molecules, and did not

deform or relax as the TLCP molecules did This

improve-ment was possible because organoclay layers could be

highly dispersed and exfoliated in the TLCP matrix This

is consistent with the general observation that the

introduc-tion of organoclay into a matrix polymer increases its

strength and modulus [26,27]

The percent elongation at break of all samples, however,

decreases from 2 to 1% and then remains constant with clay

addition

4 Conclusions

An aromatic thermotropic LCP with ethoxy side group

was synthesized and its optical texture was nematic The

addition of 2±6% C25A to a TLCP maintains liquid

crystal-linity C25A was exfoliated and dispersed homogeneously

in the matrix polymer This was direct evidence that the

C25A/TLCP hybrids formed nanocomposites This was

also cross-checked using XRD and TEM

In general, thermal behaviors (Tg, Tm, and TDi) of the

hybrids were enhanced with increasing clay content from

0 to 6 wt% On the other hand, the isotropic transition

temperatures (Ti) of the hybrids were unchanged regardless

of organoclay loading

Hybrids of different C25A contents were extruded with

DR 1 from a capillary rheometer to investigate the

mechanical properties of the hybrids The ultimate strength

and initial modulus of the hybrids increased with increasing

C25A content When the amount of organoclay in TLCP reached 6 wt%, a 1.6-fold increase in the ultimate strength and a 2.0-fold increase in the initial modulus were obtained,

as compared with the strength and modulus of the pure polymer matrix In this system, it was found that small additions of organoclay were enough to improve the proper-ties of the matrix polymer, TLCP

Acknowledgements This work was supported by Korea Research Foundation Grant (KRF-2000-041-E00358)

References [1] Kiss G Polym Engng Sci 1987;27:410.

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[3] Dutta D, Fruitwala H, Kohli A, Weiss RA Polym Engng Sci 1990;30:1005.

[4] Joseph EG, Wilkes GL, Baird DG Polymeric liquid crystals New York: Plenum Press, 1985.

[5] Heitz T, Rohrbach P, Hocker H Macromol Chem 1989;190:3295 [6] Jackson Jr WJ, Kuhfuss HF J Polym Sci A Polym Chem 1976;14:2043.

[7] Sukhadia AM, Done D, Baird DG Polym Engng Sci 1980;30:519 [8] Kenig S Polym Engng Sci 1987;27:887.

[9] Lenz W Faraday Disc Chem Soc 1985;79:21.

[10] Isayev AI, Modic M Polym Compos 1987;8:158.

[11] Chung T Plast Engng 1987;October:39.

[12] Chang J-H, Farris RJ Polym J 1995;27:780.

[13] Chang J-H, Jo B-W J Appl Polym Sci 1996;60:939.

[14] Giannelis EP Adv Mater 1996;8:29.

[15] Lagaly G Appl Clay Sci 1999;15:1.

[16] Usuki A, Koiwai A, Kojima Y, Kawasumi M, Okada A, Kurauchi T, Kamigaito O J Appl Polym Sci 1995;55:119.

[17] Yano K, Usuki A, Okada A, Kurauchi T, Kamigaito O J Polym Sci A Polym Chem 1993;31:2493.

[18] LeBaron PC, Wang Z, Pinnavaia TJ Appl Clay Sci 1999;15:11 [19] Yang F, Ou Y, Yu Z J Appl Polym Sci 1998;69:355.

[20] Kato M, Usuki A Polymer-clay nanocomposites, Wiley Series in Polymer Science New York: John Wiley & Sons, 2000 Chapter 5 [21] Chang J-H, Park D-K Polymer (Korea) 2000;24:399.

[22] Chang J-H, Jo B-W, Jin J-I Korea Polym J 1994;2:140.

[23] Lenz RW, Frukawa A, Bhowmik P, Go RO, Majusz J Polymer 1991;32:1703.

[24] Petrovic XS, Javni I, Waddong A, Banhegyi G J Appl Polym Sci 2000;76:133.

[25] Frischer HR, Gielgens LH, Koster TPM Acta Polym 1999;50:122 [26] Wang Z, Pinnavaia TJ Chem Mater 1998;10:3769.

[27] Li J-X, Wu J, Chan C-M Polymer 2000;41:6935.

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Macromolecular Nanotechnology – Short communication

Critical aspects related to processing of carbon nanotube/unsaturated thermoset polyester nanocomposites

A Tug˘rul Seyhan a, Florian H Gojny b, Metin Tanog˘lu a,*, Karl Schulte b

b

Polymer Composites, Technische Universita¨t Hamburg-Harburg (TUHH), Denickestrasse 15, 21073 Hamburg, Germany

Received 17 June 2006; received in revised form 16 August 2006; accepted 14 November 2006

Abstract

Carbon nanotubes (CNTs) have outstanding mechanical, thermal and electrical properties As a result, particular inter-est has been recently given in exploiting these properties by incorporating carbon nanotubes into some form of matrix Although unsaturated polyesters with styrene have widespread use in the industrial applications, surprisingly there is

no study in the literature about CNT/thermoset polyester nanocomposite systems In the present paper, we underline some important issues and limitations during the processing of unsaturated polyester resins with diﬀerent types of carbon nano-tubes In that manner, 3-roll mill and sonication techniques were comparatively evaluated to process nanocomposites made of CNTs with and without amine (NH2) functional groups and polyesters It was found that styrene evaporation from the polyester resin system was a critical issue for nanocomposite processing Rheological behaviour of the sions containing CNTs and tensile strengths of their resulting nanocomposites were characterized CNT/polyester suspen-sions exhibited a shear thinning behaviour, while polyester resin blends act as a Newtonian ﬂuid It was also found that nanotubes with amine functional groups have better tensile strength, as compared to those with untreated CNTs Trans-mission electron microscopy (TEM) was also employed to reveal the degree of dispersion of CNTs in the matrix

Keywords: Carbon nanotubes; Thermosetting resin; Mechanical properties; Viscosity

1 Introduction

Scientiﬁc and industrial eﬀorts have been recently

focused on nanotechnology and nanomaterials

Nanomaterials are exhibiting some superior

proper-ties, as compared to their micro or macro size

coun-terparts Carbon nanotubes (CNTs) are composed

of thin tubes with diameters of only a few nanome-ters, but a length of few microns They exhibit higher aspect ratio, extraordinary mechanical, thermal and electrical properties, which make them prime candi-dates as reinforcing constituents in various polymers for the production of nanocomposites Although there is a number of work published[1–5] on CNT reinforced polymer composites, realization of the expected enhancement in the properties of the com-posites, such as mechanical properties has not entirely been established so far This is because of

doi:10.1016/j.eurpolymj.2006.11.018

750 7890.

European Polymer Journal 43 (2007) 374–379

www.elsevier.com/locate/europolj

EUROPEAN POLYMER JOURNAL

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the fact that nanotubes have strong tendency to

exist in agglomerated form via their huge surface

area, which leads to non-homogeneous dispersion

and random distribution of the nanotubes inside

the resin Therefore, homogeneous dispersion of

CNTs in the polymer matrix is one of the key

fac-tors to enhance mechanical properties of the

com-posites [1–9] The common dispersion techniques

for processing CNT/polymer composites have been

direct mixing and sonication [1–15] In addition,

Gojny et al [6] showed that the utilization of

3-roll-milling, which applies intensive shear forces

on the processed compounds, is an appropriate

technique to exfoliate and disperse carbon

nanotubes in an epoxy resin They also concluded

that 3-roll-milling technique provided a better

dis-persion of CNTs in the epoxy resin resulting in

higher mechanical properties, as compared to those

prepared by sonication Furthermore, besides the

physical approaches for the CNT dispersion, some

other attempts including the use of surfactants and

chemical functionalization of the CNT-surfaces

have been made in order to alter the degree of

dispersion and to tailor the interface between the

matrix and carbon nanotubes In the near future,

the further development in chemical

functionaliza-tion of nanotubes may be the key challenge for

advanced nanocomposites with the desired

proper-ties Consequently, it is obvious that a better

understanding of the relationship between

process-ing, interfacial optimization, surface chemistry and

composite properties is necessary for the potential

future applications of CNTs in polymer matrices

Unsaturated polyesters (UP) with good cost/

property relation have been the most commonly

employed matrix materials for glass ﬁber reinforced

polymer composite parts UP based materials have

been utilized in many applications including

auto-motive, construction, transportation, storage tanks

and piping industry Unsaturated polyesters become

insoluble and infusible by crosslinking with a

mono-mer, which is usually styrene The miscibility of the

resin and the styrene depends on the resin

composi-tion Commercial polyester resins contain about

30–40% by mass of styrene Polyester resins are

ver-satile, quick curing, and have a long shelf life at

room temperature The disadvantages of these

ther-moset resins are self-polymerization at higher

tem-peratures and signiﬁcantly higher cure shrinkage,

as compared to epoxy Despite the fact that

polyes-ter resins have been commonly employed in many

industrial applications, to our knowledge, there is

no reported work in the literature on the processing and properties of CNT/polyester systems Thus, CNTs have a great potential to improve the proper-ties of a low cost resin like polyester at very low ﬁller content and to induce new characteristics such as electrical conductivity In this paper, we address some critical aspects on the processing of CNT/ polyester nanocomposites prepared with the use of 3-roll-milling and also sonication techniques Trans-mission electron microscopy (TEM) was employed

to reveal the degree of dispersion of carbon nano-tubes with and without functional groups in the involved resin Some rheological and mechanical properties of the composites are also discussed

2 Experimental details

An isophtalic commercial unsaturated polyester resin Cam Elyaf 266 with 35 wt.% of styrene was obtained from CAM ELYAF Inc, Turkey Also, spe-cial polyester resin blends, composed of an allylic based polyester resin Poliya 240 with negligible amount of styrene and Poliya 420 without any sty-rene were obtained from POL_IYA POLYESTER Corp., Turkey Double-wall carbon nanotubes (DWCNT) and multi-walled carbon nanotubes (MWCNT) with and without amine functional group (NH2) produced by chemical vapor deposition (CVD) were obtained from Nanocyl (Namur/Bel-gium) and used as additives in the involved resin sys-tems DWCNTs and MWCNTs have average diameters of 2.8 and 15 nm, respectively, with a length of approximately 50 lm Cobalt naphthanate (CoNAP) and methyl ethyl ketone peroxide (MEKP) were used as an accelerator and initiator, respec-tively, to polymerize the resin suspensions that con-tain various amounts of CNTs

To prepare CNT/polyester nanocomposites, the ﬁrst approach was the utilization of the 3-roll-mill-ing process, successfully employed to process epoxy resins, employing to a commercial unsaturated polyester resin In this manner, the samples were prepared under excessive shear forces for the disper-sion of 0.1, 0.3, and 0.5 wt.% of carbon nanotubes

in Cam Elyaf 266 resin, setting the dwell time of CNTs/polyester suspension on the rolls for about

2 min The resin suspensions were polymerized with the addition of 0.3 wt.% of CoNAP and 1 wt.% of MEKP into the system During the application of this technique, we have experienced some diﬃcul-ties The major concern was the styrene evaporation from the polyester resin during the processes, which

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caused a dramatic increase of the viscosity Styrene

evaporation was accelerated due to heat occurred

on the rolling mills due to higher shear eﬀect The

polyester resin with high viscosity stacked on the

rolls and it caused some diﬃculties for the collection

of the resin, due to the uncontrolled styrene

evapo-ration, and thus the ﬁnal styrene compositions

within the resin blends were unknown

Alterna-tively, the sonication method was employed with

the same CNT/resin systems Some problems with

the sonication method similar to 3-roll-milling

pro-cess were observed Even though the sonication

bath was cooled by water, the local heating due to

energy created within the resin system, caused

sty-rene evaporation from the polymer suspension,

leading to a more viscous resin In addition, it was

observed that nanotubes were agglomerated in the

volumes closer to the tip of the sonicator Van der

Waals attractive forces between the CNT-surfaces

are known to be sensitive to heat, so increasing

agglomeration occurred[2,6] To overcome the

dif-ﬁculties associated with styrene evaporation, we

switched to a resin system, containing negligible

amount of styrene (Poliya 240) and non styrene

(Poliya 420)

With the corresponding novel polyester resin

sys-tems nanocomposites were prepared by setting the

appropriate gelation time and viscosity for the

3-roll-milling processing The styrene was added to

the system after 3-roll-milling After some

experi-mental trials, polyester resin blends were formulated

based on 45 wt.% of Poliya 420, 30 wt.% of Poliya

240 and 25 wt.% of styrene with the presence of

0.2 wt.% of CoNAP and 1.5 wt.% of MEKP The

polymer mixture to be used during 3-roll-milling

process was prepared by hand-mixing of two types

of the polyester resin at the given ratio for 10 min

Nanocomposite samples were prepared by the

dis-persion of the 0.1, 0.3, and 0.5 wt.% of the carbon

nanotubes within the polyester resin blend After

collecting the CNT containing polyester suspension

by spatula from the 3-roll-milling, 25 wt.% of

sty-rene was added to the involved resin system The

whole system was then subjected to the intensive

mixing for half an hour using magnetic stirrer and

ﬁnally poured down into an aluminum mold and

cured at room temperature followed by post curing

in an oven at 110°C for 2 h Although Poliya 240

with a lower amount of styrene was introduced to

3-roll-milling, comparable viscosity increase of the

resin, and problems due to the stacking of the high

viscosity resin on the rolls were observed However,

note that in this second approach, we diminished diﬃculties with styrene evaporation and unknown styrene content in the ﬁnal product by using low sty-rene containing resin For that reason, in our fur-ther experimental investigations, we focused just

on investigating the properties of the nanomaterials prepared by the second approach

The dispersion of the CNTs within the composites was characterized by transmission electron micro-scopy (TEM) using a Philips EM 400 at 120 kV acceleration voltages The ultra thin TEM samples with a thickness of 50 nm were prepared by ultra-microtome cutting at room temperature

TA Instruments RDA III with parallel plate rhe-ometer geometry (500 lm gap, and 50 mm plate diameter) was used to analyze the rheological behaviour of the polyester suspensions with diﬀer-ent carbon nanotubes loadings Tests were per-formed in steady state modes at room temperature

in order to avoid styrene evaporation during the measurements For that reason, liquid samples were taken from the collected resin suspension from the 3-roll-mill Steady shear rates (SSS) were used to investigate the ﬂow properties of the polyester sus-pensions by considering the viscosity as a function

of increasing shear rates

Mechanical tensile properties of the composites were determined according to DIN EN ISO 527.1 Dog bone specimens were prepared by countersink-ing uscountersink-ing a Mutronic Dear Drive 2000 The tensile samples were tested using a Zwick Z010 Universal tensile testing machine at a cross head speed of

1 mm/min The elongation of the specimens during the test was also measured

3 Results and discussion The 3-roll-milling process via intensive shear forces seems to be more convenient technique than traditional ones such as sonication and direct mix-ing for the dispersion of carbon nanotubes within

a liquid polymer resin.Fig 1shows the TEM micro-graphs of MWCNTs and DWCNTs with and with-out functional groups in the polyester resin blend for 0.3 wt.% of loading MWCNTs with functional groups exhibited better local dispersion in the poly-ester matrix, as compared to DWCNTs with and without treatment In general, DWCNTs were observed to be more agglomerated form caused by their pronounced higher surface area In the litera-ture, rheological behaviour of the polymer suspen-sion was associated with the prediction state of

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CNTs dispersion within the corresponding resin

[11] Figs 2 and 3 give the viscosity as a function

of shear rate for the Poliya polyester based

suspen-sions containing MWCNTs, MWCNT–NH2,

res-pectively at diﬀerent loading rates As seen in

the ﬁgures, shear thinning behavior was observed

for the samples containing either MWCNT or

MWCNT–NH2, such that viscosity is reducing with the increase of shear rates The viscosity of polyester suspensions with MWCNT decreases sharply at 0.1 wt%, but MWCNT–NH2 has not the same behavior This might be due to the fact that nano-tubes with amine functional groups reveal better compatibility or chemical interaction with the poly-ester chains within the system Carbon nanotubes

10 -2

10 -1

10 0

10 1

10 2

Neat polyester resin 0.1wt.% MWCNT 0.3wt.% MWCNT 0.5wt.% MWCNT

Shear rate [s -1 ]

Fig 2 Viscosity of the polyester suspension with MWCNTs as

a function of shear rate.