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Carbonized Cellulose Nanofibril with Individualized Fibrous Morphology: toward Multifunctional Applications in Polycaprolactone Conductive Composites


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Drying cellulose nanofibril (CNF) from aqueous suspensions often leads to aggregated fibril morphology, negatively affecting its performance in ensuing applications. In this work, we introduced a new solvent drying approach to acquire dry CNF from aqueous suspensions and subsequently pyrolyzed the CNF precursor to obtain carbonized CNF (CCNF) without loss of its fibrous morphology. The fibrous CCNF was dispersed homogeneously in polycaprolactone (PCL) thermoplastic resin, greatly enhancing PCL composite tensile performance. After being further mixed with carbon black (CB), the CCNF helped to minimize CB aggregation due to formation of interconnected three-dimensional (3D) structures. The CCNF/CB/PCL composite exhibited superior electrical conductivity ascribed to electrons transporting more efficiently among CB aggregates. The composite is also suitable for applications such as 3D printed electromagnetic interference (EMI) shielding and deformation sensing. Specifically, the 3D printed EMI shielding composite efficiently absorbed EM radiation in the frequency range of 4–26 GHz, and the 3D printed deformation sensor exhibited excellent sensitivity, durability, and flexibility in monitoring mechanical distortions. Herein, this study sheds light on the development of multifunctional conductive composites embedded with fibrous CCNF from sustainable resources.
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Carbonized Cellulose Nanobril with Individualized Fibrous
Morphology: toward Multifunctional Applications in
Polycaprolactone Conductive Composites
Ju Dong, Xingyan Huang, Guang-Lin Zhao, Jaegyoung Gwon, Won-Jae Youe, and Qinglin Wu*
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ABSTRACT: Drying cellulose nanobril (CNF) from aqueous
suspensions often leads to aggregated bril morphology, negatively
aecting its performance in ensuing applications. In this work, we
introduced a new solvent drying approach to acquire dry CNF from
aqueous suspensions and subsequently pyrolyzed the CNF precursor
to obtain carbonized CNF (CCNF) without loss of its brous
morphology. The brous CCNF was dispersed homogeneously in
polycaprolactone (PCL) thermoplastic resin, greatly enhancing PCL
composite tensile performance. After being further mixed with
carbon black (CB), the CCNF helped to minimize CB aggregation
due to formation of interconnected three-dimensional (3D)
structures. The CCNF/CB/PCL composite exhibited superior
electrical conductivity ascribed to electrons transporting more
eciently among CB aggregates. The composite is also suitable for
applications such as 3D printed electromagnetic interference (EMI) shielding and deformation sensing. Specically, the 3D printed
EMI shielding composite eciently absorbed EM radiation in the frequency range of 426 GHz, and the 3D printed deformation
sensor exhibited excellent sensitivity, durability, and exibility in monitoring mechanical distortions. Herein, this study sheds light on
the development of multifunctional conductive composites embedded with brous CCNF from sustainable resources.
KEYWORDS: cellulose nanobril, composites, 3D printing, EMI shielding, deformation sensing, conductivity
Electrically conductive polymer nanocomposites are being
widely used as key components in batteries,
electromagnetic interference (EMI) shielding,
and electro-
static discharge protection.
Since most polymer matrix
materials are nonconductive, embedded nanollers need to
be superconductive to oer sucient conductivity to the
nanocomposites. Hence, carbon nanomaterials such as carbon
carbon nanotubes,
and graphene
are commonly
used as superior conductive nanollers for nanocomposites.
In particular, carbon nanobers are compatible with both
isotropic and anisotropic polymer matrix materials owing to
the presence of limited functional groups on the ber surface.
Besides, good chemical resistance, high aspect ratio,
and cost
eectiveness also make carbon nanobers popular for the
conductive nanocomposite.
Carbon nanobers grown from hydrocarbon feedstock or
carbon monoxide on the metal catalyst
generally have
controlled ber structure and nanoscale dimensions. How-
ever, large-scale production of carbon nanobers in such a
way is not practical due to the requirement of complicated
equipment and demanding processing conditions. An
alternative method would be similar to the manufacturing
of carbon bers that includes making nanober precursors
through electrospinning, stabilization, and graphitization.
Polyacrylonitrile (PAN),
and cellulosic materials
are three most commonly used carbon ber precursors.
Unlike PAN and pitch that need to undergo electrospinning
to form nanober, cellulose nanobril (CNF) is already in a
ber form with nanoscale diameter and microscale length.
Thus, carbon nanobers might be directly produced from
CNF through stabilization and carbonization. Carbonization
of cellulose macromolecules is actually a thermochemical
decomposition process that includes desorption of free
moisture and bounded moisture, cleavage of cellulose
polymer side groups, session of cellulose polymer backbones,
and graphitization.
During the transformation from organic
cellulose molecules to inorganic carbon, the CNF precursor
Received: March 23, 2021
Accepted: May 7, 2021
© XXXX American Chemical Society A
ACS Appl. Bio Mater. XXXX, XXX, XXXXXX
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needs to remain in a brous morphology at all stages to avoid
material heterogeneity and loss of nanomaterial properties.
Widely used CNF drying methods include freeze-drying,
oven drying,
and supercritical uid drying.
However, all of these methods are reported to have problems
to obtain dry CNF precursors with individualized brous
structures. In this work, we developed a solvent drying
approach to acquire water-free CNF from a concentrated
CNF aqueous suspension. With this novel drying approach,
we were able to prepare dry CNF precursors and carbonized
CNF (CCNF) without loss of individualized brous
morphology. Furthermore, we thoroughly studied fundamen-
tal properties of CCNF and used CCNF as a multifunctional
ller for thermoplastic polycaprolactone (PCL)-based con-
ductive nanocomposites. The CCNF-integrated conductive
nanocomposites were demonstrated to be suitable for EMI
shielding and deformation sensing applications.
2.1. Materials. Thermoplastic PCL pellets (Capa 6800) were
purchased from Perstorp Holding AB (Malmö, Sweden). CNF
suspension (containing 25 wt % CNF) was purchased from Daicel
FineChem (Osaka, Japan). Carbon black (CB-TIMCAL SUPER
C65) was purchased from MTI Corp. (Richmond, CA, USA).
Polyvinylpyrrolidone (PVP, Mw= 40,000) was purchased from
Sigma-Aldrich (St. Louis, MO, USA). Acetone, toluene, and
chloroform were purchased from Hach (Loveland, CO, USA). All
materials were used directly without further purication.
2.2. Preparation of CCNF. Three drying methods including
vacuum oven drying, freeze drying, and solvent drying, were used to
prepare the CNF precursors. Vacuum oven drying was carried out
using a DZF-6020-ETL-110 oven (MTI Corp. Richmond, CA,
USA) at 105 °C for 12 h. Freeze-drying was carried out using
Freezone 4.5(Labconco Corp. Kansas City, MO, USA) for 60 h.
Solvent drying was carried out, as shown in Figure 1. Specically,
the CNF was rst transferred from water suspension to acetone and
subsequently to toluene. The obtained CNFtoluene suspension
was treated with high-intensity ultrasonication using an MSK-USP-
3N-LD ultrasonic processor (MTI Corp. Richmond, CA, USA) for
30 min. Toluene was stripped ovia evaporation at room
temperature to acquire dry CNF. Acetone used in the solvent
drying process was recollected using a Yamato RE 300-AO rotary
evaporator (Yamato Scientic Co., Ltd., Tokyo, Japan).
The CNF precursor was thermally stabilized in a GSL-1100x tube
furnace (MTI Corp., Richmond, CA, USA) at 240 °C in air for 8 h
and then pyrolyzed at 1000 °C in nitrogen for 2 h. The CCNF
properties, including thermal stability, morphology, crystal structure,
elemental composition, and particle size, were characterized using a
thermogravimetric analyzer (TGA Q50, TA Instrument, New Castle,
DE, USA), a eld-emission scanning electron microscope (FEI
QuantaTM 3D FEG dual beam SEM/FIB system, FEI Co.,
Hillsboro, OR, USA), a high-resolution transmission electron
microscope (JEOL 2011, JEOL USA Inc., Peabody, MA, USA),
an X-ray diraction (XRD) device (PANalytical Empyrean
diractometer, Malvern Panalytical Inc. Westborough, MA, USA),
an X-ray photoelectron spectroscope (ESCA 2SR, Scienta Omicron,
Uppsala, Sweden), and a particle size analyzer (Microtrac S3500,
Microtrac Inc., Largo, FL, USA).
2.3. Fabrication and Tensile Testing of CCNF/PCL
Composite Films. The CCNF/PCL composite lms were prepared
by the solution casting method. Tensile testing specimens were
prepared in a dumbbell shape in accordance with the ASTM D638
standard (type V). The tensile test was carried out with a model
5582 Instron machine equipped witha1 kN load cell and Bluehill
software (Instron Inc., Norwood, MA, USA). The tensile speed was
at 10 mm/min with a span distance of 7.62 mm. Five specimens
were tested for each material to obtain an average value. Specimen
fracture surfaces were observed with a Zeiss Axiovert 200 optical
microscope (Carl Zeiss X-ray Microscopy, LLC., Pleasanton, CA)
and scanning electron microscope.
2.4. Preparation of the PVP-Coated CCNF/Carbon Black
Conductive Hybrid. The concentrated CNF suspension and PVP
were mixed in water and stirred for 12 h. Similar to the solvent
drying process for preparing dry CNF, PVP-coated CNF was rst
transferred from water to acetone (Figure 2). CB was then added,
and acetone in the mixture was removed by centrifuging. The PVP/
CNF/CB compound was redispersed in toluene with intensive
ultrasonication. After toluene was completely removed through
evaporation, dry PVP/CNF/CB precursors were obtained, and they
were subsequently stabilized and carbonized to prepare the PVP-
coated CCNF/CB (PVP@CCNF/CB) conductive hybrid (Figure
2.5. Fabrication of Three-Dimensional Printing Conductive
Filaments. Three-dimensional (3D) printing conductive laments
(1.75 mm in diameter) containing various conductive nanollers
(i.e., CB, CCNF/CB, and PVP@CCNF/CB) were extruded at 90
°C using a Filabot EX2 extruder (Filabot, Barre, VT, USA). The as-
extruded laments were chopped to small segments and re-extruded
3 times to ensure that conductive nanollers were dispersed
homogeneously in the PCL matrix.
The four-probe method was used to measure lament electrical
resistivity (ρ)
where lis the lament length (3 cm), Ais the lament cross-
sectional area, and Ris the resistance measured by a Keithley
2100digital multimeter (Tektronix, Inc., Beaverton, OR, USA).
2.6. 3D Printed EMI-Shielding Composites. The EMI-
shielding composites were 3D printed using a FlashForge Creator
Pro printer (Flashforge USA, City of Industry, CA, USA). The
composite has a cylindrical geometry with 1.5 mm inner diameter,
3.5 mm outer diameter, and 2 mm height (Figure S1a). The 3D
printing parameters are listed in Table S2. EMI shielding properties
were measured by a Keysight N5230c vector network analyzer
(Keysight Technologies, Santa Rosa, CA, USA). The coaxial line
method was used for the scattering parameter measurements within
the frequency range of 426 GHz. Compositesouter surface
morphologies were observed with SEM.
2.7. Fabrication of the Deformation-Sensing Composite
Film. Composite electromechanical behaviors were investigated by
measuring specimen real-time electrical resistance as a function of
applied strain. Composite sensing stability was tested on the casted
lm and 3D printed lm. The casted lm was in a rectangular shape.
The 3D printed lm was in a rectangular geometry with a hexagon
array built in the center region (Figure S1b). The 3D printing
parameters are listed in Table S2. Finite element analysis (FEA) was
performed using ANSYS Workbench 2.0 (ANSYS Inc, Canonsburg,
PA, USA) to compare stress distribution on the surfaces. The FEA
parameters (Table S3) were based on the tensile data of neat PCL
lms. The casted and 3D printed lms were meshed with 990 and
198 elements for the numerical analysis, respectively. The simulation
Figure 1. Schematic of the CNF solvent drying process.
Centrifugation was carried out at 7000 rpm for 7 min. CNF
concentration was kept at 3 wt % in water, acetone, and toluene.
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Figure 2. Schematic representation of the preparation of the PVP@CCNF/CB conductive hybrid. CNF and PVP had an equal weight in CNF/
PVP/water suspension. CNF/PVP/CB weight ratio was 1:1:3 in a CNF/PVP/CBacetone suspension. Centrifugation was carried out at 7000
rpm for 10 min. For stabilization and carbonization processes, dry CNF/PVP/CB was rst heated up to 240 °C with a heating rate of 1 °C/
min and then stabilized at 240 °C for 8 h in air, followed by heating up to 1000 °C with a heating rate of 5 °C/min, and then carbonized at
1000 °C for 2 h under the protection of N2.
Figure 3. (a) TGA curves for three types of dried CNF precursors. Insets show that oven-dried CNF and freeze-dried CNF were heavily
agglomerated, while solvent-dried CNF was in the brous form. (b) SEM image of solvent-dried CCNF. (c) CCNF size distribution from 0.1 to
100 μm. Insets show photographs of dispersion state of 0.5 wt % CCNFwater suspensions. (d) Schematic of the proposed CNF solvent drying
mechanism. (e) FTIR spectra of three types of dry CNF. Insets show magnied spectra at 36003100 and 17501550 cm1and (f) XRD
patterns of solvent-dried CNF and CCNF. Insets show deconvolution of the CCNF diraction peak (002) and an HRTEM image of a CCNF
with short-ranged carbon lattice highlighted in red boxes.
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results are presented in normal stress and von Mises stress (Figure
S2) under 3% applied strain (Figure S3) in the longitudinal
Filament specimens used for the cyclic stretchingbending test
had a length in the range of 1015 mm. Filament electromechanical
response to mechanical deformations was evaluated by stretching
and bending the laments with the Instron tensile testing setup 5
times. The maximum tensile and bending strains were set at 3 and
2%, respectively. Electromechanical data including short (i.e., 5
cycles) cyclic stretchingbending, electrical resistance, and long (i.e.,
300 cycles) cyclic stretching were collected within the composite
elastic regime (strain = 3%).
3.1. Fundamental Properties of the Solvent-Dried
CNF Precursor and Carbonized CNF. Figures 3 and 4
show fundamental properties of the CNF precursor and
carbonized CNF. In particular, Figure 3a shows thermogravi-
metric curves for three types of dry CNF precursors. Rapid
sample weight loss between 240 and 390 °C was due to CNF
thermal decomposition.
For solvent-dried CNF, slightly
lower (15 °C) onset degradation temperature and less char
residue indicated its higher sensitivity to temperature as heat
diused in and decomposed the inner CNF easily through
the largely exposed surface area.
This is veried by SEM
images that oven-dried CCNF agglomerated to particulates
(Figure 4a), freeze-dried CCNF agglomerated to akes
(Figure 4b), and both show absence of brous morphologies.
Nevertheless, solvent-dried CCNF preserved its brous
morphology (Figure 3b), showing single ber and entangled
ber bundles in SEM (Figure 4c) and HRTEM images
(Figure 4d). After being dispersed in water (Figure 3c),
solvent-dried CCNF had an average size within the
nanometer scale (<1 μm.), while majority of the freeze-
dried CCNF was in the range of 15μm. The oven-dried
CCNF had the largest sizes and the most widespread size
distribution from 5 to 20 μm.
A schematic drawing (Figure 3d) illustrates the proposed
mechanism for the formation of the brous structure of
solvent-dried CCNF. During solvent drying, the liquid
medium for CNF was changed from water to acetone and
then to toluene.
The CNFwater suspension exhibits an
opaque appearance (Figure 4e) due to the existing CNF
water molecule hydrogen bonds. In acetone, CNF self-
assembled through intrinsic intra- and interhydrogen bonds.
After the CNF was redispersed in toluene under ultra-
sonication, toluene inltrated into CNF molecular chains and
disengaged CNF interhydrogen bonds, resulting in a
transparent appearance. Toluene-occupied spaces that used
to belong to water molecules remained, and owing to the
presence of these free spaces left from removal of toluene,
solvent-dried CCNF attained brous structures instead of
agglomerations. Moreover, high-temperature carbonization
induced the cleavage of CNF side groups, which destroyed
intrahydrogen bonds
and placed furthermore distances
between CCNF backbones. Our hypothesis is supported by
the FTIR results (Figure 3e) that the decreased peak
intensity at 36003150 cm1of solvent-dried CNF
represented diminished CNF self-assembled inter hydrogen
The absorbance bands at 17501680 and 1650
1600 cm1were reported to be attributed to the CO group
from CNF impurities (e.g., pectin) and the hydroxyl group
from CNF-absorbed water molecules, respectively.
Figure 4. (a) SEM images of oven-dried CCNF. (b) SEM images of freeze-dried CCNF. (c,d) SEM and HRTEM images of solvent dried
CCNF, showing individual ber and entangled bers. Yield rates (mass of CCNF/mass of CNF precursors) are 19.0, 17.8, and 17.5%,
respectively. (e) Photographs of CNF dispersion states in dierent mediums during the solvent drying process with the CNF concentration kept
at 3 wt % in water, acetone, and toluene. (f) XPS survey spectrum of solvent-dried CCNF. Inset gure shows deconvolution of the carbon peak.
Inset table compares the C/O ratio between CNF and CCNF.
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decreased intensity of these two peaks indicated that the free
spaces prevented solvent-dried CNF from forming interhy-
drogen bonds with impurities and absorbed moisture.
Three diraction peaks (101), (101̅), and (200) in the
solvent-dried CNF (Figure 3f) represent the transverse
arrangement of the crystallites in cellulose I, while the
diraction peak (040) represents the CNF longitudinal
The newly formed broad diraction peak (002)
in CCNF was a mixture of a sharp crystalline peak and
amorphous halo associated with the graphitized carbon
amorphous carbon,
respectively. This intermediate carbon
phase of CCNF is also observed in the HRTEM image
(Figure 3f inset) in which both disordered and short-ranged
carbon structures are seen along the bers.
The XPS survey spectrum (Figure 4f) shows that CCNF is
mainly composed of carbon and oxygen, while the other
small amounts of coexisting elements such as sodium,
nitrogen, and sulfur are from alkali purication and acid
hydrolysis in CNF preparation. The increased C/O ratio was
ascribed to the loss of oxygen-based functional groups after
carbonization (i.e., dehydration and depolymerization). In
CCNF, oxygen elements cross-linked with the carbon
microstructures (i.e., COandCO), forming the
amorphous carbon region, while the graphitized carbon
region contained concentrated sp2 carbon, which was
reected by the CC peak with a broad and asymmetric
tail toward higher binding energy.
3.2. Tensile Properties of CCNF/PCL Composite
Films. The addition of solvent-dried CCNF signicantly
improved tensile performance for PCL composite lms
(Figure 5a). The maximum tensile toughness (317.0 MJ/
m3), modulus (117.7 MPa), strength (30.5 MPa), and strain
at break (1959.4%) were achieved at a 7 wt % CCNF
loading, which were 2.9, 1.4, 1.9, and 1.9 times higher
compared to the corresponding properties of the neat PCL
Figures 5b and S4acshow a comparison of composite
tensile properties in terms of the CCNF ller type. It is
clearly observed that the solvent-dried CCNF/PCL compo-
site lms had superior tensile properties over both oven-dried
and freeze-dried CCNF/PCL composite lms at the given
loading rates. It also should be noted that PCL composite
obviously benets from the solvent-dried CCNF on tensile
strength enhancement even when compared to some
commonly used nanollers (Table S1).
Improved tensile properties were strongly associated with
the uniform dispersion state of CCNF in the PCL matrix,
mainly due to the fact that solvent-dried CCNF had brous
morphology. It is evident that solvent-dried CCNF/PCL
composite lms show a relatively smooth fracture surface
(Figure 5c,d) with pulled-out bers (Figure 5e), indicating a
strong bermatrix interface. Rough fracture surfaces are
observed for CCNF/PCL composite lms with oven-dried
(Figure 5f) and freeze-dried (Figure 5h) CCNF materials,
showing a poor CCNF dispersion state resulting from ber
agglomerations in the fracture region. Moreover, not only the
oven- and freeze-dried CCNF failed to reinforce the PCL
matrix but also tended to initiate failures at ber-rich regions
(Figure 5gi).
3.3. Conductivity of CCNF/CB Hybrids. CCNF derived
from a low carbonization temperature (<1400 °C) lacks of
sucient electrical conductivity. Hence, a combination with
Figure 5. (a) Tensile stressstrain curves for solvent-dried CCNF/PCL composite lms, with a loading rate from 0 to 7 wt %. (b)
Comparisons on tensile toughness for PCL composite lms containing dierent types of CCNF. (c) Photograph of composite lms after the
tensile test. and (di) fracture surface morphologies of composite lms containing 5 wt % CCNF. Optical images: (d) Solvent-dried CCNF/
PCL. (f) Oven-dried CCNF/PCL. (h) Freeze-dried CCNF/PCL. SEM images: (e) solvent-dried CCNF/PCL. The arrows highlight pulled-out
bers. (g) Oven-dried CCNF/PCL. (i) Freeze-dried CCNF/PCL.
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other conductive materials
is necessary to boost CCNF
In this work, we physically blended CCNF
and CB in an attempt to have CCNFs serving as bridges to
connect CB aggregates (Figure 6a). Furthermore, in the
precursor preparation step, we coated PVP on top of CNFs
through strong hydrogen bonds. After carbonization, both
organic CNF and PVP converted into inorganic carbon. The
XRD patterns show that the PVP-coated CCNF (PVP@
CCNF) (Figure S5a) has a crystallite structure identical to
pure CCNF, indicating that the resulting PVP-derived carbon
was only in an amorphous state.
The HRTEM images
(Figure 6b,c) show that a 17.3 nm thickness of PVP derived
carbon laid on the CCNF continuously, which plays a role as
a binding agent that linked CB particles to CCNF. Therefore,
in addition to PVP@CCNF, a pathway through the linked
CB particles on the CCNF helped electrons to move freely
(Figure 6d), which led to a signicantly enhanced electrical
conductivity for the PVP-coated CCNF/CB hybrid (PVP@
Figure 6e shows a comparison of conductive laments
powering up LEDs in a closed loop (Figure S5b). It is clearly
observed that the PVP@CCNF/CB/PCL lament lighted up
all LEDs, while yellow and green LEDs barely emitted lights
for CB/PCL and CCNF/CB/PCL laments. A conductivity
comparison between the CCNF-based laments and
commercial CB-based laments is shown in Figure 6f. It
should be noted that the electrical resistivity for the PVP@
CCNF/CB/PCL lament signicantly dropped from 1116.44
to 0.0159 Ωm as the conductive ller loading increased from
10 to 30 wt %, making 30 wt % PVP@CCNF/CB/PCL the
most conductive lament. The extraordinary conductivity
indicates a great potential for 3D printing, EMI shielding, and
deformation-sensing composites.
3.4. Composite EMI Shielding Properties. Conductive
composites are able to weaken or completely shield EM
radiation from dielectric loss.
The nature of materials (e.g.,
permittivity), EM radiation characteristics (e.g., frequency),
and geometries (e.g., thickness) all together control EMI
shielding eectiveness.
Figure 7a,b shows composite real (ε) and imaginary (ε)
dielectric permittivity within the frequency range of 426
GHz. As the PCL matrix is nonconductive, εis only
associated with composite interfacial polarization ability,
which is closely related to the ller surface area and free
charge mobility within composites.
Positive εintended the
interfacial polarization arising from the dierence in electrical
conductivities between the ller and matrix.
tended to interact with CB and form a chain-like
interconnected structure,
which further delocalized polarized
charge at the llermatrix interface instead of causing
electron accumulations.
The decreased εand εvalues
with frequency were due to the possible interfacial polar-
izationa phenomenon resulting from insulated PCL,
creating boundaries within the CCNF/CB network.
the dielectric loss tangent (ε/ε) in the frequency range of
415 GHz, ε/εvalues were in the order CCNF/CB/PCL
> PVP@CCNF/CB/PCL > CB/PCL > neat PCL (Figure
7c). At frequencies higher than 15 GHz, the ε/εranking
changed to PVP@CCNF/CB/PCL > CCNF/CB/PCL >
CB/PCL > neat PCL, leading to a higher EM attenuation
factor for PVP@CCNF/CB due to the signicantly increasing
Radiation reection and absorption are two dominant EMI
shielding mechanisms.
Reection occurs when radiation
interacts with shielding-material surface charges.
is a dissipation of EM radiation energy into other internal
EM radiation can be reected multiple times, and
the multireected EM radiation residual is eventually
reabsorbable by the material.
The total EMI shielding
eectiveness (SET) is a sum of eectiveness by absorption
(SEA) and reection (SER), which is expressed as
Figure 6. (a) SEM image showing CCNF connecting CB aggregates in the CCNF/CB hybrid. (b) TEM image of the PVP@CCNF/CB hybrid
structure. (c) TEM image of the PVP-derived carbon layer on the CCNF surface. (d) Schematic of surface charges moving between CB
aggregates in the PVP@CCNF/CB hybrid. (e) 30 wt % conductive llerPCL composite laments power up light-emitting diodes (LEDs). (f)
Resistivity comparison between CCNF-based laments (this work) and commercial CB-based laments.
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R=− − (5)
where PIand PTare the powers of incident and transmitted
EM radiations, respectively.
The reection coecient R, transmission coecient T, and
absorption coecient Aare calculated by
1=− − (8)
Measured Sparameters, S11 [reected (PR) power over
incident power (PI)] and S21 (transmitted power over
incident power) are expressed by
11 R
21 T
Figure 7d shows the composite total shielding eectiveness
in terms of conductive ller types. As expected, neat PCL
barely showed shielding performance within the entire testing
frequency, while the CCNF-added CB/PCL composites
showed much improved total eectiveness. For example,
the maximum SETvalue for the PVP@CCNF/CB/PCL
composite reached above 15 dB at 26 GHz. Composite
absorption eectiveness (Figure 7e) followed the trend of the
total eectiveness; while composite reective eectiveness
was in a bell shape, as shown in Figure 7f, with the maximum
reection reaching 6 GHz. It is clearly observed that
radiation absorption had the major contribution to the total
shielding eectiveness. Limited radiation reection was
partially related to the poor surface morphology of 3D
printed composites. Composite defects such as voids between
adjacent layers (Figure 7g), excessive stringing, and layer
Figure 7. Comparison of EMI shielding properties for composites with dierent conductive llers: (a) ε. (b) ε. (c) Loss tangent. (d) SET. (e)
SEA. (f) SER. (i) Skin depth. SEM surface morphologies of 3D printed EMI shielding cylindrical specimens: (g) Voids between two successive
layers and (h) seams at a single layer.
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seams (Figure 7h) caused multiple EM radiation scattering
and thereby negatively impacted composite reection
The skin depth is dened as the distance up to which EM
radiation attenuates by 1/e or 37%.
It is expressed as
where fis the EM radiation frequency, μ0is the absolute
permeability of free space, which equals 4π×107H/m1,
relative permeability μrequals 1 for the nonmagnetic material,
and σis the composite electrical conductivity. It is known
that the δvalue is inversely proportional to SEA; as a result,
the PVP@CCNF/CB/PCL composite with the minimum δ
value exhibited the maximum absorption, and alternatively,
the CB/PCL composite exhibiting the minimum absorption
had the maximum δvalue.
3.5. Composite Deformation Sensing Properties.
Figure 8ac shows CCNF/CB/PCL lament electrical
resistance measured by cyclic (ve cycles) stretching
bending test at 1.5, 3.0, and 4.5%/min strain rates,
respectively. The ratio of real-time resistance to the initial
value (R/R0) dropped after the rst cycle but reached a
relatively stable state within the rst ve cycles. As the
applied strain recovered to 0, the R/R0value dropped below
1, indicating that the lament electrical resistance was lower
than the initial value, which could be ascribed to the
formation of distortions and cracks during the straining
As expected, the strain rate had no direct impact on
the maximum R/R0value as the electromechanical property
was mainly determined by the geometrical eect.
Due to
the uniform dispersion of conductive llers, structural
integrity of conductive networks was well preserved as long
as stretching maintained in the elastic regime. Filament
electrical resistance exhibited a signicant drop before the
Figure 8. Electrical resistance (R/R0) and FEA results of normal stress distribution on the sensor surface in response to the stretchingbending
force for 20 wt % CCNF/CB/PCL laments. (a)1.5, (b) 3.0, and (c) 4.5%/min1, (d) lm composite, (e) 3D printed composite, (f)
Comparative R/R0stability data for the lm type and 3D printed composites under cyclic tensile test for 500 cycles at a strain rate of 1.5%/
min1. Insets show R/R0values for the 50th and 250th cycles, and (g) sensitivity and reliability to deformations with insets show lm exibility
to bending.
ACS Applied Bio Materials Article
ACS Appl. Bio Mater. XXXX, XXX, XXXXXX
yield point as the CCNF/CB hybrid aligned along with PCL
molecular chains, after which it started to increase following
the geometrical eect and nally reached innity at fracture.
It should be noted that a hysteresis phenomenon was
observed due to the viscoelastic nature of the PCL matrix
(Figure S6).
Furthermore, the performances of the casted lm and 3D
printed honeycomb composites were compared in order to
understand composite geometry eect on the sensing
response to various mechanical deformations. Both types
had the same composition (i.e., 20 wt % CCNF/CB/PCL)
and dimension (i.e., length, width, and thickness), except that
the 3D printed one was lighter in weight due to the presence
of hexagonal holes in the center region. The FEA results
show that under the uniaxial force, the lm type (Figure 8d)
carried load mostly at four corners. On the other hand,
stresses distributed uniformly across the entire honeycomb
structure (Figure 8e), and there was noticeable deformation
upon applied force and quick shape recovery upon the release
of force. This also aligns with the reproducible R/R0values
for the 3D printed composite from cycle to cycle (Figure 8f).
For example, a stable R/R0value of 0.93 was observed for
both the 50th and 250th cycles. Figure 8g also demonstrates
the exibility and durability of the 3D printed composite in
response to bending and its ability to recover to the original
position upon the release of bending force. The 3D printed
composite showed equal sensitivities, regardless of the
bending direction as the R/R0value shifted to the same
level when being bent to the same angle in two directions.
In this work, we introduced a novel solvent drying approach
for acquiring dry CNF precursors. Through solvent drying,
the majority of the CNF self-assembled interhydrogen bonds
were displaced by toluene molecules, which resulted in
sucient free spaces between CNF molecular chains and
brous morphology for dried CNF. The obtained CCNF was
in an intermediate state between graphitized and amorphous
carbons, having both short-ranged and disordered carbon
lattices distributed along a single ber or entangled ber
bundles. The retained brous morphology was essential for
CCNF serving as reinforcing llers to disperse homoge-
neously in the PCL matrix and for remarkably improving
PCL composite lm tensile performance. CCNF also helped
in forming a 3D network when blending with CB, creating
bridges for free charges to move between CB aggregates.
Conductive composites containing CCNF/CB-based llers
(i.e., CCNF/CB and PVP@CCNF/CB) were suitable for
EMI shielding. EM radiation was mainly absorbed owing to
the CCNF interconnected structure eectively delocalizing
free charges. The conductive composites also nd applica-
tions for deformation sensing, which exhibited excellent
stability and durability in response to various types of forces.
sıSupporting Information
The Supporting Information is available free of charge at
Autodesk Fusion 360 designs of 3D printed compo-
sites, equivalent (von Mises) tensile stress distribution
on the surface and the FEA method with 3% tensile
strain applied in the uniaxial direction, comparison of
tensile constants for PCL composite lms containing
dierent types of CCNF, XRD pattern for PVP@
CCNF and the photograph of a closed LED loop
(Figure S5), and stresstime curve under the cyclic
stretchingbending test for 20 wt % CCNF/CB/PCL
laments, a summarized comparison of tensile strength
of CCNF/PCL composite lms (this work) vs PCL
composites lled with other types of reinforcing agents,
3D printing parameters for EMI shielding and
deformation sensing components, and PCL tensile
parameters used for FEA (PDF)
Corresponding Author
Qinglin Wu School of Renewable Natural Resources,
Louisiana State University, Baton Rouge, Louisiana 70803,
United States;;
Ju Dong School of Renewable Natural Resources, Louisiana
State University, Baton Rouge, Louisiana 70803, United
Xingyan Huang School of Renewable Natural Resources,
Louisiana State University, Baton Rouge, Louisiana 70803,
United States; College of Forestry, Sichuan Agricultural
University, Chengdu 611130, China
Guang-Lin Zhao Physics Department and Nano Materials
Laboratory, Southern University and A&M College, Baton
Rouge, Louisiana 70813, United States;
Jaegyoung Gwon Department of Forest Products, National
Institute of Forest Science, Seoul 130-712, Korea
Won-Jae Youe Department of Forest Products, National
Institute of Forest Science, Seoul 130-712, Korea
Complete contact information is available at:
Author Contributions
The manuscript was written through contributions of all
authors. All authors have given approval to the nal version
of the manuscript.
The authors declare no competing nancial interest.
This collaborative study was carried out with support from
the National Institute of Forest Science (Seoul, Korea)
through a cooperative project to the LSU AgCenter and
Louisiana Board of Regents [LEQSF(202023)-RD-B-02 and
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With the speedy advancement of wireless communications, electromagnetic wave (EMW) pollution and electromagnetic interference (EMI) are gradually austere, which makes it an urgent need to develop high-performance EMW absorbers and shields. Carbon-based materials are important members of the family of EMW absorbing and shielding materials, due to their satisfactory electrical conductivity, low density, and good corrosion resistance. Especially, wood biomass-derived carbon materials, including carbonized cellulose, lignin and wood monolith, are acclaimed for their rich conductive network, unique porous structure, and effective loading of functional fillers, which promotes high-efficiency EMW absorption and shielding, and provides a new idea for using wood biomass resources in a high-value way. Additionally, due to the high yield of wood biomass, the mass production of carbonized wood biomass-derived EMW absorbers and EMI shields becomes more promising compared with other carbon materials. Hence, relevant studies should be summarized. Meanwhile, the functions of carbon materials need to be clarified and the importance of carbon materials needs to be highlighted. In this review, the roles of wood biomass-derived carbon in various EMW absorbing and shielding materials were emphasized, along with an analysis of related studies. Meanwhile, the main obstacles and the prospects of wood biomass-derived carbon were proposed.
Present surgical and repair treatments for articular cartilage abnormalities do not produce long-term results that are adequate due to its avascular and alymphatic system. Polycaprolactone (PCL) is a synthetic polyester biomaterial with characteristics of bioresorbable, controlled biodegradation, nontoxic and excellent biocompatibility used for articular cartilage and subchondral bone regeneration. There are various therapeutic potentials for articular cartilage tissue regeneration methods, which have the access to fabricate the replicable biodegradable PCL-based 3D scaffold. Over the past years, natural polymer-based scaffold with poor mechanical properties fails to show significant success in pre-clinical trials, thereby, the mechanical properties of these scaffolds have been enhanced by combining with PCL and it also increases the stability of developed scaffolds. PCL is also being added with natural polymer to increase hydrophilicity and better cell attachment. PCL possesses both semicrystalline and hydrophobic properties ensuring long time support for the scaffold to degrade approximately 3-4 years thereby making it a suitable choice for bone and cartilage tissue engineering involving long-term implants. Some other vital properties of PCL like low melting point, ability to solubilize in organic solvents, heat molding, etc. can be channelized in a proper way to help in recovering the articular cartilage damages.
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Advanced functional materials that are highly efficient in shielding electromagnetic radiation and sensing applications are primarily lightweight polymeric materials. In recent years, several research works on the development of polymer‐based sensors and electromagnetic interference (EMI) shielding materials have been reported. Cellulosic materials have been extensively investigated for fabricating EMI shielding gadgets and sensors. Cellulose is a naturally abundant renewable polymeric material, and the EMI shielding, and sensing performances of cellulose‐based materials depend on their conductive network architecture. Incorporating conducting nanofillers can improve the conductivity of the cellulose matrix in composites. However, a comprehensive understanding of the electrical response of nanofillers in cellulose‐based composites is necessary for the design of EMI shielding materials and sensor devices. Therefore, this work provides a critical overview of the types of processing methods used, an insight into the effects of incorporating conductive nanofillers on the architectural structure of cellulose, and the obtained shielding and sensing properties of the cellulose‐based composites. This article is expected to provide guidelines for developing sustainable polymer materials for advanced applications in the future. This article is protected by copyright. All rights reserved
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Inspired by the ultralight and structurally robust spider webs, flexible nanofibril‐assembled aerogels with intriguing attributes have been designed for achieving promising performances in various applications. Here, conductive polyimide nanofiber (PINF)/MXene composite aerogel with typical “layer‐strut” bracing hierarchical nanofibrous cellular structure has been developed via the freeze‐drying and thermal imidization process. Benefiting from the porous architecture and robust bonding between PINF and MXene, the PINF/MXene composite aerogel exhibits an ultralow density (9.98 mg cm⁻³), intriguing temperature tolerance from ‐50 to 250 °C, superior compressibility and recoverability (up to 90% strain), and excellent fatigue resistance over 1000 cycles. The composite aerogel can be used as a piezoresistive sensor, with an outstanding sensing capacity up to 90% strain (corresponding 85.21 kPa), ultralow detection limit of 0.5% strain (corresponding 0.01 kPa), robust fatigue resistance over 1000 cycles, excellent piezoresistive stability and reproductivity in extremely harsh environments. Furthermore, the composite aerogel also exhibits superior oil/water separation properties such as high adsorption capacity (55.85 to 135.29 g g⁻¹) and stable recyclability due to its hydrophobicity and robust hierarchical porous structure. It is expected that the designed PINF/MXene composite aerogel can supply a new multifunctional platform for human bodily motion/physical signals detection and high‐efficient oil/water separation.
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Interactions of water with cellulose are of both fundamental and technological importance. Here, we characterize the properties of water associated with cellulose using deuterium labeling, neutron scattering and molecular dynamics simulation. Quasi-elastic neutron scattering provided quantitative details about the dynamical relaxation processes that occur and was supported by structural characterization using small-angle neutron scattering and X-ray diffraction. We can unambiguously detect two populations of water associated with cellulose. The first is “non-freezing bound” water that gradually becomes mobile with increasing temperature and can be related to surface water. The second population is consistent with confined water that abruptly becomes mobile at ~260 K, and can be attributed to water that accumulates in the narrow spaces between the microfibrils. Quantitative analysis of the QENS data showed that, at 250 K, the water diffusion coefficient was 0.85 ± 0.04 × 10⁻¹⁰ m²sec⁻¹ and increased to 1.77 ± 0.09 × 10⁻¹⁰ m²sec⁻¹ at 265 K. MD simulations are in excellent agreement with the experiments and support the interpretation that water associated with cellulose exists in two dynamical populations. Our results provide clarity to previous work investigating the states of bound water and provide a new approach for probing water interactions with lignocellulose materials.
Poly(lactic acid) (PLA) and PLA grafted cellulose nanofibers (PLA-g-CNFs) mixture were extruded into filaments, and subsequently 3D printed into composites. As-3D printed composites were then thermally annealed at a temperature above PLA glass transition temperature (Tg). Dynamic mechanical analysis, including temperature ramp, frequency sweep, and creep for annealed composites, confirmed the enhanced responses to various viscoelastic factors. Such enhancements were ascribed to the presence of PLA crystalline regions containing both ɑ and ɑʹ phases, which were induced and developed through the annealing treatment. After 3-point bending test at 70 °C, unannealed composites were partially damaged, while annealed composites preserved the originally well-integrated layer structures. Experimental creep and recovery data essentially fitted to the Burger’s model and Weibull’s distribution function, respectively. The calculated parameters (e.g., moduli) from numerical fitting curves demonstrated the synergetic effect of PLA-g-CNFs and annealing treatment on the enahncement of flexural properties for 3D printed PLA composites.
Filament fused fabrication (FFF) is an extrusion-based 3D printing technology for manufacturing thermoplastic components. One major obstacle facing 3D printed thermoplastic material is the reduced crystallinity resulting from a fast quench when material exiting the 3D printer hot nozzle solidifies quickly at the low-temperature platform, leading to weak mechanical performance. Here, we report an accelerated annealing strategy with the assistance of microwave heating, aiming to enhance crystallinity and mechanical performance of FFF 3D printed polylactic acid (PLA) composite. We selected naturally abundant cellulose fibers as precursors for producing carbonized cellulose nanofibers (CCNFs), and compounded CCNFs with PLA to produce bi-component filament for 3D printing final composite. After being irradiated with microwave, the embedded CCNFs in composite selectively absorbed microwave energy and generated heat. Subsequently, the localized heat transferred to the adjacent PLA regions, triggering amorphous PLA chains to repack and convert to new crystallites. In this work, annealing conditions, including heating method (i.e., oven annealing vs. microwave annealing), time (0–120 min), and temperature (80 vs. 120 °C), were systematically studied to understand the relevant effects on the resulting parameters including composite crystallinity and tensile strength. Microwave annealing method was also compared with conventional oven annealing method and results shows that microwave annealing significantly reduced the required annealing time to reach the maximum crystallinity and tensile strength. Notably, microwave annealing performed below cold crystallization temperature was exceptionally suitable to develop an optimized crystallinity and tensile strength for 3D printed PLA composite.
The electromagnetic interference (EMI) shielding effectiveness of epoxy based composite samples, containing two different sizes of high entropy AlCoCrFeNi alloy powders, was examined under Kα band [26–40 GHz] in this paper. High-energy ball milling processes were used to fabricate AlCoCrFeNi alloy powders of two different sizes, HEAL for larger sizes and HEAS for finer powders, respectively. Comparing with the 8.44 dB value of the HEAL containing sample, the HEAS sample has a maximum total shielding effectiveness (SET) of 20 dB due to the smaller powder sizes and flake-like morphology. Shielding mechanism was studied by resolving the total EMI SET into absorption and reflection portions. Absorption was found to be the major shielding mechanism and reflection had a secondary shielding effect for both HEAL and HEAS cases.
We report the microstructures, electrical conductivity, and electromagnetic interference (EMI) shielding effectiveness of a series of hybrid cellulose papers coated alternatively with silver nanowire (AgNW) and multi-walled carbon nanotube (MWCNT), which are fabricated by controlling the dip-coating sequence and cycle. SEM images and EDS data reveal that AgNWs and/or MWCNTs are sequentially coated on the surfaces of the cellulose papers with increasing the dip-coating cycle and the coating density of the particles decreases gradually in thickness direction of the papers. This result is supported by the anisotropic apparent electrical conductivity of AgNW/MWCNT/cellulose hybrid papers in in-plane and thickness directions. In addition, the apparent electrical conductivity of the hybrid papers in the in-plane direction increases significantly from 0.17–0.22 S/cm to 2.55–2.83 S/cm with increasing the coating cycle from 2 to 10, although it is higher for the hybrid cellulose papers with AgNW top-coating layers than the hybrid papers with MWCNT top-coating layers at the same coating cycle. This result indicates that a highly effective and conductive AgNW/MWCNT network is formed on the cellulose fibers in a layer-by-layer manner. For the hybrid papers with 2.55–2.83 S/cm, high EMI shielding effectiveness of ∼23.8 dB at 1 GHz is achieved.
CNFs) via ring-opening polymerization, forming poly(lactic acid) grafted cellulose nanofibers (PLA-g-CNFs). PLA-g-CNFs and pristine PLA were then blended in chloroform and dried to prepare a master batch. PLA-g-CNFs/PLA composite filaments targeted for 3D printing were produced by compounding the master batch in PLA matrix and melt extrusion. The as-extruded composite filaments were subsequently thermal annealed in a conventional oven, and their morphological, thermal, and mechanical properties were evaluated. PLA was successfully grafted on the surface of CNFs as demonstrated by elemental analysis, and the concentration of grafted PLA was estimated to be 33 wt %. The grafted PLA were highly crystallized, contributing to the growth of crystalline regions of PLA matrix. The incorporation of PLA-g-CNFs improved storage modulus of the composite filaments in both low temperature glassy state and high temperature rubbery state. Postextrusion annealing treatment led to 28 and 63% increases for tensile modulus and strength of the filaments, respectively. Simulated Young's moduli from the Halpin-Tsai and Krenchel models were found comparable with the experimental values. The formed composite filaments are suitable for use in 3D printing. © 2017 Wiley Periodicals, Inc. J. Polym. Sci., Part B: Polym. Phys. 2017