CARBON
4 8 ( 2 0 1 0 ) 1 0 1 2 –1 0 2 4
available at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/carbon
Changes in electrical and microstructural properties
of microcrystalline cellulose as function
of carbonization temperature
Yo-Rhin Rhim a, Dajie Zhang b, D. Howard Fairbrother c, Kevin A. Wepasnick c,
Kenneth J. Livi d, Robert J. Bodnar e, Dennis C. Nagle b,*
a
Applied Physics Laboratory, Johns Hopkins University, 11100 Johns Hopkins Road, Laurel, MD 20723, USA
Advanced Technology Laboratory, Johns Hopkins University, 810 Wyman Park Drive, Baltimore, MD 21211, USA
c
Department of Chemistry, Johns Hopkins University, 3400 North Charles Street, Baltimore, MD 21218, USA
d
Department of Earth and Planetary Sciences, Johns Hopkins University, 3400 North Charles Street, Baltimore, MD 21218, USA
e
Department of Geosciences, Virginia Tech, 4044 Derring Hall, Blacksburg, VA 24061, USA
b
A R T I C L E I N F O
A B S T R A C T
Article history:
AC and DC electrical measurements were made to better understand the thermal conver-
Received 9 February 2009
sion of microcrystalline cellulose to carbon. This study identifies five regions of electrical
Accepted 8 November 2009
conductivity that can be directly correlated to the chemical decomposition and microstruc-
Available online 13 November 2009
tural evolution of cellulose during carbonization. In Region I (250–350 C), a decrease in
overall AC conductivity occurs due to the loss of the polar oxygen-containing functional
groups from cellulose molecules. In Region II (400–500 C), the AC conductivity starts to
increase with heat treatment temperature due to the formation and growth of conducting
carbon clusters. In Region III (550–600 C), a further increase of AC conductivity with
increasing heat treatment temperature is observed. In addition, the AC conductivity demonstrates a non-linear frequency dependency due to electron hopping, interfacial polarization, and onset of a percolation threshold. In Region IV (610–1000 C), a frequency
independent conductivity (DC conductivity) is observed and continues to increase with
heat treatment due to the growth and further percolation of carbon clusters. Finally in
Region V (1200–2000 C), the DC conductivity reaches a plateau with increasing heat treatment temperature as the system reaches a fully percolated state.
2009 Elsevier Ltd. All rights reserved.
1.
Introduction
Electrical properties of various forms of carbon, including
graphite, anthracite carbons, carbon nanotubes, and graphene materials are of great interest in many technical areas.
The electrical properties of carbon materials derived from organic precursors have been studied extensively over the years
and have been shown to vary widely depending on the nature
of the precursor and the heat treatment temperature (HTT)
[1–8]. Microcrystalline cellulose is a basic component of wood
that has been highly refined to remove all the inorganic ash.
By using this high purity precursor material, more definitive
observations could be made on the carbonization mechanisms of organic compounds as they are converted to nongraphitizing hard carbons. Since cellulose molecules do not
contain any aromatic structures and does not exhibit an
intermediate mesophase, they yield highly amorphous
carbon even when heated to extremely high HTTs. This study
* Corresponding author: Fax: +1 410 516 7249.
E-mail address: [email protected] (D.C. Nagle).
0008-6223/$ - see front matter 2009 Elsevier Ltd. All rights reserved.
doi:10.1016/j.carbon.2009.11.020
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focuses on measuring changes in AC and DC electrical conductivities of microcrystalline cellulose during its transformation into carbon over the temperature range from 250 to
2000 C. Changes in electrical conductivity changes were also
correlated with and rationalized by the structural evolution of
microcrystalline cellulose carbon.
Earlier studies of electronic properties of soft carbons described by Pinnick [1] and Seldin [9] have shown that electrical
resistivity decreased by nine orders of magnitude as HTT increased from 600 to 3000 C, resulting in a non-metal to metal
transition. In an attempt to describe such observations, a
band model was proposed [10] in which resistivity varied as
a function of an energy band gap. As the HTT was increased
to 1200 C, the resistivities of soft carbons were shown to decrease as electrons moved to the conduction band. Later work
(Giuntini et al. [2], Guintini and Zanchetta [10]) described an
electron hopping between localized states as the main mechanism for electric conduction for samples heat treated to
600 C and between localized states near the Fermi level for
samples heat treated to 650 C. Such studies were attempted
to better describe the non-metal–metal transition observed
between HTT of 600 and 650 C for anthracene carbons. By
calculating the number of density of states as a function of
HTT, the authors have shown that the number of localized
states increases as HTT is increased within the range of
600–700 C, thus explaining the non-metal to metal transition
observed.
More recent studies by Kercher and Nagle [4] showed a
change of DC electrical conductivity of carbonized mediumdensity fiberboard heat treated to temperatures of 600 C
and above. A percolation model based on the formation and
growth of highly conducting turbostratic sites embedded
within a dielectric was used to explain the varying conductivities [11]. AC studies by Sugimoto and Norimoto [5,12] of carbonized wood materials heat treated to temperatures lower
than 600 C showed frequency dependent conductivity values. Dipole polarization and Maxwell–Wagner interfacial
polarization were used to explain such results. Furthermore,
a heterogeneous structure exhibiting interfacial polarization
that seemed to occur at the interface of isolated conductive
regions was described to support both percolation and interfacial polarization by Kercher and Nagle [11].
Two models of two-phase composites consisting of conductive regions embedded within an insulating matrix have
been proposed by McLachlan et al. [13–18]. Systems composed
of conductive regions separated interstitially by an insulating
medium with separation distances large enough to limit electron hopping and tunneling are best described by the Maxwell–Wagner model:
rm rc
ð1 /Þðri rc Þ
¼
;
rm þ 2rc
ri þ 2rc
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ð1Þ
where rm, rc, and ri denote the mean conductivity of the system, the conductivity of conductive clusters, and the conductivity of the insulating matrix, respectively, and u denotes the
volume fraction of the conductive phase.
Electrical properties of binary composites have also been
well described by the effective media and percolation theories
[13,15,18]. In such systems the conductivity is shown to increase sharply when the volume fraction of the conductive
phase reaches a critical volume fraction, uc. At this point
the percolation threshold is reached and the conductivity of
the composite changes from that of an insulator to conductor.
Systems with the conducting phase surrounded by a mixture
with an effective conductivity of the medium are well characterized by the general effective media (GEM) equation:
1=s
1=t
ð1 /Þ ri1=s rm
/ rc1=t rm
þ
¼ 0;
ð2Þ
1=s
1=t
ri1=s þ Arm
rc1=t þ Arm
cÞ
, s and t are exponents that dependent on the
where A ¼ ð1/
/c
dimensionality and geometry of the system, and the conductivity of the composite, rm, is characterized by finite values of
conductivity for the conductive phase and insulating phase,
rc, and ri, respectively. The GEM equation above yields two
limits below and above the critical volume fraction:
s
/c
for / < /c ;
ð3:aÞ
rm ¼ ri
/c /
t
/ /c
for / > /c :
ð3:bÞ
rm ¼ rc
1 /c
Above the critical volume fraction, uc, Eq. (2) reduces to Eq.
(3.b) and has the same form as the percolation equation:
rm ¼ rc ð/ /c Þt
for / > /c :
ð4Þ
A two-phase system of conductive regions in contact within an insulating matrix, where an increasing volume fraction,
u, of the conductive phase facilitates hopping and tunneling
of electrons, may best described by the percolation model.
Percolation of many two-phase composites varying in conductive and insulating volume fractions has been extensively
studied [13–21]. The reported AC conductivities followed the
general relationship:
r ¼ r0 þ rðxÞ;
ð5Þ
where the total AC conductivity observed, r, displayed a frequency independent term at low frequencies (x ! 0 Hz) or
DC conductivity, rdc, and a frequency dependent term, r(x),
whose real term followed a power law relationship of Axs
[20,22]:
r ¼ rdc þ Axs :
ð6Þ
For low fractions of conductive material below the percolation threshold, the system behaved as an insulator; and for
high fractions above the percolation threshold, the system behaved as a conductor. The conductivity of the overall system
was observed to increase dramatically as the fraction of conductive material reached the percolation threshold. Percolation of conductive sites was shown to assist electron
hopping and tunneling [7,21,23–25], therefore increasing the
overall conductivity of the system.
Although the percolation model successfully describes
conductivity changes of two-phase composites, it may not
accurately describe the varying conductivities of complex
microstructured carbon materials. Previous electrical studies
of heat treated carbon materials by Hernandez et al. [26]
and Emmerich et al. [27] incorporate the effects of porosity
and apply a modified form of the percolation model. These
studies assume a granular structure of conductive and nonconductive phases. In such system, the conductive phase is
composed of ordered microcrystallite structures, and the
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non-conductive phase includes void spaces, organic material,
and ash [27]. The volume fraction of the conductive phase is
increased with heat treatment, until finally percolation of
the conductive phase aids electron tunneling between microcrystallites [26], resulting in dramatic increases in overall conductivity. Thus, changes in porosity that occur during this
transformation must be taken into account when applying
the percolation model.
This study was undertaken to fully characterize the electrical property changes that take place during the carbonization process involving both compositional and structural
changes. By using a high purity homogeneous carbon precursor such as microcrystalline cellulose, more definitive observations can be made regarding the organic matter to carbon
conversion process.
2.
Experimental
2.1.
Sample preparation
Avicel microcrystalline cellulose powder was mechanically
pressed and compacted into 2 inch diameter disks at 10 tons.
Pressed disk samples were placed between flat graphite
sheets to ensure uniform temperature distribution during
heating. Samples were then heat treated in an argon filled
inconel-lined retort furnace to final HTTs using the following
heating schedule:
100 C/h to 250 C, 3 h dwell; 5 C/h to 275 C, 2 h dwell;
5 C/h to 325 C, 2 h dwell; 50 C/h to 450 C, 1 h dwell;
100 C/h to final HTT, 12 h dwell; 100 C/h to 25 C.
Samples for electrical conductivity studies were heat treated to final temperatures ranging from 250 to 1000 C. Samples of higher HTTs of 1200–2000 C were prepared by
further heating the 1000 C retort furnace samples in a hightemperature graphite furnace. Before the high-temperature
treatments, these samples were first outgassed in vacuum
at 1300 C for 12 h (except for 1200 C samples, which were
outgassed at 1000 C instead).
2.2.
Electrical measurements
AC conductivity measurements were performed using an
HP4194A impedance/gain analyzer with the HP16451B dielectric test fixture. Samples were placed in between two 38-mm
electrodes (guarded/unguarded). Samples for AC measurements included microcrystalline cellulose heat treated at final
temperatures of 250–650 C. Parallel conductance was recorded as a function of frequency (ranging from 1000 Hz to
1 MHz) at room temperature.
DC conductivity was measured for samples heat treated
from 650 to 2000 C at several temperatures between 50
and 150 C using a four-point probe apparatus following the
standard procedure for measuring electrical resistivity as described by ASTM C 611-98, ASTM 2003 [28]. Samples were machined and polished to fit specimen dimensions and to
reduce errors from surface imperfections. A Lodestart 8300
variable power supply was used as the current source and
two HP 34401A multimeters were used for current and voltage
measurements.
2.3.
Density measurements
Skeletal density measurements were conducted using the
Accupyc 1330 helium pycnometer for all samples heat treated
from 250 to 2000 C. Since the effect of closed micropores becomes significant for samples heat treated at 650–2000 C,
these samples were first activated in a convection furnace
at 290 C for 24 h before skeletal density measurements to
minimize the effects such closed pores and adsorption.
2.4.
Microstructural characterization
Energy dispersive X-ray spectroscopy (EDS), X-ray photoelectron spectroscopy (XPS), Raman spectroscopy, and highresolution transmission electron microscopy (HR-TEM), were
used to characterize the chemical and microstructural evolution of the precursor with heat treatment. Microcrystalline
cellulose samples heat treated to different temperatures
ranging from 250 C up to 2000 C were analyzed for each of
these studies.
EDS and XPS studies were conducted to observe the chemical changes during carbonization. For EDS, a JEOL JSM-6700F
SEM equipped with the Silicon SUTW-Sapphire EDAX detector
was used. Analysis of the EDS data to extract atomic identification and percent information was performed using the EDS
Genesis software via background subtraction and Gaussian
peak fitting. Microcrystalline cellulose samples that were untreated and heat treated to HTTs between 250 and 600 C were
coated with platinum before testing to reduce surface charge
buildup. Samples heat treated to HTTs between 650 and
2000 C did not require any preparation since they possessed
sufficient electrical conductivity that charge buildup did not
occur.
XPS analyses were conducted on untreated and heat treated microcrystalline cellulose samples. Samples were
mounted onto a sample stub using double-sided copper tape
and placed into an XPS (PHI 5400) analysis chamber
(109 Torr). Samples were subjected to Mg Ka irradiation
(1253.6 eV) generated from a 04-500 dual-anode X-ray source.
During analysis, an electron flood gun (Specs FG20) was used
to compensate for differential charging due to heterogeneities
in the samples’ conductive properties. The kinetic energy of
the flood electrons was varied until the dominant C(1s) peak
was centered at 284.6 eV. All spectra were collected using
an electron energy analyzer operating with pass-energy of
44.75 eV and resolution of 0.125 eV/step; spectra were aligned
by referencing the C(1s) spectral envelope to the CAC/C@C
component at binding energy of 284.6 eV [29,30]. Spectra were
analyzed and deconvolved using spectral analysis software,
CasaXPS. All reported areas were calculated by using sensitivity factors of 0.296 for C(1s) and 0.711 for O(1s) [31].
Raman spectroscopy relies on the interaction of photons
with a sample resulting in scattered radiation that emerge
with different wavelengths compared to the incident beam
[32]. Its sensitivity to lengths, strengths, and arrangement of
bonds in a material make it a popular choice for structural
characterization. A Horiba LabRam HR800 visible Raman system with a 15 mW 633 nm He–Ne laser with excitation line set
to k0 = 514.57 nm and magnification set to ·40 was utilized for all
tests. Detailed scans from 600 to 2500 cm1 were conducted on three
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different spots of each sample. Grams AI software was used to analyze all spectra through background subtraction and peak fitting.
HR-TEM has been used to pictorially describe the microstructural evolution of carbon clusters during the conversion
of microcrystalline cellulose to carbon with increasing HTT.
100 nm thick samples were microtomed and placed onto holy
carbon grids prior to analysis. Images for samples were collected using the HR-TEM CM300 FEG by Philips operating at
297 keV, Gatan Image Filter 200 electron energy-loss spectrometer, and 1 K CCD camera. Images were collected on regions in the holes of the holy carbon grids to avoid
interference from the amorphous carbon film. Energy-filtered
elastic images were collected to reduce inelastically-scattered
electrons and increase image contrast [33].
3.
Results and discussions
3.1.
AC and DC electrical conductivity
AC and DC electrical conductivities of the carbonized microcrystalline cellulose materials are summarized in Fig. 1. Five
distinct regions of electrical conductivity can be identified.
These regions can be correlated with the microstructural
changes that take place during carbonization.
Region I (Fig. 1a) comprises samples heat treated to final
HTTs from 250 to 350 C. The electrical conductivity of this region increases with increase in frequency of applied electric
field. As indicated by Sugimoto and Norimoto [5,12], at this
HTT range, the hydrocarbon is not fully depolymerized and
still possesses high concentrations of hydroxyl and carboxyl
functional groups that dominate the overall electrical con-
1015
ductivity in this region. These organic groups have dipole moments that respond to an applied AC field and such dipole
interaction increases with frequency. The overall AC conductivity of this group of materials decreases as a function of
increasing HTTs since the increase of HTTs leads to the removal of more and more hydroxyl and carboxyl groups from
the original precursor as gaseous products.
In Region II (Fig. 1b), as HTT is progressively increased from
400 to 500 C, the electrical conductivity of the materials is
dominated by the formation and growth of conducting carbon
clusters. Hydroxyl and carboxyl groups, which are responsible
for the electrical conductivity of the Region I materials
through dipole polarization, are now almost fully driven off
due to the higher HTTs. Highly conductive carbon nano-clusters begin to form at HTT of 400 C and the concentration and
size of these conductive clusters grow with the increase of
HTT from 400 to 500 C, explaining the overall rising trend
of conductivity with increasing HTTs.
Region III (Fig. 1c) includes samples heat treated to final
temperatures of 550 and 600 C. In this region, non-linear
electrical conductivity as a function of frequency is clearly observed. Continuous growth of conductive carbon nano-clusters within the dielectric matrix of disordered hydrocarbon
groups gives rise to interfacial polarization under an applied
AC electric field and a quasi-percolated system aiding electron hopping. Sugimoto and Norimoto [5] observed similar results with carbonized wood specimens with HTTs of 500 C,
550 C, and 600 C, proposing that the non-linear increase in
conductivity with frequency was due to the existence of interfacial polarization. At HTT higher than 550 C, no more polar
functional groups or permanent dipoles are left; and instead,
Fig. 1 – Five regions of electrical conductivity measurements and microstructure evolution originated from microcrystalline
cellulose carbonization at final HTTs of 250–2000 C; (a) Region I – dipole polarization (HTT 250–300 C); (b) Region II – linear
AC conductivity, carbon cluster formation and growth (HTT 400–500 C); (c) Region III – non-linear AC conductivity, electron
hopping (HTT 550–600 C); (d) Region IV – DC conductivity, percolation (HTT 650–1000 C); and (e) Region V – DC conductivity,
percolation threshold (HTT 1200–2000 C).
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polarization is caused by the localized movement of free electrons causing dipole moments. An applied AC field induces
polarization across this interfacial phase when the polarization energy barrier is overcome at higher frequencies. Interfacial polarization is a result of the localized movement of free
electrons of the carbon clusters as suggested by [5,11,34–36],
and such frequency dependent polarization across the interfacial phase causes a dipole moment. Maxwell Wagner relaxation observed due to interfacial polarization is caused by a
heterogeneous nature of the material.
As suggested by Barrau et al. [19], electron hopping between conducting nano-clusters is aided by the continuous
development of a percolated system reaching a percolation
threshold. Simultaneous growth of both conducting carbon
clusters and shrinkage of carbon monolithic structure with
increasing HTTs reduces the distance between the carbon
clusters. At this stage these carbon clusters become close enough for electron hopping and quasi-percolation of conducting sites to occur, as predicted by Kercher and Nagle [4,11].
As anticipated, elemental analysis obtained by EDS and
XPS show a decrease in O/C ratio with rising HTTs (shown
in Fig. 2a). A rapid decrease in O/C ratio was observed as HTTs
increased from room temperature to 400 C due to dehydration and depolymerization. The O/C ratio continued to decrease slightly from 400 to 600 C, but remained almost
constant at 95% C as the HTT was increased from 600 to
2000 C. Thus, trace oxygen concentration still remained even
at very high HTTs. These residual oxygen atoms are involved
in the cross-linking of the carbon microstructure, yielding a
non-graphitizing hard carbon [37].
XP spectra of C(1s) regions for untreated microcrystalline
cellulose and heat treated samples up to 2000 C are shown
in Fig. 2b. Four peaks were used to fit the C(1s) region, with
binding energies of 284.6 eV, 286.1 eV, 287.6 eV, and 291.2 eV,
corresponding to CAC/C@C, CAOAC, C@O, and a p–p* shakeup feature [38,39] respectively. Only one peak was used to fit
the graphitic (C@C) and aliphatic (CAC) carbon atoms due to
the close proximity of their binding energies [29].
Extracted peaks appeared are shown in Fig. 2b using consistent fitting protocols for each one of the C(1s) spectral
envelopes. Integrated areas of individual components were
computed to quantify the change in chemical composition
Fig. 2 – (a) Observed and deconvolved C(1s) XP spectra of microcrystalline cellulose samples untreated and heat treated to
temperatures from 250 to 2000 C; (b) The O/C ratios derived from EDS (bulk) and XPS (surface) analysis as a function of HTT;
and (c) calculated component percentages as a function of HTT.
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of microcrystalline cellulose as a function of HTT. Results
from this analysis are shown in Fig. 2c.
As the HTT increased, the conversion of C@O and CAOAC
species in the cellulose starting material to CAC/C@C group is
denoted by the decrease of integrated areas of peaks corresponding to CAOAC and C@O species along with a concomitant increase in intensity for the peak centered at a binding
energy of 284.6 eV due to CAC and C@C groups. The dramatic
decrease of polar CAOAC and C@O groups up to 450 C accounts for the decrease in AC electrical conductivity observed
in Region I.
For samples heat treated to temperatures of 450 C and
above a p–p* shake-up satellite peak appears around
290–294 eV (see insert boxes in Fig. 2b). The appearance of
the p–p* shake-up feature in the C(1s) region is the effect of
polycondensed carbon cluster development that leads to the
formation of a delocalized p electron system. The formation
of such carbon clusters also leads to the increase in the number of p electrons and the increase in both AC and DC electrical conductivity [38] beyond this HTT of 450 C.
AC electrical conductivities of samples with HTTs ranging
from 600 to 650 C were closely analyzed to further study the
insulator to semiconductor transition and the results are
shown in Fig. 3. As Fig. 3 shows, the electrical conductivity becomes frequency independent at HTT of 610 C. However,
these results differed from the gradual transition from frequency dependent to frequency independent AC conductivities observed before and after the percolation threshold for
carbon nanotube polyepoxy [19], carbon anthracites [2], and
carbon-insulator composite systems described [14,16,17], in
which the power law stated in Eq. (6) was followed. AC conductivity of our materials begins to become independent of
frequency abruptly between HTTs of 600–610 C, suggesting
that the percolation threshold volume, uc, of conducting carbon nano-clusters has been reached between these two very
close HTTs. In traditional two-component conducting filler/
insulating matrix composite systems [14,15,21], the arrival
at the percolation threshold is achieved by increasing the concentration of conductive fillers. In our system, the arrival at
percolation threshold is also achieved by the increase of carbon nano-cluster concentration by its continuous growth and
overall volumetric shrinkage with rising HTTs. The increase
of HTTs also leads to the increase of intrinsic electrical conductivities of the conducting phase; while in traditional
two-component systems the intrinsic electrical conductivities of the conductive phase remains constant.
As HTT increases from 600 to 1000 C (Region IV shown in
Fig. 1d), the overall DC conductivity increases by about 5 orders of magnitude as the result of further percolation and rising intrinsic conductivity with HTT increase. As shown in
Fig. 3, the greatest increase of bulk conductivity occurs at
HTT 600–610 C, signaling the arrival of the percolation
threshold and the onset of DC conductivity due to electron
hopping between conductive carbon nano-clusters. The increase in conductivity due to the growth and interconnectivity of the carbon clusters can be described by percolation
theory. Similar results of increasing u with HTT and a sharp
increase in conductivity at uc have been observed for anthracite carbons by Giuntini et al. [2] and Celzard et al. [3]. In
addition to anthracite carbons, composites consisting of
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Fig. 3 – Electrical conductivity measurements near the
percolation threshold, HTT 600, 610 and 650 C.
wood-based carbons [40], carbon blacks [14,41], and graphite
flakes [42,43] within a polymer matrix also showed such conductivity behaviors and these systems were well described by
percolation theory. However, it is important to note that for
our materials, the conductive carbon clusters grow and
intrinsically become more organized with higher HTTs, thus
leading to the increase of the intrinsic conductivity of the
conductive phase.
Figs. 1e and 3 shows that the overall conductivity of our
materials becomes constant with further increase HTTs from
1200 to 2000 C. This indicates that the system has reached
full percolation with the intrinsic conductivity of the carbon
clusters being constant in this HTT region. Although we can
firmly propose the arrival of full percolation at the HTT of
1200 C, we cannot determine the exact HTT range (below
1200 C) where the intrinsic conductivity of the carbon
nano-clusters becomes constant.
3.2.
Microstructural evolution
Raman spectroscopy is a widely used tool to evaluate the
microstructure of carbon materials, particularly the distribution and state of sp2 bonded carbon. In the present study
we postulate that highly conductive clusters, composed of
crystalline sp2 carbon, are embedded in a disordered matrix
of both amorphous sp3 and sp2 carbon. The growth of such
conductive clusters due to rising HTT leads to increases in
electrical and thermal conductivity as observed in Fig. 1 and
in previous studies [44]. Information on the concentration of
crystalline sp2 carbon is thus necessary to describe the formation, growth, and percolation of the conductive carbon clusters. Similarly, Raman characterization can also support the
existence of conductive carbon clusters and help to understand their evolution as a function of HTT.
Raman spectra obtained for microcrystalline cellulose
samples heat treated over the range from 250 to 2000 C are
presented in Fig. 4. A spectrum for untreated microcrystalline
cellulose could not be obtained using the excitation line line
k0 = 514.57 nm due to an overwhelming background signal
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Fig. 4 – (a) Raman spectra of microcrystalline cellulose samples heat treated to temperatures ranging from 300 to 2000 C; (b)
deconvolved Raman spectra; (c) Depiction of the two-dimensional growth of carbon clusters; and (d) I(D)/I(G) ratios as a
function of HTT.
due to fluorescence. Two broad spectral features centered at
Raman shifts of 1350 and 1575 cm1 were observed for heat
treated microcrystalline cellulose. As HTT increases and as
shown in Fig. 4a, both spectral features become sharper and
increase in intensity due to the growth of crystalline sp2 carbon atoms [45–47].
Raman spectra were linear background subtracted by
defining regions between 800 and 2000 cm1 and deconvolved
using five peaks, centered at 1620, 1575, 1500, 1350, and
1100 cm1. Results of the Raman peak fitting using these five
peaks are shown in Fig. 4b. The G-band centered at 1575 cm1
arises from the in-plane vibrations of the sp2-bonded crystallite carbon and has been observed for single crystal graphite
[45,48–50]. For polycrystalline graphite, however, another
peak denoted as the ‘‘disorder’’ peak (or D-band) centered at
1350 cm1 is typically observed. This band is attributed to
in-plane vibrations of sp2 bonded carbon within structural defects [45,50]. Additional, smaller sized Raman peaks centered
at 1620 cm1, between 1500 and 1550 cm1, and near 1100 cm
1
were also used based on their observation in Raman spectra for wood-based carbons [51,52], soot [53], carbon black
[50,54], glassy carbons [55,56], and other graphitic materials
[46,49]. The D 0 -peak at 1620 cm1 is often observed for wellorganized carbonaceous materials with surface p-electrons
[49] while the broad D00 -peak centered between 1500 and
1550 cm1 is associated with an amorphous sp2 carbon
bonded phase. Peaks centered at 1100–1190 cm1 are attributed to sp3 carbon [57]. Lorentzian peaks were chosen to fit
the G-, D, D 0 -, and sp3-band(s) [53,58,59] because of line broadening effects [50,58] while a Gaussian distribution was used
for the fitting of the D00 -band [50,53,58,59].
Analyses of the D- and G-bands can be used to provide
information of the carbon microstructure [47,51,53,60,61].
Crystalline sp2 carbon clusters are often defined in terms of
their average size parallel, La, and perpendicular, Lc, to the orientation of the individual graphene layers as depicted in
Fig. 4c. A parameter often used as a measure of La is the ratio
of the integrated areas of the D-band and G-bands, I(D)/I(G)
[47,48,62,63]. This is a consequence of the fact that G-band
absorption may be regarded as a measure of the average volume of the ordered graphite crystals and the D-band absorption is a measure of the surface area of carbon atoms at the
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edge of graphitic planes [49]. Thus, analysis of the I(D)/I(G) ratio provides a measure of the average crystallite thickness.
The merit of such an approach for providing microstructural
information has been shown explicitly in studies by Tunista
and Koening [48], where the I(D)/I(G) ratio was shown to be inversely related to the crystalline thickness, La, of graphitic
materials.
1019
For heat treated microcrystalline cellulose, the variation in
the I(D)/I(G) ratio as a function of HTT is summarized in Fig. 4d.
An increase of I(D)/I(G) is observed as HTT increases from 300
to 650 C due to depolymerization. The consequence of depolymerization is the formation of isolated sp2 carbon atoms; the
same species that are responsible for the increase in overall
AC conductivity in this temperature range. For increasing
Fig. 5 – HR-TEM images of microcrystalline cellulose samples heat treated to (a) 250 C; (b) 450 C; (c) 600 C; (d) 1000 C; (e)
1500 C; and (f) 2000 C.
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HTTs between 650 and 2000 C, Fig. 4d shows that the I(D)/I(G)
ratio decreases systematically with increasing temperature.
This effect can be ascribed to the conversion of disordered to
ordered sp2 carbon crystallites which increases the crystalline
thickness, La. Therefore, the critical volume for carbon cluster
percolation determined from the sharp increases in thermal
[44] and electrical conductivities observed above 650 C
(Fig. 1) is correlated with the growth of sp2 crystalline carbon.
High-resolution TEM images of microcrystalline cellulose
heat treated to various temperatures between 250 and
2000 C are presented in Fig. 5. For 250 C (Fig. 5a) and 450 C
(Fig. 5b), the precursor resembles an amorphous carbon
phase, typically observed for organic materials. By 600 C
(Fig. 5c), carbon atoms arranged in onion-like graphitic structures [64,65] can be seen, which become more prevalent with
increasing HTTs above 600 C. These onion-like carbon structures are observed with increasing curvature as the HTT increases to 800 C (Fig. 5d) and then to 1000 C (Fig. 5e).
Elongation of the carbon onion layers are observed for samples heat treated to 1500 C (Fig. 5f) and 2000 C (Fig. 5g), consistent with increasing crystalline thicknesses observed in
Raman spectroscopy and XRD studies [44]. Thus in summary,
TEM images show the evolution from amorphous carbon at
low HTT values to a structure that ultimately contains numerous carbon onion clusters at high HTT values.
Previous Raman studies for wood-based carbons have revealed the existence of a small peak between 1100 and
1190 cm1 [51] and have attributed this to sp3 bonded carbon.
In this study, a small peak was observed in this region for
microcrystalline cellulose samples. With increases in HTT,
this peak shifted from 1190 to 1100 cm1 and becomes more
distinctive, due to partial conversion of amorphous to crystalline sp3 carbon for samples heat treated to 1500 C and above
[51]. This is shown explicitly for microcrystalline cellulose
heat treated to 2000 C, by the insert in Fig. 4b. It has been
suggested that the curvature of graphene layers is due to
the inclusion of sp3 crystalline bonds in the graphene layers
[57,66]. Thus, the presence of crystalline sp3 carbon atoms
at higher HTTs is consistent with the existence of the curved
carbon onion structures observed by TEM.
3.3.
Porosity effects
In our carbon materials the connectivity path of carbon conducting nano-clusters among the insulating matrix and the
porosity within the matrix are equally important in determining the overall conductivity. In Fig. 6, electrical conductivity is
shown to increase with increasing skeletal density (decreasing porosity) as expected. As HTT is increased, samples shrink
in volume while conductive carbon nano-clusters form and
grow.
In order to analyze the net contribution to overall electrical
conductivity by the carbon conducting nano-clusters and
insulating matrix, skeletal conductivity is chosen over bulk
conductivity in order to neglect the effects of porosity. Skeletal
conductivity can be calculated based on skeletal and bulk densities as indicated by Kercher and Nagle [4] Ashby et al. [67]:
1:5
r
1 q
2 q
¼
þ
;
ð7Þ
rs
3 qs
3 qs
Fig. 6 – Extrapolated frequency independent electrical
conductivity (r0,dc) of heat treated microcrystalline cellulose
samples as a function of HTT from AC and direct DC
conductivity measurements, d (as shown in Fig. 1); and
calculated skeletal electrical conductivities, m.
where r is the measured bulk conductivity, rs is the skeletal
conductivity, q is the bulk density, and qs is the skeletal conductivity. Bulk and calculated skeletal conductivities as a
function of HTT are shown in Fig. 6. The skeletal density, qs,
is solely dependent upon the densities and volumes of conductive carbon clusters, qc and Vc, and those of the non-conductive amorphous phase, qi and Vi. The total volume of the
skeletal phase, Vtotal, is sum of the volumes of conductive
and insulating phases, Vc, Vi. The volume fraction of the conductive phase, u, can be derived from:
Vtotal ¼ Vc þ Vi ;
q Vc þ qi Vi
;
qs ¼ c
Vtotal
ð8:aÞ
ð8:bÞ
c
with / ¼ VVtotal
, Eq. (8.b) reduces to
/¼
qs qi
:
qc qi
ð8:cÞ
To implement Eq. (8) several assumptions should be taken
into account:
(1) The skeletal density of our materials only depends on
the volume fraction and density of conductive and
non-conductive phases. Since we used a purified form
of cellulose as a precursor, our materials do not have
ashes as found in other forms of carbons such as
anthracites [3].
(2) The density of the amorphous phase, qi, is assumed to
be that of the measured skeletal density of HTT 400 C
sample, 1.37 g/cm3, based on the assumption that at
the of HTT of 400 C, most of the non-carbon elements
have been volatilized and a high purity amorphous carbon has been formed.
(3) The density of the conductive phase, qc, is assumed to
be that of the measured skeletal density of the HTT
2000 C sample, 2.16 g/cm3, based on the assumption
CARBON
4 8 ( 20 1 0 ) 1 0 1 2–10 2 4
Fig. 7 – Electrical conductivity dependence on skeletal He
density.
that at the HTT of 2000 C, the only existent phase is the
conductive phase and the amorphous phase has been
eliminated (see Fig. 7).
Taking all these assumptions into account, the calculated
skeletal conductivity, rs, as a function of u is shown in
Fig. 8. A sharp increase in conductivity occurs when u equals
0.39, corresponding to the HTT of 600 C sample and to the
critical volume fraction, uc. Above this uc value, rs continues
to increase with increasing u, eventually reaching a plateau.
For our system, the increase in conductivity at uc, is not as
sharp as it has been observed for materials in previous studies [3,40–43], where the exponent, t, is reported to have had
values of about 2. The GEM and percolation equations have
been fitted to the skeletal conductivity rs vs. volume fraction
u data and a value of 3.94 resulted for t and these results are
summarized in Fig. 8. A high value of t greater than 2 is
caused by geometrical effects involving the connectivity of
1021
Fig. 8 – Skeletal electrical conductivity (rs) as a function of
volume fraction of conductive phase (u), with GEM equation
regression.
some conductive phases [13,18] and may also be attributed
to the changing intrinsic conductivity of the conductive carbon clusters. The calculated value for t exponent agrees with
those observed for carbon composites by Deramn et al. [68].
3.4.
Energy band gap
The electrical conductivity for samples with HTTs above
600 C is related to the energy gap associated with electron
hopping between conductive sites [1,6,7,23,25] and can be described by the Arrhenius equation:
EHTT
;
ð9Þ
r ¼ r0 exp kT
where the conductivity, r, is a function of an activation energy
associated with each HTT (EHTT), k is the Boltzman constant,
and T is the testing temperature. kT must surpass the energy
Fig. 9 – (a) DC electrical conductivity as a function of inverse testing temperatures for samples with HTTs of 800, 900, 1000,
1200, 1500, and 2000 C; and (b) calculated activation energies from slopes as a function of HTT.
1022
CARBON
4 8 ( 2 0 1 0 ) 1 0 1 2 –1 0 2 4
gap EHTT to observe DC conduction. Electrical conductivities,
r, as a function of testing temperatures, T, are shown in
Fig. 9a. Samples heat treated to 800 and 900 C exhibit
increasing conductivities with T, showing behaviors typical
for semiconductors. For samples heat treated to 1200 C and
higher, electrical conductivities exhibit little variation with
testing temperatures, T.
The activation energy (EHTT) values are calculated from the
slopes of linear regressions of the r vs. T data and are shown
in Fig. 9b. As HTT increases from 800 to 1200 C, the energy
gap decreases from 0.015 to 0.002 eV (agreeing with previous
results observed [1,8] for carbons) as the degree of percolation
of conductive carbon clusters increases. For samples heat
treated to temperatures of 1200–2000 C the calculated energy
gap becomes constant as the system reaches a fully percolated state and as the effects of thermal scattering increases.
Therefore, as shown in Fig. 6, the conductivities reach a plateau for samples with HTTs of 1200–2000 C due to a constant
energy gap and a fully percolated state.
4.
Conclusion
AC and DC electrical conductivity measurements have provided new insights to the decomposition and conversion of
cellulosic materials to carbon. A model is presented that describes the depolymerization of microcrystalline cellulose
and the evolution of the carbon structure with thermal treatment. Five distinct electrical conductivity regions were identified and are directly linked to the microstructure changes
during the cellulose-to-carbon conversion process. In Region
I (HTTs of 250–350 C), AC electrical conductivity is the result
of polarization of organic functional groups and is observed to
decrease with increase in HTT as the cellulose starts to
decompose and lose the polar organic functional groups.
Region II (HTTs of 400, 450 and 500) samples show increasing AC electrical conductivities with rising HTTs due to the
homogenous formation and further growth of highly conducting carbon nano-clusters. Also observed is a linear frequency response that is attributed to the carbon clusters
being sufficiently separated from each other so there is no
interaction. In this stage, the material is described as a system consisting of well dispersed and highly conducting
nano-clusters embedded in a matrix of low conducting amorphous carbon.
In region III (HTTs of 550 and 600 C), a non-linear frequency dependence of AC conductivity is observed as the
conductive nano-clusters grow in size and move sufficiently
close to each other to interact. In this region, electron hopping
and interfacial polarization between the conductive carbon
clusters lead to an order of magnitude increase in the AC
conductivity.
The growth of the conductive carbon clusters continues to
increase the volume fraction of the conductive phase and the
system reaches a percolation threshold at HTT between 600
and 610 C with a critical conductive phase volume fraction
of 0.39, where frequency independent AC conductivity (DC
conductivity) is observed. For samples heat treated to temperatures of 610–1000 C (Region IV), a five order magnitude increase of DC conductivity with rising HTTs is observed due
to percolation and the rise of intrinsic conductivity of the carbon clusters. Finally, in Region V (HTTs from 1200 to 2000 C),
a constant electrical conductivity with rising HTTs is observed due to a fully percolated state and constant intrinsic
carbon nano-cluster conductivities.
Microstuctural and chemical studies were conducted to
understand and correlate the chemical and structure–property relations during carbonization. EDS, XPS, Raman spectroscopy, and HR-TEM were used to follow the chemical
evolution of the precursor and for the identification of carbon
clusters. Chemical analysis indicated a decrease in O/C ratios
on the bulk and on the surface with increasing to 600 C due
to dehydration and depolymerzation, resulting in an increase
in sp2 character and condensation into polycondensed units.
Carbon clusters were thus identified as aggregates of curved
‘‘onion-like’’ carbon structures similar to those observed in
soot and referred as carbon onions.
This research has shown that by using highly pure carbon
precursors such as microcrystalline cellulose and by carefully
controlling the carbonization conditions, a more comprehensive and definitive description can be made for the decomposition and carbonization mechanism of hydrocarbon and
carbohydrate carbon precursors.
Acknowledgments
The authors would like to express our appreciation to the
Department of Energy of the USA for providing funding for
this research under contract number DEFC07-05ID14676. We
would also like to acknowledge David G. Drewry, III for helping with skeletal density measurements and Sun-Hee Park
for helping with resistivity measurements at different testing
temperatures. Finally, we would like to acknowledge the
Material Science Department at Johns Hopkins University
for use of the surface analysis laboratory.
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