C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

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Effect of processing technique on the transportand mechanical properties of vapour grown carbonnanofibre/rubbery epoxy composites for electronicpackaging applications

Mohsin Ali Raza a,*, Aidan Westwood a, Chris Stirling b

a Institute for Materials Research, University of Leeds, Leeds LS2 9JT, UKb Morgan AM&T, Swansea SA6 8PP, UK

A R T I C L E I N F O

Article history:

Received 4 May 2011

Accepted 9 August 2011

Available online 16 August 2011

0008-6223/$ - see front matter � 2011 Elsevidoi:10.1016/j.carbon.2011.08.010

* Corresponding author.E-mail address: mohsinengr@yahoo.com

A B S T R A C T

Vapour grown carbon nanofibre (VGCNF)/rubbery epoxy (RE) composites were produced, by

either mechanical mixing, three-roll milling (RM) or combined ultrasonication/mechanical

mixing. Incorporation of VGCNFs resulted in significant enhancements in the thermal and

electrical conductivities of the material. Appropriate selection of processing technique and

parameters can help to maximise the potential of VGCNF additions by improving their dis-

persion in the matrix. The composites produced by RM have superior transport properties

compared with those produced by other techniques. The thermal conductivity of such

composites at 40 wt.% VGCNFs reached 1.845 W/m K, a 10-fold increase compared to RE

alone. The thermal conductivity data of VGCNF/RE composites best fits to the Hatta–Taya

model. The lowest electrical percolation threshold is at 2 wt.%, obtained for composites

produced by RM. The thermal conductivity of VGCNF/glassy epoxy (GE) composites at

12 wt.% is 10% lower than the corresponding RE composite but its electrical conductivity

is 2 orders of magnitude higher than the corresponding RE composite. VGCNFs at

40 wt.% increase the compressive strength of rubbery epoxy by �5· but the compressive

modulus of 40 wt.% VGCNF/RE composite is 12 times lower than that of 12 wt.% VGCNF/

GE composite, demonstrating highly compliant nature of RE composites.

� 2011 Elsevier Ltd. All rights reserved.

1. Introduction

Thermally and electrically conductive polymer composite

materials are of great importance for electronics packaging

applications (EPAs), such as thermal interface materials [1],

electrically conductive adhesives [2], electromagnetic inter-

ference shielding and electrostatic dissipation [3]. Multifunc-

tionality of the polymer composites depends on the size,

shape and nature of the fillers [4]. For example, not all fillers

er Ltd. All rights reserved(M.A. Raza).

can impart both thermal and electrical conductivity to poly-

mer composites; only metallic and carbon based fillers such

as carbon nanotubes (CNTs), carbon fibres and graphite flakes

can enhance both transport properties.

Vapour grown carbon nanofibres (VGCNFs) have been

studied in recent years as fillers to improve mechanical, elec-

trical and thermal properties of the polymers [5,6]. They have

high aspect ratio and are available with diameters ranging

from 50 to 200 nm and lengths of 50–100 lm. Although the

.

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 85

properties of VGCNFs are inferior to those of CNT, their lower

cost and close resemblance to CNTs makes them attractive

for reinforcement in polymers [4].

VGCNFs are produced by catalytic vapour deposition from

hydrocarbons or carbon monoxide over a metal catalyst.

VGCNF consists of a tubular structure comprised of single

or double layers of stacked, truncated graphene cones which

intersect the central hollow along the fibre axis at highly ob-

lique angles [4,7]. Depending upon the cone angle and stack-

ing, various structures of the fibres are possible including

bamboo like [8] and cup-stacked [4]. The outer surface of

the carbon fibres consists of a layer of vapour-deposited car-

bon which thickens the fibres and bonds them together in

large aggregates. The exterior layer of the fibre is not as gra-

phitic as the interior layers and consists of disordered

embryonic graphene planes [7]. The graphitised VGCNF

exhibits very high electrical and thermal conductivity val-

ues. The electrical resistivity of VGCNF after graphitisation

at room temperature is therefore low, at about 5 · 10�5 O cm, which is close to the resistivity of graphite. The

thermal conductivity of as-grown VGCNF is reported to be

20 W/m K but this can exceed 1950 W/m K upon graphitisa-

tion [9]. The thermal conductivity of highly graphitic fibres

is the highest among all commercial fibres [4]. The major

drawback of VGCNFs is their poor tensile properties com-

pared to those of carbon nanotubes. This is probably due

to the weakness of the conic stacks under the shear which

is applied when the fibre is in tension in combination with

their large diameter which increases the likelihood of a flaw

being present in the section under test. Tibbetts and Beetz

[10] believe that the dependence of the fibre mechanical

properties on the diameter is also due to non-uniform car-

bon deposition rates along the length of the fibre during

growth. The large diameter fibres have lower modulus than

the small diameter fibres [4]. Patton et al. [11] estimated

the tensile modulus and tensile strength of Pyrograph-III

nanofibres (having diameter from 100 to 300 nm) to be in

the range of 88–166 GPa and 1.7–3.38 GPa, respectively. In

addition to the diameter, the tensile properties of the fibres

also depend on morphology and crystallinity.

VGCNF composites have been developed by dispersing

them in both thermoplastics [12,13] and thermosetting poly-

mers such as epoxy [5,11,14–16] as well as in elastomers

[17,18]. Two review articles [3,4] give a good overview of the

fabrication and properties of these composites. Among all

these VGCNF/polymer composites, the epoxy-based compos-

ites are the most promising for electronics packaging applica-

tions. Epoxy resins have secured a firm place in engineering

applications due to their ease of processing and application,

low viscosity (before curing), good adhesion, high strength,

low coefficient of thermal expansion, high thermal stability

and low cost [19,20]. The articles published on VGCNF/epoxy

composites to date have mainly addressed the effect of

VGCNF on the mechanical properties of the composites. Only

a few authors have addressed the thermal properties of these

composites. For example, Chen and Ting [21] developed

aligned VGCNF/epoxy composites for thermal management

applications using VGCNF mats. Beircuk et al. [22] compared

the thermal conductivities of single walled carbon nanotube

(SWCNT)/epoxy composites and VGCNF/epoxy composite at

1 wt.% loading of the filler. They found that, compared to

VGCNF, SWCNTs offer three times greater enhancement of

composite thermal conductivity. Patton et al. [11] reported

that the thermal conductivity of the composites ranges from

0.6287 to 0.8072 W/m K at 17–39 vol.% fibre, respectively.

Tibbetts et al. [3] reported the work of Lafdi and Maztex

who measured a thermal conductivity of 2.8 W/m K for a

VGCNF/epoxy composite at 20 wt.% loading, an increase of

1300% compared to the neat epoxy (0.2 W/m K). Choi et al.

[14], Prasse et al. [16] and Kotaki et al. [23] investigated the

electrical properties of such composites. The transport prop-

erties of the composites depend strongly on the processing

methods and conditions because these effect the filler disper-

sion, distribution and aspect ratio [4]. Therefore, it is

important to select appropriate processing routes and condi-

tions to produce VGCNF/polymer composites for different

applications.

Epoxy resins, despite their good properties, have some

drawbacks. Most of the epoxy-based composites reported in

the literature or used for engineering applications are highly

crosslinked materials, have high modulus and lack the con-

formability [24] which, for example, is required for thermal

interface applications. These highly crosslinked epoxy resins

are known as ‘‘glassy epoxy’’. The high modulus and brittle-

ness of glassy epoxy does not allow internal stresses to dissi-

pate which leads to delamination and thermal fatigue when

they are used as adhesives [2]. The addition of filler into such

epoxy resins generally further increases the modulus and re-

duces conformability. Therefore, it is highly desirable to pro-

duce composites for EPAs using a lower modulus matrix.

Although it is possible to decrease the modulus of the epoxy

resin by addition of flexibilisers [25], there is a commercially

available type of epoxy resin, namely rubbery epoxy, which

may be used as an alternative to state of the art glassy epoxy

resins to fabricate superior composites for EPAs. Rubbery

epoxy has a glass transition temperature (ca. 238 K) which is

below ambient temperature and it has a very low modulus

compared to glassy epoxy [18,26]. It is not a true elastomer

but its mechanical properties resemble those of an elastomer

to some extent. There is only one report of VGCNF/rubbery

epoxy composites and this has its emphasis on the mechan-

ical properties [18].

The main focus of this research was to produce VGCNF/

rubbery epoxy composites and to explore their potential for

EPAs. The transport and mechanical properties of the

VGCNF/rubbery epoxy composites are also compared with

those of corresponding composites produced using glassy

epoxy. The issue of fibre dispersion in the matrix was also

studied so as to determine the optimum production condi-

tions for these composites. Composites have been produced

via three different processing routes, namely conventional

mechanical mixing, 3-roll mill mixing and combined ultra-

sonication and mechanical mixing, in order to study the ef-

fect of mixing route on the dispersion and distribution of

the VGCNF filler and on the transport and mechanical proper-

ties of the resulting composites.

86 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

2. Experimental

2.1. Materials used

VGCNFs (Pyrograf III, PS and HT grade) were purchased from

Applied Sciences Inc. These fibres have diameters in the

range of 70–200 nm and lengths about 50–100 lm. These fibres

were used as-received in fabrication of composites.

Epoxy resin, EPIKOTETM EP828, was kindly supplied by

Hexion Specialty Chemicals and two aliphatic polyether-

amine curing agents, Jeffamine D2000 and Jeffamine T403

(ex Huntsman Corporation) were used in this work.

EP828 is an undiluted clear difunctional epoxy bisphenol

A/epicholorhydrin-derived liquid epoxy resin. Jeffamine

D2000 polyetheramine is characterised by repeating

oxypropylene units in the backbone. It is a difunctional pri-

mary amine with a molecular weight of about 2000. Jeffamine

T403 polyetheramine is a trifunctional primary amine having

a molecular weight of about 440. The chemical structures of

EP 828, Jeffamine D2000 and Jeffamine T403 are shown in

Fig. 1.

Rubbery and glassy epoxies were used as matrices for

VGCNF/epoxy composites. Rubbery epoxy (RE), so-called be-

cause (after curing) it has a glass transition temperature be-

low ambient temperature, was produced by mixing epoxy

resin, Epon 828, and curing agent, Jeffamine D2000, at a

weight ratio of 25:75, respectively. The rubbery nature of the

epoxy is attributed to the moderate cross-linking developed

by the Jeffamine D2000. Glassy epoxy (GE) matrix, so-called

because (after curing) it has a glass transition temperature,

Tg, of �80 �C [26], was produced by mixing Epon 828 and

Jeffamine T403, at a weight ratio of 100:42, respectively. The

glassy nature of the resultant epoxy develops due to the

extensive cross-linking of the epoxy resin by the Jeffamine

T403.

2.2. Fabrication of composites

The VGCNF/epoxy composites were prepared by three differ-

ent mixing techniques. These are described in the following

section.

2.2.1. Conventional mechanical mixing (MM)When preparing samples with minimum dimensions of

40 · 25 · 10 mm3, 40–50 g batches were prepared by mixing

Fig. 1 – (a) Chemical structure of EPIKOTETM 828 (b) Chemical stru

T403.

VGCNFs and epoxy. All of the composite dispersions were pre-

pared at room temperature. VGCNFs were dried in an oven at

80 �C for a prolonged period to remove any moisture adsorbed

on their surface. The dried VGCNFs were then mixed at

appropriate percentages with rubbery epoxy by using a

conventional mechanical mixer with a high-speed motor at-

tached to a shaft with a propeller. This was rotated in the

mixture at 4500 rpm for 15 min. These parameters had been

‘‘optimised’’ (i.e. providing material with the highest electrical

conductivity without unduly prolonged mixing). After mixing,

the batch was degassed under vacuum to remove any trapped

air and was poured into a custom-made aluminium mould.

The filled mould was again degassed for half an hour to com-

pletely remove any trapped air. The VGCNF/RE composites

were prepared with a loading of 2–15 wt.% of VGCNFs. It

was not feasible to incorporate VGCNFs at loadings higher

than 15 wt.% by MM due to the high viscosity of the resulting

dispersion. The temperature of the dispersions increased to

60 �C during mixing. This temperature is safe for the rubbery

epoxy as it does not start to cure even after prolonged treat-

ment at 80 �C. VGCNF/GE composite was prepared in the same

way using 12 wt.% of VGCNFs. Both VGCNF/RE and VGCNF/GE

were cured at 80 �C for 2 h and 120 �C for additional 3 h. Sam-

ples of ‘‘neat’’ RE and GE were also produced by the MM

method.

2.2.2. Combined sonication and mechanical mixing (CSM)In this method, VGCNFs were first ultrasonicated in acetone

at concentration of 1.5 g/100 ml for 1, 5 and 10 h to deagglom-

erate the fibres so as to achieve better dispersion in the epoxy.

After sonication, curing agent (Jeffamine D2000) was added at

an appropriate percentage into the VGCNF/acetone disper-

sion. This dispersion was kept stirring on a hot plate at

80 �C until evaporation of the acetone was complete. Once

the acetone had been completely eliminated, resin (Epon

828) was added and mixed by the MM method at 4500 rpm

for 15 min. The rest of the procedure for fabrication of com-

posites was same as described for composites produced by

MM. VGCNF/RE composites were produced at 10 wt.% fibre

loading by the CSM method.

2.2.3. Three-roll mill mixing (RM)A three-roll mill (EXAKT GmbH) was used for fabrication of

the composites. The roll mill has the ability to achieve good

filler dispersion by extremely high shearing force resulting

cture of Jeffamine D2000 (c) Chemical structure of Jeffamine

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 87

from different roller speeds and filler deagglomeration using

the crushing force of the rollers. VGCNF/RE dispersions were

first prepared by MM at 1000 rpm for 2 min. This dispersion

was poured between the feed (n1) and central rollers (n2) as

shown in Fig. 2. The mixture was transported between the

n2 and n3 rollers where it was dispersed to desired degree of

fineness. The scraper system removes the finished product

from the (n3) apron roller. The VGCNF/RE dispersions were

passed through the rolls three times in gap mode operation

and two times in force mode operation at a speed of

200 rpm. In gap mode, the rollers were operated in the first

pass with gaps of 120 and 40 lm between n1/n2 rollers and

n2/n3 rollers, respectively. This gap was reduced in the second

and third passes to 60/20 lm and 15/5 lm, respectively. In

force mode operation, the gap between the rollers was kept

at less than 5/<1 lm so that rollers could apply force (approx.

2 N/mm2) on the filler to break the agglomerates. After the

fifth pass, the VGCNF/RE dispersion was collected directly

into the custom-made aluminium mould. The remaining pro-

cedure for the fabrication of the composites was same as de-

scribed for composites produced by MM. VGCNFs were

dispersed at 2–40 wt.% in RE to form composites. There was

no significant temperature increase observed for dispersions

produced by RM due not only to the low speed of the rollers

but also to their acting as heat-sinks.

2.3. Calculation of volume fraction of fibres

The volume fraction of fibres can be calculated from weight

fractions according to following equation [27]:

Vf ¼qmWf

qf Wm þ qmWfð1Þ

and

Vm ¼ 1� Vf ð2Þ

In Eqs. (1) and (2), Vf and Vm are the volume fraction of the

VGCNF and matrix, respectively, Wf and Wm are the weight

fraction of the VGCNF and matrix, respectively and qf and

qm are the densities of the VGCNF (1.95 g/cm3) and rubbery

epoxy matrix (1.025 g/cm3), respectively.

Fig. 2 – Schematic of exact 3-roll mill (http://www.exakt.de).

2.4. Characterisation

2.4.1. Electron microscopyThe morphology of the composites was observed using a LEO

1530 field emission gun scanning electron microscope (FEG-

SEM). The images were obtained using secondary electrons

at 3 kV with a working distance of 3 mm. The samples for

SEM analysis were prepared by cooling strips of the compos-

ites in liquid nitrogen and then brittle-fracturing them. The

fractured surface of the sample was sputter-coated with a

thin layer (5 nm) of Pt/Pd alloy prior to the SEM analysis. All

of the samples studied by SEM were sectioned in such a man-

ner that a surface parallel to the direction of gravity (during

moulding) was exposed for the analysis. TEM images of

VGCNF were obtained using an FEI CM200 FEGTEM with Gatan

GIF 200 Imaging filter running at 197 kV. Samples were pre-

pared by dispersing in methanol, with a drop placed on a

holey carbon-coated copper grid.

2.4.2. Thermal conductivityThe thermal conductivity of the neat epoxy resin and com-

posites was measured by the hot disk thermal constant ana-

lyser method (Hot Disk� AB), which is a transient plane

source technique [28]. The sensor, which acts both as heat

source and temperature recorder, with a radius of 3.180 mm,

was sandwiched between the two halves of each sample.

For measurement, each sample half was prepared so as to

be 8–10 mm thick (in the direction of gravity that applied dur-

ing curing in the mould, i.e. the measurement direction) and

�20 · 20 mm in the other directions, with a flat surface. The

measurements were made by applying a heating power of

0.1–0.2 W for between 10 and 80 s, depending on the thermal

conductivity of the sample. The temperature increase of the

sample as a function of time was recorded to compute the

thermal conductivity of the sample, based on solution of

the transient heat conduction equation.

2.4.3. Electrical conductivityFor electrical conductivity measurement, approximately

cuboidal samples (�5 · 5 · 2 mm) of the composites were

placed between two copper electrodes having dimensions

slightly greater than that of the sample. The electrodes were

connected to an Agilent mulitmeter (34401A) which measured

the resistance of the sample according to the two probe meth-

od. To ensure good contact between the sample and the

copper electrodes, the samples were slightly compressed

between the electrodes. To observe the effect of orientation

of the VGCNFs in the composites, the electrical conductivity,

r, was measured through the sample, designated rz (in the

direction of gravity that applied during curing in the mould),

and along the sample, designated rx and ry, as shown in Fig. 3.

2.4.4. Compression and hardness testingCompression testing of the neat epoxy and composites was

carried out an Instron universal testing system (Model No.

3382 with a 100 kN load cell). Cuboidal samples (�10 · 10 ·8 mm3) were compressed at a strain rate of 0.5 mm min�1.

The compression tests were performed on the samples so

that compression occurred parallel to the direction of gravity

in the original curing moulds (z-direction in Fig. 3). A typical

Fig. 3 – Schematic of electrical conductivity measurement

set up.

88 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

compression test was carried out until the sample fractured.

Hardness testing of the samples was carried out using a shore

hardness tester (Zwick) and values were measured on scale A.

3. Results and discussion

3.1. Morphology

TEM images of the VGCNFs used in this work are shown in

Fig. 4(a) and (b). It can be seen that the fibre walls consist of

stacked graphitic planes in one or two distinct layers and

have cup-stacked (Fig. 4(a and b)) or bamboo-like (Fig. 4(k))

morphology.

The SEM micrographs of the 12 wt.% VGCNF/RE compos-

ites produced by MM and RM are presented in Fig. 4(c and d)

and Fig. 4(e and f), respectively. The dispersion quality of

the fibres in the rubbery epoxy composite produced by MM

is relatively poor compared to that in the composite produced

by RM. Clusters of agglomerated fibres (indicated by arrows in

Fig 4(c)) and some regions of the matrix without fibres can be

seen in the SEM image of the composite produced by MM. On

the other hand, the SEM image of the composite produced by

RM (Fig. 4(e and f)) showed no significant agglomerates but

uniform dispersion of the fibres in the matrix. In the compos-

ites produced by RM at 25 and 40 wt.% loading, the fibres

seem to be homogeneously dispersed in the matrix as shown

in Fig. 4(i–k). A few voids can be observed in the SEM images

of the composites produced by RM at 40 wt.% loading. These

voids arose due to the slightly higher viscosity of the disper-

sions at such high loading, which did not allow the matrix

to completely homogenise during the curing of the compos-

ite. However, it is believed that these voids can be easily

avoided either by introducing the dispersion into the mould

under pressure or by compression moulding of the compos-

ites. In contrast, the high viscosity of the VGCNF/GE compos-

ites, even at modest fibre loadings, results in a high void

content in these composites as can be seen from SEM images

(Fig. 4(g and h)). In all cases, it appears that there is little or no

orientation effect of the VGCNFs in the matrix and that the fi-

bres are randomly orientated in the matrix at loadings of 12,

25 and 40 wt.% to form isotropic composites. RM processing

also decreased the length of fibres significantly as observed

by SEM imaging (Fig. 4(l)) of processed VGCNFs (obtained after

dissolving the uncured 40 wt.% VGCNF/RE dispersion in ace-

tone and then evaporating the solution on an aluminium

SEM stub). Examination of the SEM image showed that the fi-

bre length after RM processing is as low as 5 lm but averages

about 20 lm.

SEM analysis of the composites therefore demonstrates

that the ability of the RM to deagglomerate the fibres and dis-

perse them uniformly is far superior to that of conventional

mixing (MM). This capability is also evidenced by the fact that

RM enabled the dispersion of VGCNFs in the matrix even at a

loading of 40 wt.%, which could not be achieved by any other

means. Indeed, this composition could also be coated easily

onto a substrate prior to curing, suggesting that the fibres

were well wetted with the resin. In contrast, it was observed

that it was difficult to produce the VGCNF/RE composite at

loadings as low as 15 wt.% of fibres by MM because the high

viscosity of the dispersion hindered processing.

3.2. Thermal conductivity and its correlation withprocessing of composites

The thermal conductivities at room temperature of the

VGCNF/RE and VGCNF/GE composites as a function of wt.%

of VGCNFs are presented in Fig. 5. The thermal conductivity

of the VGCNF/RE composites increases (approximately) line-

arly with increasing fibre loading. The thermal conductivity

of the VGCNF/RE produced by RM is proportional to the vol-

ume fraction of fibre according to the equation obtained by

linear fitting of measured thermal conductivity data versus

volume fraction of fibres: k = 6.69 Vf + 0.21 (where k is the

thermal conductivity of the composites and Vf is the volume

fraction of the fibre). The thermal conductivity of the neat

RE is 0.1769 W/m K (approximately consistent with the afore-

mentioned equation). The composite produced by RM at

15 wt.% loading has a thermal conductivity of 0.811 W/m K

which is a �4.5-fold increase cf. neat RE. The thermal conduc-

tivity of the composites with 25 wt.% (15 vol.%) of VGCNFs

reached 1.31 W/m K, representing a 7-fold increase, and at

40 wt.% (26 vol.%) reached 1.85 W/m K, representing a 10-fold

increase over neat RE.

Compared with the composites produced by MM, the ther-

mal conductivities of the corresponding composites produced

by RM are 11–38% higher, depending upon the wt.% of

VGCNFs incorporated (up to 12 wt.%). This increase can be

attributed to the improved distribution and dispersion of

VGCNF in the matrix achieved by RM, as observed by SEM

analysis. However, at 15 wt.% of VGCNFs the thermal conduc-

tivity of the composite produced by RM is only 8% higher than

the corresponding composite produced by MM. It is possible

that the intensity of shearing by MM increased due to the ob-

served high viscosity of the 15 wt.% VGCNF/RE dispersion and

that this resulted in much more deagglomeration of the fibres

and improved their dispersion in the matrix. However this

does not undermine the advantage of producing VGCNF com-

posites by RM because RM not only offers somewhat superior

dispersion of the filler in the matrix but also allows a higher

loading of the filler in the matrix to be processed without

compromising the uniform dispersion of the filler. In this

way, RM enabled production of composites with very high

thermal conductivities at 25 and 40 wt.% loading of VGCNFs.

These thermal conductivities are not only comparable to

those of many of the commercially available thermal inter-

face materials used for thermal management in EPAs but

Fig. 4 – TEM images of (a and b) single and double layer VGCNF, SEM images of (c and d) 12 wt.% VGCNF/RE produced by MM

(arrows pointing towards agglomerates), (e and f) 12 wt.% VGCNF/RE produced by RM, (g and h) 12 wt.% VGCNF/GE produced

by MM (arrows pointing toward voids) (i) 25 wt.% VGCNF/RE produced by RM, (j and k) 40 wt.% VGCNF/RE produced by RM

(arrow pointing towards bamboo-shaped VGCNF). Some voids in (i–k) are due to fibre pull-out during fracture (l) solvent-

extracted VGCNFs, showing evidence of fibre shortening to an average length of ca. 20 lm by roll milling.

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 89

were also obtained at 20–30 wt.% lower loading of the filler

[1,29].

Patton et al. [11] reported that the thermal conductivity of

VGCNF/epoxy (using Pygrograf III fibres) increases linearly

with the increase of volume fraction of fibre. This is in agree-

ment with our work. However, they reported very low thermal

conductivity values for these composites, finding that com-

posites with VGCNF loadings of 30 and 39.2 vol.% had thermal

conductivities of only 0.6587 W/m K and 0.8072 W/m K,

respectively. In contrast to their work, similar thermal con-

ductivity values were achieved in this work at much lower

loadings (6–9 vol.%) of the fibres. This is attributed to more

intense shearing of the mixtures by RM, resulting in good

dispersion of the fibres in the matrix.

In view of the simplicity and common availability of the

conventional MM technique, the present study did neverthe-

less investigate the effect of the propeller’s mixing speed

and duration on the thermal conductivity of VGCNF/RE

Fig. 5 – Thermal conductivities of the VGCNF/epoxy

composites as a function of wt.% of VGCNFs.

90 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

composites. The results for the 12 wt.% VGCNF/RE composite

produced by MM are presented in Fig. 6. It can be seen that the

thermal conductivity of the composite increases with

increasing mixing speed and time. One interpretation is that

both higher speed and longer mixing time contribute to

deagglomeration of the fibres and subsequently to their good

dispersion in the matrix. The thermal conductivity of 12 wt.%

VGCNF/RE produced by MM at 4500 rpm for 15 min is almost

the same (0.561 W/m K) as was obtained by RM (0.540 W/

m K). In addition, the thermal conductivity of the composite

produced at 4500 rpm for 60 min has 14% higher thermal con-

ductivity than the corresponding composite produced by RM.

These results clearly show that, at lower filler loading, the

conventional mixing technique is equally capable of produc-

ing composites with higher thermal conductivities, provided

that the time and speed of mixing are carefully selected.

On the other hand, the thermal conductivity (0.483 W/m K)

of the 10 wt.% VGCNF/RE composite produced by CSM is al-

most the same as that of the corresponding composite pro-

Fig. 6 – Effect of mixing speed (mixing time is 15 min.) and

time (mixing speed is 4500 rpm) on the thermal

conductivity of the 12 wt.% VGCNF/RE composite produced

by MM.

duced by MM (0.459 W/m K), as shown in Fig. 5. Thus there

is no clear benefit from sonication in this case. The effect of

sonication time on the thermal conductivity of the 10 wt.%

VGCNF/RE composite produced by CSM is shown in Fig. 7.

However, since the standard deviations in the thermal con-

ductivity data are high and overlap each other, it is not possi-

ble to draw any conclusion regarding the effect of sonication

on the thermal conductivity of the composites.

Neat glassy epoxy (GE) has 27% higher thermal conductiv-

ity than the rubbery epoxy (RE) but the thermal conductivity

of 12 wt.% VGCNF/GE produced by MM (0.51 W/m K) is 10%

lower than that of the 12 wt.% VGCNF/RE composite produced

by MM; the increase in thermal conductivity of the GE com-

posite upon loading with 12 wt.% of VGCNFs is 85% lower

than that for the RE composite. GE is a highly cross-linked

polymer and this imparts a higher thermal conductivity than

that of the lightly cross-linked RE polymer. It might therefore

be expected that, upon incorporation of VGCNFs, the resulting

GE composite should have a higher thermal conductivity than

the corresponding RE composite due to GE’s inherent higher

cross-linking density. However, this is not the case for two

reasons: (i) GE has a higher viscosity than RE which inhibits

uniform dispersion of VGCNFs in GE and (ii) air trapped in this

dispersion during mixing due to this higher viscosity results

in excessive void formation in the final GE composite. It was

also observed that the maximum possible loading of VGCNFs

in GE by MM could not exceed 12 wt.% due to the high viscos-

ity of GE. The heat generated by working the high viscosity

12 wt.% VGCNF/GE dispersion increased the temperature to

above 70 �C and if mixing was prolonged this would be hot en-

ough to initiate curing of the GE which might hinder the dis-

persion of VGCNFs. Therefore, for production of VGCNF/GE

composite at higher filler loadings steps must be taken to en-

sure that the temperature does not exceed a safe limit (e.g.

50 �C). In contrast, RE overcomes the problems associated

with GE since RE not only has a low modulus in its cured state

but its lower viscosity before curing (643 cP at a shear rate of

15 s�1) allows higher loading of the filler and poses no risk of

curing initiation due to heat of working. The final RE compos-

ite also contains fewer voids compared to GE composites.

These advantages make RE more attractive than GE for mak-

ing composites with higher loadings of fillers.

Fig. 7 – Effect of sonication time on the thermal conductivity

of the 10 wt.% VGCNF/RE composite produced by CSM.

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 91

3.2.1. Theoretical models for thermal conductivityThe thermal conductivity of VGCNF/RE composites can be

predicted by various theoretical models. However, the estima-

tion of thermal conductivity by these models depends mainly

on using the correct thermal conductivity values for the fill-

ers. Although the theoretical values of thermal conductivities

of individual VGCNFs and carbon nanotubes are very high, in

reality the values are 2–3 orders of magnitude lower due

imperfections in the structure [30]. In polymer composites, a

large fraction of the filler is dispersed in the matrix resulting

in contacts between the filler particles to form conducting

networks. The contacts between the fillers also exhibit ther-

mal contact resistance which affects the overall thermal con-

ductivity of the composites. Yu et al. [30] calculated the

thermal conductivity of carbon nanofibres (grown on silicon

wafers) from thermal contact resistance measurements to

be about 14 W/m K. The structure of these fibres is same as

that of Pyrograf-III (used in this work) with diameters of

�50 nm. Patton et al. [11] estimated the thermal conductivity

of Pyrograf-I fibres having diameter in the range of 1–100 lm

to be about 20 W/m K. Therefore, the thermal conductivity

of VGCNFs used in this work was assumed to be about

20 W/m K and this value was used for the prediction of com-

posite thermal conductivity.

The thermal conductivity of VGCNF/RE composites pre-

dicted according to various models is presented in Fig. 8.

The rule of mixtures (Eq. (3)) gave overestimated values of

thermal conductivity of VGCNF/RE composites and the paral-

lel electric circuit model (Eq. (4)) gave underestimated values

[31].

Kc ¼ Vf Kf þ VmKm ð3Þ

Kc ¼KmKf

Vf Km þ VmKfð4Þ

In Eqs. (3) and (4), Kc is the thermal conductivity of the

composite, Vf and Vm are the volume fraction of the VGCNF

and matrix, respectively and Kf and Km are the thermal

Fig. 8 – Theoretically predicted and experimentally

determined thermal conductivities of VGCNF/RE composites

vs. volume fraction of the fibres.

conductivities of the VGCNF (20 W/m K) and the matrix

(0.1769 W/m K), respectively.

The former assumes that fillers are able to form conduc-

tive networks in the matrix and ignores any effect of filler as-

pect ratio. The latter model assumes that the fillers are

homogeneously dispersed and are completely isolated in

the matrix. Consequently, phonons would be damped in the

insulating matrix and will not propagate from one particle

to another. The assumption of the parallel electric circuit

model cannot be envisaged as realistic because it is not pos-

sible to produce polymer composites without filler–filler con-

tacts at high filler loadings. Hence, both models are unable to

predict accurately the thermal transport behaviour of the

VGCNF/RE composites.

A more realistic model for 3-dimensional randomly ori-

ented short fibre reinforced composites was developed by

Hatta et al. [32] based on the Eshelby’s equivalent inclusion

model [33] (Eq. (5)) and is known as the Hatta–Taya model.

This model takes into account the aspect ratio of the fibres,

which is one of the key parameters governing the thermal

transport properties of the fibre based composites.

Kc ¼½1þ VffðKf � KmÞð2S33 þ S11Þ þ 2Kmg� � Km

Jð5Þ

where,

J ¼ 3ð1 � Vf ÞðKf � KmÞS11S33 þ Kmf3ðS11 þ S33Þ � Vf ð2S11 þ S33Þgþ3K2

mðKf�KmÞ, S11 and S33 are the tensors for thermal conduction, a

function of fibre’s aspect ratio g.

S11 ¼g

2ðg2 � 1Þ32

fgðg2 � 1Þ12 � cos h�1gg

S33 ¼ 1� 2S11

g ¼ LD, where L is the length and D is the diameter of the fibre.

It was observed by SEM that VGCNF/RE composites contain

randomly oriented VGCNFs and therefore, the thermal con-

ductivity of the composites is assumed to be isotropic at all

loadings. The average length and diameter of the fibres were

determined from SEM analysis of VGCNF/RE composites pro-

duced by RM and were about 20 lm and 116 nm (Fig. 4(l)),

respectively, i.e., the aspect ratio (g) is 172. Using the value

of g = 172 and assuming the thermal conductivity of VGCNF

is 20 W/m K, our thermal conductivity data for VGCNF/RE

composites produced by RM correlates well to the Hatta–Taya

model. It can be seen from Fig. 8 that by assuming a fibre

length, L = 1 lm (i.e., g = 9), the thermal conductivity of the

resulting composites would be significantly lower than for

composites having our observed fibre length, L = 20 lm (i.e.,

g = 172). However, it was also observed that when the value

of g was set to be >200 (e.g. g = 1030, L = 120 lm and

D = 116 nm) this had no significant effect on the thermal con-

ductivity of the resulting composites compared to the case for

g = 172. This shows that the Hatta–Taya model is valid only for

a randomly oriented short fibre composite (i.e. the present

case) when the aspect ratio is below a certain value [33].

3.3. Electrical conductivity

The electrical conductivities (measured in the direction paral-

lel to that of gravity during curing of the resin, rz) of the

92 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

VGCNF/epoxy composites produced by various techniques as

a function of wt.% of VGCNFs are presented in Fig. 9. As

observed for the thermal conductivity, the highest increase

in electrical conductivity with VGCNF loading is observed

for the composites produced by the RM and lowest increase

is observed for composites produced by CSM. The percolation

threshold (Pc), defined as the filler content required to achieve

a conductivity of P10�6 S.m�1 [31], for VGCNF/RE composites

produced by RM was found to be just less than 2 wt.%. On the

other hand, Pc may be as high as 4 wt.% for composites pro-

duced by MM since the electrical conductivity of 2 wt.%

VGCNF/RE composite produced by MM was not detectable

by the instrument used (which can detect resistance up to

100 MO).

In contrast to the thermal conductivity behaviour of the

composites produced by MM, increasing the mixing time at

4500 rpm caused no significant increase in the electrical con-

ductivity of the 12 wt.% VGCNF/RE composite produced by

MM. This was the main reason that all of the composites pro-

duced by MM used a speed of 4500 rpm (for no more than

15 min.). However, it should be noted that unlike thermal con-

ductivity, the electrical conductivity of 12 wt.% VGCNF/RE

composite produced by RM is an order of magnitude higher

than that of the corresponding composite produced by MM.

Hence, the lower Pc obtained for the composites produced

by RM and higher electrical conductivity indicates that the

RM not only dispersed the VGCNFs in the matrix well but also

distributed them uniformly. The good distribution of these fi-

bres in the matrix decreases their separation by the polymer

matrix which increases the probability of electron tunneling

and hence increases the electrical conductivity of the result-

ing composite.

Fig. 9 – Electrical conductivities of VGCNF/RE composites produce

inset shows the fit of RM-derived VGCNF/RE composites’ data to

The electrical conductivity of 10 wt.% VGCNF/RE composite

produced by CSM is one order of magnitude lower than the cor-

responding composite produced by MM. This is thought to be

due to a decrease in aspect ratio of the fibres or breaking of car-

bon bridges between the fibres during the sonication process

[18]. It was also observed for 10 wt.% VGCNF/RE composite pro-

duced by CSM that longer sonication times (up to 10 h) caused

no increase in electrical conductivity and this is similar to what

is observed for its thermal conductivity.

The electrical conductivities of the VGCNF/RE composites

were measured in different directions through samples (as

shown in Fig. 3) to study the effect of fibre orientation on

the electrical conductivity and these are plotted in Fig. 10 as

a function of wt.% of VGCNF.

The electrical conductivities of VGCNF/RE composites pro-

duced by RM are almost the same in every direction at all

loadings of VGCNFs. This clearly indicates that the compos-

ites produced by RM have no preferred orientation of the

VGCNFs, which is thus oriented randomly in the matrix.

Hence, all of the composites produced by RM are isotropic.

On the other hand, the composites produced by MM have

slightly higher electrical conductivities in the x or y direction

than in the z-direction. This indicates that fibres are slightly

more oriented in the x or y directions at VGCNF loadings of

up to 15 wt.%. These results show that the composites pro-

duced by RM are more isotropic than those produced by MM.

The electrical conductivity comparison between 12 wt.%

VGCNF/GE composite produced by MM and the corresponding

RE composite is also shown in Fig. 10. The GE composite has

an electrical conductivity which is 2 orders of magnitude

higher (in the z-direction) than for the RE composite. The

electrical conductivities of GE composites are almost the

d by various techniques as a function of wt.% of VGCNF. The

the power law.

Fig. 10 – Electrical conductivities of VGCNF/RE composites as a function of wt.% of VGCNF produced by RM and MM measured

in different directions. The electrical conductivity of VGCNF/GE composite is also presented for comparison.

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 93

same in all directions unlike RE composites, and the GE com-

posite produced by MM at 12 wt.% of VGCNFs is highly isotro-

pic in nature. The higher electrical conductivity of the VGCNF/

GE composite is attributed to the inherent high electrical con-

ductivity (6.3 · 10�9 S.cm�1) and high cross-linking density of

GE compared to RE [2]. The extensive cross-linking of GE

brings the molecular chains of epoxy close to each other

and, compared to RE, this could decrease the gap between

the fibres allowing more electrons to tunnel/transport be-

tween the fibres and through the thin layer of GE.

At the percolation threshold composites undergo a transi-

tion from insulator to conductor due to the ability of the fillers

to form conductive networks. The current–voltage behaviour

of the composites below the percolation threshold results in

a straight line according to Ohm’s law which indicates that

particle–particle contact is the dominant mechanism of con-

duction. Electrical conductivity of the composites thus in-

creases by several orders of magnitude compared to the

neat polymer. However, above Pc the current–voltage relation-

ship of the composites does not follow a linear relationship

[4]. This indicates that electron tunneling between filler parti-

cles is the dominant conduction mechanism. The behaviour

of such composites above Pc can be described with a percola-

tion-like power law (Eq. 6) [34].

r / ðP� PcÞt ð6Þ

Where, r is electrical conductivity of the composite, P is the

volume fraction of the VGCNFs, Pc is the volume fraction of

the VGCNFs at the percolation threshold and t is the critical

exponent which depends on the structure of the conductive

networks.

The inset in Fig. 9 shows that the VGCNF/RE composites’

data fit to the power law, indicating that the dominant mecha-

nism is electron tunneling at Pc. The t value for VGCNF/RE com-

posites is 2.5, higher than the universal value of 2 predicted by

statistical percolation theory [35]. Values of t between 1.7 and 2

are observed for spherical fillers while the values of t > 2 ob-

served for VGCNF [16] and carbon nanotube-based composites

[36] are attributed to their high aspect ratio [4]. The Pc value for

VGCNF/RE composites found in the present work (ca. 2 wt.%)

corresponds well to that reported by Kotaki et al. [23] but is sig-

nificantly higher than that reported by others [16,37]. Allaoui

et al. [37] reported significantly lower Pc (0.064 wt.%) and t

(1.94) values for VGCNF/epoxy composites. The fibre–fibre

and fibre–matrix contacts play a significant role in controlling

the electron tunneling behavior. The tunneling phenomenon

is influenced not only by inter-particle distance (i.e. thickness

of the polymer layer between fillers) but also by the tunneling

resistivity of the polymer [37]. The higher Pc and t values in

the present work might be due to the nature of rubbery epoxy

matrix, which has lower cross-linking density and presumably

lower electrical conductivity than the epoxy used by Allaoui

et al. [37] and Prasse et al. [16]. We have already noted that

the 12 wt.% VGCNF/GE composite has 2 orders of magnitude

higher electrical conductivity than the corresponding RE com-

posite. This suggests that RE has much lower electrical conduc-

tivity than GE and hence contributes to higher Pc and t values,

as compared with those observed by others.

The relationship between electrical and thermal conductiv-

ity as a function of volume fraction of the VGCNF for VGCNF/RE

composites produced by RM is presented in Fig. 11. It can be

seen from Fig. 11 that a significant increase in both electrical

and thermal conductivity of the composites occurs above the

percolation threshold of �1.1 vol.%. This indicates that the

conducting networks formed facilitate both electrical and ther-

mal transport in these composites at Vf above Pc. The inset in

Fig. 11 also shows that the electrical conductivity increases

exponentially with the increase of thermal conductivity (at Vf

between 0.032 and 0.26). The experimental data fit best to Eq.

(7) as shown in the inset

Fig. 11 – Thermal and electrical conductivity of VGCNF/RE composites vs. vol.% of VGCNF. The inset shows that the electrical

and thermal conductivity data of VGCNF/RE composites fit to the exponential model.

94 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

Y ¼ yo þA expðRoxÞ ð7Þ

Where, y is electrical conductivity and x is thermal con-

ductivity, yo is the intercept and A and Ro are constants. The

greater increase in electrical conductivity, as compared to

thermal conductivity, with increasing Vf is related to the in-

crease in electron tunneling between the fibres due to

decreasing gaps, and thus thinner layers of the polymer, be-

tween them.

VGCNF/RE composites showed significant enhancements

in electrical conductivity compared to neat RE but these val-

ues are still significantly lower than required for developing

electrically conductive adhesives for EPAs. However, the good

dispersion of the fibres in the matrix and the isotropic nature

of these composites makes them promising candidates for

electrostatic dissipation and electromagnetic interference

shielding applications.

3.4. Comparison of VGCNF/rubbery epoxy compositeswith carbon nanotube/epoxy composites

Significant efforts have been directed towards harnessing

the high thermal conductivity and electrical conductivity

of CNT in polymer matrix composites. Published data sug-

gest that incorporation of CNT offers significant improve-

ments in epoxy resin transport properties but there are

significant variations in the values of thermal conductivity

reported.

Previously, Thostenson and Chou [38] reported the thermal

conductivities of multiwalled CNT/epoxy composites pro-

duced by RM. They reported that the composites with 5 wt.%

CNT have thermal conductivities of �0.35 W/m K, i.e. �60%

higher than for epoxy alone. It is appropriate to compare the

thermal conductivity of VGCNF/rubbery epoxy composite, also

produced by RM, in this work and this comparison shows that

the incorporation of 5 wt.% VGCNF into rubbery epoxy can pro-

duce �96% enhancement in thermal conductivity. On the other

hand, Yu et al. [39] reported a thermal conductivity increase of

250% upon incorporation of purified single walled CNT in epoxy

and this increase is twice as high as that achieved by incorpora-

tion of VGCNF in rubbery epoxy at an equivalent loading of

6 wt.%. However, Yu et al. [39] reported almost same percolation

threshold (1 vol.%) and furthermore the electrical conductivity

of their 4 vol.% purified single walled CNT/epoxy composite

was the same as that measured in this work for the VGCNF/rub-

bery epoxy composite produced by RM at equivalent loading. In

contrast to the findings of Thostenson et al. and Yu et al., Gojny

et al. [31] reported only a 4% increase in thermal conductivity

upon incorporation of 1.2 wt.% of double walled CNT into epoxy.

Thus, comparison of VGCNF/rubbery epoxy composites with

CNT/epoxy composites suggests that, overall, VGCNFs and CNTs

can give comparable transport properties when used as fillers at

equivalent loadings in epoxy. Thus, the lower cost of VGCNF and

the ease with which they may be processed at higher loadings to

form composites affords them a major advantage over CNT

when producing composites for EPAs.

Fig. 13 – Compression stress–strain diagram of neat RE, neat

GE, 12 wt.% VGCNF/RE, 12 wt.% VGCNF/GE and 40 wt.%

VGCNF/RE composites.

C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7 95

3.5. Compression testing

The uniaxial compression stress–strain curves of neat RE and

VGCNF/RE composites produced by RM at VGCNF loadings of

12, 15, 25 and 40 wt.%s are presented in Fig. 12. It can be seen

that neat RE has very low modulus (6.77 MPa at a compressive

strain of 20%) and it experiences a compressive strain of �28%

before failure. This demonstrates the conformable nature of

rubbery epoxy which resembles that of elastomers. The addi-

tion of VGCNFs in RE at loadings of 12 and 15 wt.% increases

the compressive modulus by 2· and increases the compres-

sive strength and strain to failure (at which a crack appeared

in the sample) by �3· and �1.5·, respectively. The compres-

sive modulus and failure strength of the 25 wt.% VGCNF/RE

composite increased by �3· and �4·, respectively compared

to neat RE. At 40 wt.% loading of VGCNF, the compressive

modulus, strength and strain to failure increased by �3·, 5·and 1.8·, respectively, again compared to neat RE. Overall,

the compressive modulus increased by only 2–3· with VGCNF

loadings up to 40 wt.%, demonstrating the highly compliant

nature of the RE composites. Thus, compression testing of

VGCNF/RE composites shows that VGCNF addition signifi-

cantly improves the mechanical properties of RE without

compromising the original compliant nature of RE. The

VGCNF increases the compressive strength of RE by acting

as a barrier to crack propagation. It was observed during com-

pression testing that neat RE completely crumbles into pieces

at failure. This behaviour was not observed in the case of

VGCNF/RE composites as the samples remained intact even

when a crack appeared thereby avoiding catastrophic failure.

For comparison, the compressive stress–strain diagram for

VGCNF/RE and VGCNF/GE composites produced by MM is

shown in Fig. 13. The compression stress–strain curves of neat

GE and VGCNF/GE composites show that these materials have

very high moduli compared to those of RE composites. It can be

seen from Fig. 13 that the VGCNF/GE composite has �2· lower

compressive strength at failure compared to neat GE. In con-

trast to RE composites, VGCNFs do not act as a reinforcing filler

for GE composites. The VGCNF/GE composites inherit their

high stiffness from GE and further addition of VGCNF was det-

Fig. 12 – Compression stress–strain diagram of neat RE and

VGCNF/RE composites produced by RM.

rimental to the compressive properties of GE, with the fibres

effectively acting as flaws in this case. This behaviour is the

opposite to what is observed for VGCNF/RE composites. The

compressive modulus of VGCNF/RE composites (measured at

a strain of 20%), even at a loading of 40 wt.%, is�12· lower than

for the 12 wt.% VGCNF/GE composite (Fig. 13).

The compressive stress–strain curves for 12 wt.% VGCNF/

RE composites produced by RM and MM are similar showing

no significant difference at macro-scale (Fig. 13). VGCNF/RE

composites have very low modulus and high compressive

strain to failure even at loadings of up to 40 wt.%. These re-

sults suggest that the effect of VGCNF addition on the

mechanical properties of the material mainly depends upon

the original nature of the polymer. Hence, VGCNFs are highly

effective fillers for improving the mechanical properties of RE

without sacrificing its compliant behaviour. It is believed that

the mechanical properties of VGCNF/RE can be further

enhanced by compression moulding of these composites as

this can be expected to further reduce their void content.

3.6. Hardness testing

The shore hardnesses of neat RE and VGCNF/RE composites

produced by RM and MM are presented in Fig. 14. In parallel

Fig. 14 – Shore hardnesses of neat RE and VGCNF/RE

composites produced by RM and MM.

96 C A R B O N 5 0 ( 2 0 1 2 ) 8 4 – 9 7

to the compression behaviour, the hardness of VGCNF/RE

composites increases with the increasing wt.% of VGCNF.

The hardness of VGCNF/RE composites produced by RM at

25 and 40 wt.% loading of VGCNF increases by 22% and 28%,

respectively, compared to neat RE. The hardness of 12 wt.%

VGCNF/RE produced by MM is �10% higher than the corre-

sponding composite produced by RM and this might be due

to the presence of agglomerated fibres. The small increases

in hardness and compressive modulus upon incorporation

of VGCNF into RE clearly show that VGCNF/RE composites

are highly compliant materials.

4. Conclusions

• Variations of VGCNF concentration and processing within

the rubbery epoxy matrix have direct and appreciable

effects upon the transport and mechanical properties of

the resulting composites and this allows for the potential

of VGCNFs to be more fully exploited than previously dem-

onstrated in polymer composites.

• Composites produced by RM have much better dispersion and

distribution of VGCNFs in RE than those produced by either

MM or CSM methods. This is attributed to good shear mixing

and crushing action in the RM. The maximum possible load-

ing of VGCNFs in RE depends on the mixing technique. The

VGCNFs can be loaded conveniently into RE up to 40 wt.% by

RM and up to 12–15 wt.% by MM. SEM analysis and electrical

conductivity data show that the VGCNF/RE composites pro-

duced by RM are isotropic in nature while some anisotropy

is observed in the composites produced by MM.

• Increasing the speed and time of mixing in the MM process

significantly improves the thermal conductivities of the

resulting composites. However, increasing the mixing time

does not have an appreciable effect on the electrical con-

ductivities of the resulting composites. On the other hand,

increasing sonication time during CSM mixing deteriorates

the transport properties of the resulting composites. This is

thought to be due to a decrease in the aspect ratio of the

fibres and/or breaking of carbon bridges between the fibres.

• The composites produced by RM have 13–38% higher ther-

mal conductivities and an order of magnitude higher elec-

trical conductivities than the composites produced by MM.

However, MM is capable of producing composites with

thermal conductivities comparable to those of composites

produced by RM (at low loadings of VGCNFs) by increasing

the speed and time of mixing.

• The thermal conductivities of VGCNF/RE composites increase

linearly with increasing wt.% of VGCNF. The thermal conduc-

tivity of VGCNF/RE composites produced by RM with 15 wt.%

and 40 wt.% of VGCNF resulted in a�4-fold increase (0.811 W/

m K) and �10-fold increase (1.854 W/m K), respectively, com-

pared to RE alone (0.1769 W/m K). The thermal conductivity

data for VGCNF/RE composites (produced by RM) best fit to

the Hatta–Taya model, when assuming the thermal conduc-

tivity of VGCNF is about 20 W/m K.

• The electrical conductivity of VGCNF/RE composites also

increases with increasing wt.% of VGCNF. The lowest per-

colation threshold (2 wt.%) is obtained for composites pro-

duced by RM rather than for those produced by MM. It is

observed that above Pc, the electrical behaviour of VGCNF/RE

composites follows the percolation-like power law. A correla-

tion between electrical and thermal conductivity with

increasing wt.% of VGCNF is observed above Pc. It is also

observed that the electrical conductivity of VGCNF/RE com-

posites increases exponentially with increasing thermal con-

ductivity (i.e. with increasing wt.% of VGCNF) due to a stronger

electron tunneling effect at higher loading of VGCNF.

• The electrical conductivity of a VGCNF/RE composite was

found to be 2 orders of magnitude lower than the corre-

sponding GE composite produced by MM. This behaviour

is attributed to the inherent high electrical conductivity

and high cross-linking density of GE.

• Comparison of these VGCNF/rubbery epoxy composites

with carbon nanotube/epoxy composites indicates that

VGCNF can be used to produce composites with thermal

and electrical transport performance comparable to that

of corresponding composites based on carbon nanotubes.

• Compression testing of VGCNF/RE composites showed that

VGCNF incorporation up to 40 wt.% significantly increased

the compressive strength and strain to failure without sig-

nificant detriment to the original compliant nature of RE.

On the other hand, it was observed that VGCNF incorpora-

tion into GE was detrimental to its mechanical properties.

A VGCNF/RE composite with 40 wt.% loading of VGCNF has

a compressive modulus 12 times lower than that of the

12 wt.% VGCNF/GE composite. This shows that VGCNF/RE

composites are much more compliant materials than

VGCNF/GE composites. The shore hardness of the VGCNF/

RE composite at a loading of 40 wt.% increased by only

28% compared to that of neat RE.

• VGCNF/RE composites, with their unusual combination of

high thermal and electrical conductivities and conformable

nature (for gap filling applications) are promising candidates

for electronics packaging applications, particularly for ther-

mal interface applications. VGCNF/RE composites offer high

compliance, ease of processing (even at high filler loading),

low viscosity and long shelf-life before curing, and fewer voids

(which could act as stress concentrators) after curing. These

features make them attractive as adhesives for various EPAs.

In future work, this concept can be developed by testing the

thermal management performance of these composites in

‘‘demonstrator’’ electronic packaging assemblies which

include relevant working electronic components.

Acknowledgements

The authors thank Morgan AM&T and EPSRC for funding

M.A.R.’s Dorothy Hodgkin Postgraduate Award Scholarship.

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