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

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Observation of mechanical percolation in functionalizedgraphene oxide/elastomer composites

Khalid Nawaz a, Umar Khan b, Noaman Ul-Haq a, Peter May b, Arlene O’Neill b,Jonathan N. Coleman b,*

a Department of Chemical Engineering, School of Chemical and Materials Engineering (SCME),

National University of Sciences and Technology, H-12 Islamabad, Pakistanb School of Physics and CRANN, Trinity College University of Dublin, Ireland

A R T I C L E I N F O

Article history:

Received 30 January 2012

Accepted 11 May 2012

Available online 19 May 2012

0008-6223/$ - see front matter � 2012 Elsevihttp://dx.doi.org/10.1016/j.carbon.2012.05.029

* Corresponding author: Fax: +353 1 6711759.E-mail address: colemaj@tcd.ie (J.N. Colem

A B S T R A C T

We have covalently functionalized graphene oxide (GO) with octadecylamine (ODA) to form

GO–ODA. This material can be dispersed in tetrahydrofuran (THF) and subsequently

formed into composites with polymers such as thermoplastic polyurethane (TPU). We have

characterized the mechanical properties of composites of GO–ODA in TPU. No increase in

stiffness was observed at loading levels below 2.5 vol%. Reinforcement appeared to start

sharply at this volume fraction and subsequently increased as a power law with increasing

volume fraction. This behavior is typical of percolation and shows that the low-strain

stress is not increased until the functionalized graphene flakes form a percolating network.

Slightly different behavior is observed for properties related to material failure. The ulti-

mate tensile strength increased linearly with graphene content up to the percolation

threshold before subsequently falling off. Similarly the strain at break was constant below

the percolation threshold but fell off dramatically above it. This work shows the impor-

tance of network formation in the reinforcement of elastomeric materials.

� 2012 Elsevier Ltd. All rights reserved.

1. Introduction

Graphene oxide (GO) [1] is an exceptional material which has

shown great promise in a number of application areas [2–5].

One area where GO is expected to make a big impact is as a

filler in composites. Incorporation of GO into polymers such

as polyvinyl alcohol can result in significant increases in

mechanical properties [6], while composite formation with

reduced GO can give large increases in electrical properties

[1]. However, due to its polar nature, as-produced GO is not

a natural filler material for non-polar polymers. This can be

addressed by the covalent functionalization of GO with non-

polar polymers or molecules. These groups alter the surface

chemistry thus compatiblising the functionalized GO with

both non-polar solvents and polymer matrices.

er Ltd. All rights reservedan).

Of particular interest is the functionalization of graphene

oxide with octadecylamine (ODA). The reason for this is that

the amine group of the ODA should react readily with epoxide

or carboxylic acid groups in the GO [7–9]. Critically, this reac-

tion can be followed with FTIR [10], allowing researchers to be

confident that they have produced GO–ODA. Because of the

presence of covalently attached non-polar groups, GO–ODA

should disperse well in non-polar polymers. This makes it a

good model system for the exploration of the properties of

polymer–graphene composites processed using organic

solvents.

One interesting class of composite consists of thermoplas-

tic elastomers filled with nano-fillers such as carbon nano-

tubes or graphene [11–14]. These are of great interest as

addition of the nano-filler can result in the scaling of strain

.

C-O

GO-ODA

ODA

GO

C=O

C-OH

AlkylC-H

N-H

Amide C=O

(alkoxy)

C-O (alkoxy)

C-O (carboxy)

Tra

nsm

issi

on

4000 3000 2000 1000

Wavenumber (cm-1)

Fig. 1 – FTIR spectra of GO (top), ODA (middle) and GO–ODA

(bottom).

4490 C A R B O N 5 0 ( 2 0 1 2 ) 4 4 8 9 – 4 4 9 4

at break and low-strain stress all the way from those of an

elastomer to a rigid thermoplastic [15,16]. However, the rein-

forcement mechanism at work in these systems is unclear.

We suggest that because of its dispersibility and ease of com-

posite formation, GO–ODA is an ideal model system to inves-

tigate the nature of this mechanism. In this work, we

functionalize GO with ODA. FTIR shows the attachment to

be covalent while SEM analysis shows the GO–ODA to be well

dispersed, even at very high loading levels. Mechanical mea-

surements show no reinforcement at volume fractions below

2.5%. Above this filler content, both stiffness and low-strain

stress increase significantly in line with percolation theory.

However, we find that once a percolating network is first

formed, the ultimate tensile strength and strain at break be-

gin to fall with increasing graphene content.

2. Experimental

Graphite powder (Sigma Aldrich) was oxidized using a modi-

fied version of the Hummers method [17]. In brief, graphite

powder was oxidized using NaNO3, H2SO4 and KMNO4 in an

ice bath [17]. The material obtained was centrifuged at

2000 rpm for 10 min. The supernatant, containing inorganic

salts, was decanted and the deposit washed several times

with deionized water until the pH was neutral. This sediment

was dried overnight in a vacuum oven at 80 �C.

The dried GO was treated with ODA as previously de-

scribed [8] with slight modification. Briefly, 300 mg of GO

was dispersed in 30 ml of water while 300 mg of ODA was dis-

solved in 30 ml of dimethyl formamide (DMF). Both the dis-

persions were mixed together in round bottom flask and

refluxed for 48 h at 100 �C under continuous stirring. The

resultant material was filtered through a nylon membrane

of pore size 0.45 lm (Sterlitech). The filtered material was

thoroughly washed with THF to remove any unreacted, free

ODA.

In next step the ODA functionalized graphene (GO�ODA)

was dispersed in THF by sonication for 30 min in a sonic bath

(Branson 1510-MT). The dispersion was centrifuged at

500 rpm for 45 min using a Hettich Mikro 22R to remove large

aggregates or unfunctionalized GO. The supernatant was col-

lected and filtered through a nylon membrane of pore size

0.45 lm (Sterlitech). The membrane supported GO–ODA was

dried in a vacuum oven overnight at 60 �C to facilitate accu-

rate mass determination. The thermoplastic polyurethane

(TPU) used in this work was Morthane which was obtained

from Huntsman Polyurethanes (Morthane PS455-203 - an aro-

matic polyester based thermoplastic polyurethane). This

polymer (6 g) was dissolved in 100 ml of THF (60 mg/ml) by

stirring at 40 �C for 24 h. GO–ODA was added to the TPU solu-

tion at the required concentration and sonicated for 30 min.

A range of ODA/GO/polymer dispersions were prepared by

adding ODA–GO to the polymer solutions to give various mass

fractions of ODA–GO from 1 to 50 wt.% (but keeping the poly-

mer concentration at 60 mg/ml). These dispersions were son-

icated for four hours in a sonic bath with no significant

dependence of viscosity on ODA–GO content. The dispersions

were stable over the timeframe of the sample preparation.

However, if stored for days, the higher concentration ODA/

GO/polymer dispersions showed visual aggregation of the

ODA–GO. Films were formed by casting the dispersions into

in Teflon trays of dimensions 4 · 4 · 2 cm. These cast samples

were dried at room temperature and 900 mbar in a vacuum

oven for 24 h. The dried films were then placed in an oven

at 60 �C for 48 h to remove trace THF. All samples were of con-

stant mass (�150 mg). The film thicknesses were in the range

of 50–60 microns. A reference sample of TPU was also

prepared.

Mechanical testing was performed using a Zwick Roell ten-

sile tester with a 100 N load cell at a strain rate of 50 mm/min.

In all cases, the Young’s modulus was measured in the linear

region between 0% and 2% strain. Scanning electron micros-

copy (SEM) was performed using a Carl Zeiss Ultra Plus Field

Emission Scanning Electron Microscope. FTIR was measured

on crushed powder on a glass slide in transmittance mode

using Nexus Nicolet FTIR.

3. Results and discussion

Initially, it is important to ascertain whether the ODA has

been covalently attached to the GO. We expect the reaction

to be of the form [8]:

R� CO� OHþ R�NH2 ! RCO�NH� Rþ H2O

where the nitrogen of the ODA bonds to the carbonyl carbon

of the GO, although reaction with the basal plane epoxides

has also been suggested [7–9]. To test this hypothesis, we per-

formed FTIR spectroscopy on the GO, neat ODA and GO–ODA

samples (Fig. 1). The GO exhibits the characteristic bands at

1720, 1630, 1390, 1220 and 1050 cm�1, which can be associated

with C=O, C=C, C–O (carboxy), COH and C–O (alkoxy) groups,

and a broad absorption band between 3000 and 3500 cm�1

associated with the hydroxyl groups [18–21]. The FTIR spec-

trum of the neat ODA shows a peak characteristic of alkyl

groups at 2850 cm�1 and a peak at 2918 cm�1 due to asym-

metric C–H stretching. The peaks around 1500 cm�1 are due

to methylene scissoring deflection and other effects associ-

ated with the alkyl chain [22]. Most importantly, we observe

a band at 3333 cm�1 associated with the amine group (N–H)

0 200 400 600 800 1000

0

10

20

30

40

10%

50%Stre

ss, σ

(MPa

)

Strain, ε (%)

TPU

1%

4%

40%

0 2 4 6 8 1002468

10

Fig. 3 – Representative stress strain curves for GO–ODA/TPU

composites with weight fractions between 0% and 50%.

Inset: the low strain regime.

0.00.10.20.30.40.5

E (G

Pa) A

02468

10σ ε=

3% (M

Pa) B

01020304050

σ B (MPa

)

C

0 10 20 30 40 50

0

400

800

1200

ε B (%)

Graphene content (wt%)

D

C A R B O N 5 0 ( 2 0 1 2 ) 4 4 8 9 – 4 4 9 4 4491

[4,17]. In the GO–ODA FTIR spectra, the disappearance of this

primary amine group and the emergence of a new peak at

3425 cm�1 indicate the formation of –C–NH–C– units due to

the reaction between the epoxide group and the amine group.

This suggests that ODA not only reacts with carbonyl carbons

but also with epoxides of the GO to form secondary amines

[8,23]. Further evidence of functionalization comes from the

fact that carbonyl C=O band at 1720 cm�1 in the GO sample

is obscured in the GO–ODA spectra and a new band has ap-

peared at around 1640 cm�1. This is consistent with amide

C=O stretching as would be expected from the reaction above

[18]. Thus, we believe that FTIR spectra clearly confirm that

the ODA molecules were attached to the graphene nano-

sheets through chemical modification.

Unlike GO, the GO–ODA was readily dispersible in THF.

This allows the formation of composites by blending disper-

sions of GO–ODA with solutions of TPU in THF. These GO–

ODA/polymer dispersions can then be formed into composite

films. In all cases, including 50 wt.% GO–ODA, composites

were very uniform to the naked eye with no appearance of

aggregates. Fig. 2 shows an SEM image of the fracture surface

of a 40 wt.% GO–ODA composite. The inset shows a high mag-

nification image showing a homogeneous dispersion of GO–

ODA within the polymer.

We performed mechanical characterization of GO–ODA/

TPU composites for a range of GO–ODA mass fractions. Repre-

sentative stress strain curves for a subset of the composites

are shown in Fig. 3. It is clear from this graph that the pres-

ence of GO–ODA has a significant effect on the polymer. For

low mass fractions the stress at all strains appears to increase

while for higher mass fractions, both stress and strain at

break fall dramatically. From the stress strain curves, one

can extract four main mechanical properties; the Young’s

modulus or stiffness, E, the stress at low (3%) strain, re¼3%,

the ultimate tensile strength, rB, and the strain at break, eB.

The average values of these quantities are shown as a func-

tion of the filler mass fraction in Fig. 4. The Young’s modulus

appeared to increase almost linearly with increasing GO–ODA

mass fraction from 9.6 MPa for the polymer to 335 MPa for the

50 wt.% composite. Similarly, the stress at 3% strain increased

in an almost linear fashion from 0.3 MPa for the polymer to

�10 MPa for the 50 wt% composite. In contrast, the ultimate

tensile strength initially increased from 27 MPa for the poly-

mer to 38 MPa for the 3 wt.% composite before falling steadily,

Fig. 2 – SEM images of the fracture surface of a 40 wt.% GO–

ODA/TPU composite film.

Fig. 4 – The effect of GO–ODA on mechanical properties of

composites for a range of loading levels. (A) Young’s

modulus, (B) stress at 3% strain, (C) ultimate tensile

strength, (D) strain at break.

reaching 10 MPa for the 50 wt% composite. Interestingly, the

initial increase was linear with a slope of 330 MPa. Although

a linear increase is predicted by the rule of mixtures, this

slope is low compared with typical values of �GPa usually

found for composites of thermoplastic polymers filled with

nanotubes or graphene [24,25]. The strain at break decreased

steadily with increasing GO–ODA mass fraction from �1000%

for the polymer to �10% for the 50 wt.% sample.

Broadly speaking, these results are typical of what is ob-

served for elastomers filled with high volume fractions of

0.01

0.1

E (G

Pa) A

1

10

σ ε=3%

(MPa

) B

01020304050

σ B (MPa

)

C

0.01 0.10

200400600800

1000

ε B (%)

Graphene volume fraction, φ

D0.6

Fig. 5 – The same data as in Fig. 4 but plotted as a function of

GO–ODA volume fraction on logarithmic plots. The lines in

(A) and (B) illustrate percolation behavior while the line in

(C) illustrates linearity. The vertical arrows illustrate the

percolation threshold while the horizontal arrows show

that value of each property displayed by the polymer.

4492 C A R B O N 5 0 ( 2 0 1 2 ) 4 4 8 9 – 4 4 9 4

nanotubes [11,13,16] or graphene [15]. These results appear

similar to those of Khan et al. who mixed pristine solvent-

exfoliated graphene with thermoplastic polyurethane [15].

The starting polyurethane had very similar properties while

their 50 wt% sample had values of E � 1 GPa, re¼3% � 10 MPa,

rB � 30 MPa and eB � 10%, values almost identical to those

found here. This implies that the ODA functional groups (or

indeed the oxides) play no significant role in the reinforce-

ment but ensure good dispersion.

We can also compare these results with other work on

composites of thermoplastic elastomers filled with other

nanomaterials. Liff et al. [26] increased the modulus of poly-

urethane by a factor of 20 by addition of 20 wt% nanoclay.

By the inclusion of functionalized nanotubes in TPU, Koerner

et. al. [27] reported a thirteen-fold increase in modulus and a

significant enhancement of the strength. Sahoo et al. [28] fab-

ricated functionalized SWNT/PU composites and achieved an

increment of 3 times in modulus at 20% loading. While seven

fold increase in modulus was reported by Cheng et al. in PU

composites reinforced by functionalized MWCNTs [29] the

strain at break of the composites was reduced significantly.

It is worth looking at the dependence of mechanical prop-

erties on graphene content in more detail. To do this we plot

the same data as in Fig. 4A and B but in double logarithmic

format. In addition, we have transformed the mass fraction

to volume fraction assuming a GO–ODA density of 1800 kg/

m3 and a TPU density of 1100 kg/m3. This data is shown in

Fig. 5A and B and immediately illustrates some behavior not

apparent from the linear graphs. It is clear from Fig. 4A and

B that the modulus and stress at low strain hardly increase

at all for graphene contents below �2.5 vol%. However, above

this threshold, both properties increase with graphene con-

tent as a power law. Such non-linear increases in mechanical

properties of elastomeric composites are generally attributed

to either the effects of strong matrix-filler interaction or the

formation of a filler network [30]. Here the fact that essen-

tially no increase in modulus (or re¼3%) are observed below

2.5 vol.% suggests the latter mechanism to be at work. Indeed,

this lack of reinforcement at low volume fraction implies that

neither polymer-filler stress transfer nor hydrodynamic ef-

fects are important [30]. The increase in E and re¼3% above

2.5 vol.% implies that network formation begins at this criti-

cal volume fraction. Network formation and growth can be

described by percolation theory [31]. Where the formation of

networks of the filler controls the mechanical properties of

composites, it has been shown that certain mechanical prop-

erties are described by a percolation-like scaling law [21,26]:

A ¼ AP þ A0½/� /c�t

Here A can represent either E or re¼3%, / is the filler volume

fraction, /c is the percolation threshold, i.e. the critical filler

volume fraction when a network first forms and AP is the va-

lue of the relevant mechanical property at the percolation

threshold. A0 and t are constants. As can be seen from

Fig. 5A and B, both E and re¼3% are described very well by this

expression. In both cases, the percolation threshold was

/c = 2.5 vol.% while the percolation exponent was t = 0.8.

While the percolation threshold is similar to that found in

the literature for high aspect ratio fillers, the exponent is very

low. This can be seen by comparison to data on percolation

behavior of the elastic modulus for polymer/nanoclay com-

posites as reported by Liff et al. and Ramorino et al. These pa-

pers reported percolation thresholds of 1.4 vol% and 6 wt%

and exponents of 1.5 and 2 [26,30]. In addition, studies on

the storage modulus of singlewall nanotubes in epoxy elasto-

mers showed a percolation threshold of 0.41 wt% and an

exponent of 1.65 [32]. We note that these percolation expo-

nents are typical of what is generally found during studies

on electrical percolation [33]. However, we note that one study

on composites of single walled nanotubes in polymethyl-

methacrylate, showed percolation scaling with an exponent

of 0.7 (percolation threshold 0.12 wt%). This result is very sim-

ilar to the value found here. However, it remains unclear

which properties of the network result in such a low

exponent.

These results imply that reinforcement of TPU is not due

to standard mechanisms which rely on stress transfer at the

polymer-filler interface or the effect of the filler particles on

polymer flow under stress. Rather, the critical factor is the for-

mation of a network which mechanically stiffens the mate-

rial. We interpret this by assuming that once the network

forms, the GO–ODA flakes form a jammed system. It has been

shown that for high concentration carbon black dispersions,

such a system can be described by a shear modulus which

obeys a percolation scaling law [34]. We believe that in the

C A R B O N 5 0 ( 2 0 1 2 ) 4 4 8 9 – 4 4 9 4 4493

case of a solid state composite, this behavior manifests itself

as the tensile modulus (and low-strain tensile stress) obeying

a percolation scaling law.

The formation of such a network is likely to impact on

other mechanical properties. We note that E and re¼3% are

low strain properties. In Fig. 5C and D we consider high strain

properties i.e. those associated with fracture; rB and eB. As de-

scribed above, the ultimate tensile strength, rB, increases lin-

early (solid line) at low strain. However, it is clear from Fig. 5C

that the point where linearity ceases and rB begins to fall off

coincides with the percolation threshold. In addition, as dem-

onstrated in Fig. 5D, the strain at break is reasonably constant

up to the percolation threshold, after which it falls dramati-

cally with GO–ODA volume fraction. It is not clear why this

should be the case. However, it may be that polymer chains

in the vicinity of the graphene surface tend to have reduced

mobility. This may reduce the scope of large strain deforma-

tion resulting in premature failure.

With this in mind, it is worth considering the nature of the

interfacial region. Previously, it has been shown that ODA

functionalized SWNTs selectively interact with the hard seg-

ments of thermoplastic elastomers as indicated by differen-

tial scanning calorimetry and other techniques [13].

Interestingly, these composites showed very similar mechan-

ical properties to those studied here; an apparent linear in-

crease in modulus and a peak in strength close to 2%

volume fraction. Also they displayed a slow decrease in strain

at break as ODA-SWNTs were added, reaching �600% for a

nanotube content of 10%, very similar to what was observed

here. Thus it is likely that the ODA functionalized graphene

sheets interact predominately with the PU hard segments.

Anchoring of these hard segments at the GO–ODA surface

may immobilize adjacent soft segments, limiting polymer

flow and ultimately causing premature failure.

4. Conclusion

We have used graphene oxide covalently functionalized with

octadecylamine as a model filler in polyurethane based com-

posites. We find no appreciable increases in either stiffness or

low-strain stress for loading levels below 2.5 vol%. However

above this threshold, both mechanical quantities increase

as a power law. This behavior is consistent with mechanical

percolation. This implies that the graphene oxide platelets

are effectively isolated at low volume fractions but begin to

form a network at a volume fraction of 2.5 vol%. This loading

level can be thought of as a percolation threshold. As the

loading level is increased, the network becomes more exten-

sive and the stiffness and low-strain stress increase as de-

scribed by the percolation scaling law. Interesting the

ultimate tensile strength initially increases but reaches a

maximum at the percolation threshold. Similarly the strain

at break is invariant with graphene content up to the percola-

tion threshold, after which it falls steadily. This work shows

that the mechanical properties of elastomers reinforced with

graphene can depend on parameters other than interfacial

stress transfer. For example, the formation of a network of fil-

ler particles which acts like a jammed system can dominate

the mechanical properties of the system.

Acknowledgements

We acknowledge financial support from Science Foundation

Ireland through the Principle Investigator scheme (grant

number 07/IN.7/I1772).

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