walters2021

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Article
Chiral Nematic Cellulose Nanocrystal/Germania and Carbon/
Germania Composite Aerogels as Supercapacitor Materials
Christopher M. Walters, Gunwant K. Matharu, Wadood Y. Hamad, Erlantz Lizundia,*
and Mark J. MacLachlan*
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ABSTRACT: The development of novel aerogel materials with
chiral nematic ordered structures offers exciting pathways for the
fabrication of multifunctional hybrid materials with enhanced
functionality. Aerogels prepared from cellulose nanocrystals are
especially interesting due to their unique structural properties. To
promote applicability in energy storage materials, it is often
necessary to incorporate metals and metal oxides into threedimensional porous nanostructures. In this study, germania was
incorporated into a chiral nematic cellulose nanocrystal aerogel
using a sol−gel method. Interestingly, our approach does not
disturb the order of the original chiral nematic CNC aerogels,
providing hybrid aerogels with a large concentration of randomly
distributed GeO2 nanoparticles and specific surface areas up to 705
m2 g−1. Carbonization of the composite material afforded a highly ordered material with no collapse during compression and good
shape recovery after release. The combination of the electrochemical double layer capacitance provided by the carbonaceous
skeleton and the pseudocapacitive contribution from the GeO2 nanoparticles resulted in materials with a maximum Cp of 113 F g−1
that exhibited good capacitance retention. To push the boundaries of safer energy storage devices based on renewable resources, we
demonstrate the preparation of an all-cellulose solid-state symmetric supercapacitor.
■
INTRODUCTION
There is a growing need for multifunctional materials made
from inexpensive, readily available renewable resources in an
effort to address rising environmental concerns.1−3 Naturally,
cellulose nanomaterials have become prime candidates for the
development of such materials due to the abundance and
processability of cellulose, as well as the impressive mechanical
properties of these nanomaterials.4−6 In 1949, Rånby
discovered that cellulosic biomass could be treated under
mild acidic conditions to selectively hydrolyze the amorphous
regions of cellulose to yield cellulose nanocrystals (CNCs).7
CNCs are regarded as an attractive nanomaterial owing to their
unique chemical and mechanical properties, which include
easily functionalized surfaces, biocompatibility, high strength,
and low density. Depending on the source and reaction
conditions, prepared CNCs can vary in size (5−20 nm wide
and 100−300 nm long)6 and differ in their surface chemistry
and reactivity. For example, CNCs prepared through sulfuricacid catalyzed hydrolysis have a negative surface charge
imparted through the installation of sulfate half-ester groups
on the surface of the nanorods.8 This surface charge promotes
colloidal stability in water and provides a target for further
chemistry. One intriguing feature of such CNC suspensions is
that they show anisotropic liquid crystalline behavior above a
© 2021 American Chemical Society
critical concentration, forming a left-handed chiral nematic
structure.9,10 This assembly process has been exploited in a
number of ways to create new materials with long-range
hierarchical ordering including photonic films and glass,11−13
stimuli-responsive materials,14,15 mesoporous substrates for
chiral separation and catalysis,16,17 and stretch-responsive
color-changing elastomers,18 to name a few. Ultimately,
CNCs provide a reliable, well understood scaffold for novel
carbon-based mesoporous materials with long-range order.19
Germanium is seeing increased use in electrochemical
storage systems owing to its impressive electrochemical
properties and large theoretical capacity of 1620 mAh g−1
when implemented as a lithium-ion battery (LIB) anode.20
Although germanium and its oxide (GeO2) have been widely
used as electrode materials in LIBs,21,22 very little research has
been done to explore its suitability in supercapacitors.23,24
Usually, metal oxides based on manganese, vanadium, iron, or
Received: April 13, 2021
Revised: May 30, 2021
Published: June 15, 2021
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Article
Table 1. Comparison of CNC/GeO2 Alcogel Formulations and Time Taken for the Prepared Sols to Gel
sample
CNC solution [g]
DMF [g]
H2O:DMFa
TPOG [μL]
TPOG [mmol]
gelation timeb [days]
CNC/GeO2-50
CNC/GeO2-100
CNC/GeO2-200
CNC/GeO2-300
2
2
2
2
4.3
4.1
3.8
3.6
5:95
10:90
15:85
20:80
50
100
200
300
0.17
0.33
0.67
1.0
no gelation after 5 days
no gelation after 5 days
4, soft gel
3, firm gel
a
Ratio of H2O:DMF used relative to the amount of TPOG added to 2 g of phase-separated anisotropic aqueous CNCs. bGelation time refers to the
minimum time required to form a viscous sol.
study provides a possible route toward incorporating
germanium into well ordered carbonaceous materials for
energy storage applications.
Our group has recently shown the potential of mesoporous
chiral nematic structures derived from CNCs to function as
long-life cycle rechargeable lithium-ion battery anodes,46 where
the architecture originating from CNCs provides a robust
physical framework able to accommodate the volume changes
occurring during successive Li+ insertion/extraction processes
(or charging/discharging). On this basis, we decided to explore
the versatility of mesoporous chiral nematic aerogels derived
from CNCs as supercapacitor materials so the functionality of
CNC-based chiral nematic structures as electrochemical energy
storage systems is better understood.
Here, we report the synthesis of mesoporous CNC/GeO2
composite aerogels exhibiting long-range chiral nematic order
with a high surface area. By carefully controlling the hydrolysis
and condensation of the germanium alkoxide precursor, we
were able to grow germania nanoparticles (86 ± 22 nm) inside
a cellulose matrix without aggregation and obtain composite
aerogels. Further carbonization of the composite materials
afforded carbon/GeO2 aerogels that retained their original
chirality and high surface area. Mechanical and liquid
electrolyte uptake studies demonstrate the benefit of having
the germania embedded in a well-ordered carbon support,
while cyclic voltammetry studies suggest these materials
perform well when used as electrodes in symmetric supercapacitor cells. As an added outlook to the future, and to the
best of our knowledge, we prepare the first solid-state
symmetric supercapacitor based entirely on cellulose. Overall,
this work demonstrates a synthetic route that can potentially
be applied to other water-sensitive inorganic precursors,
expanding the horizons for other cellulose/metal oxide-based
aerogel materials for use in energy storage.
cobalt, for example, are preferred due to their reversible
electrochemical redox processes, despite good evidence that
other group IV oxides may be useful.25,26 For example, tin
oxide/graphene nanocomposites have shown a maximum
specific capacitance of 818 F g−1,25 and like germanium
oxide, tin oxide is a wide bandgap n-type semiconductor.27
Although to the best of our knowledge no precedent on the
use of germanium oxide supercapacitors exists, related
materials such as Zn2GeO4 have been proven to show
reversible faradic redox reactions when applied as supercapacitor electrodes.28 Likely, the redox pseudocapacitance of
GeO2 is generally considered unattractive in comparison with
transition metal oxides;29 however, we hypothesize that the
semiconductor characteristics of GeO2 can help to overcome
the constraints seen in many metal oxides, where the poor
electrical conductivity limits the contribution of the bulk
material to the total capacitance (mostly surface contributions
through reversible redox reactions occur).30−33 Additionally,
GeO2 has a relatively high dielectric permittivity, and this
property has been previously shown to enhance the energy
storage capacity of barium titanate-based supercapacitors.34,35
With respect to its mechanical properties, GeO2 is structurally
flexible and has an increased resistance against volume
expansion/shrinkage upon charge/discharge processes, which
gradually deteriorate the electrochemical performance of
energy storage devices.21,22 Enlarging the active surface area
by anchoring GeO2 nanoparticles (NPs) into a mesoporous
carbon support will shorten the ion diffusion path and enhance
the available area to interact with the electrolyte, further
boosting the energy storage ability.36,37 For these reasons, it
would be interesting to explore the applicability of carbonsupported GeO2 in supercapacitors.
Aerogels are perfect candidates for supercapacitors since
they are low-density, lightweight materials with a high porosity
and surface area.38,39 Although monolithic germania aerogels
have been reported, these materials have a relatively low
surface area (155 m2 g−1) and lack long-range order.40
Moreover, because it has been repeatedly shown that Ge/
carbon or GeO2/carbon based electrode materials perform
better than their monolithic counterparts, it is of interest to
incorporate a carbon support, such as an organic polymer or
CNCs for enhanced electrode performance.22,41,42 Despite this,
there are few reports on the preparation of such materials.43,44
One possible reason for this is that commonly used germanium
alkoxide precursors such as tetraethoxygermane (TEOG) or
tetraisopropoxygermane (TPOG) undergo rapid hydrolysis
and condensation, limiting successful incorporation into
mesoporous ordered frameworks.40 In 2018, Xu et al. reported
a sol−gel method for incorporating siloxanes into CNC
suspensions to yield the corresponding chiral nematic CNC/
silica aerogels with surface areas exceeding 700 m2 g−1.45
Although the siloxanes used in that study hydrolyze and
condense much slower than their germanium analogues, the
■
EXPERIMENTAL SECTION
Materials. All chemicals were used as received without further
purification. Aqueous CNC suspensions (4.0 wt %, pH ∼ 2.5) were
supplied by FPInnovations, prepared according to a previously
published procedure.13 N,N-Dimethylformamide (DMF) (HPLC
grade), tetraisopropoxygermane (TPOG) (97%), carboxymethyl
cellulose (CMC) with a Mw of 250 000 g mol−1 and a degree of
substitution of 1.2, sodium sulfate (Na2SO4) (ACS reagent, ≥99.0%),
citric acid (ACS reagent, ≥99.5%), and hydrochloric acid (HCl)
(ACS reagent, 37%) were purchased from Sigma-Aldrich.
Synthesis of CNC Aerogels with Periodic Structure (CNCControl). Following a procedure similar to that of Xu et al.,45 an
aqueous CNC suspension (4 wt %, 15 mL) was allowed to phase
separate in a capped vial at room temperature for 96 h. A portion of
the lower anisotropic phase (2 g) was removed by syringe and
transferred into a 20 mL scintillation vial, which was sealed and stored
in a refrigerator (4 °C) for 48 h to allow the CNCs to reassemble.
Cold ethanol (16 mL, 100%, 4 °C) was slowly layered on top of the
suspension to avoid disrupting the CNCs, and the vial was
subsequently sealed and stored in a refrigerator. After 24 h, half of
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the ethanol was removed by a syringe and carefully replaced with fresh
cold ethanol (4 °C). This step was repeated for a total of 7 days. To
avoid swelling of the gel, HCl (30 μL, 0.12 M) was added to the vial
on days 2−4. The resultant alcogel was clear and was carefully
removed from the vial by scoring and breaking the vial base. The
alcogel was stored in pure ethanol and immediately dried by
supercritical drying.
Synthesis of CNC/GeO2 Aerogels with Periodic Structure
(CNC/GeO2). The phase-separated anisotropic aqueous solution of
CNCs (2 g) was transferred into a 20 mL scintillation vial to which
DMF was slowly added (Table 1). The suspension was swirled until
uniform and was subjected to rotary evaporation (25 mbar, 30 °C) to
remove water until the mass of the CNC solution returned to 2 g.
Once the appropriate ratio was achieved, the vial was sonicated in a
bath sonicator to release accumulated CNCs from the wall of the vial
and then allowed to rest for 1 h. The vial containing the CNC
suspension was then stirred at 700 rpm, and TPOG stored under N2
was added via a glass syringe. The solution was stirred for 15 min and
occasionally agitated to dissolve any white globular precipitate that
formed. The clear, colorless sol was stored in the refrigerator (4 °C)
for 72 h to allow reorganization of the CNCs. As before, cold ethanol
(16 mL, 100%, 4 °C) was slowly layered on top of the sol to prevent
agitation and then sealed and stored in the fridge. After each 24 h
period, half the ethanol was removed via syringe and replaced with
fresh cold ethanol. This step was repeated for a total of 7 days. The
resultant alcogels were translucent, but their opacity increased with
increasing TPOG concentration. The alcogel was stored in pure
ethanol and immediately dried by supercritical drying.
Supercritical CO2 (sc-CO2) Drying of Alcogels. Alcogels were
dried using a Tousimis Autosamdri 815B Critical Point Dryer.
Samples were purged for 5 min and at the beginning of the heating
stage the instrument was turned off for a 1.5 h stasis step. This process
was repeated 2 additional times prior to heating to the supercritical
point. The resultant aerogels were transparent after drying but turned
opaque after a few days of storage under ambient conditions (Figure
S1). This is likely due to index matching of water in the pores as water
is adsorbed from the atmosphere.
Carbonization of Aerogels (C/GeO2). The aerogels were affixed
between two glass slides and secured with a metal wire before being
placed into a XST-3-0-18-1C tube furnace (Thermcraft, Inc.) with a
2404 temperature controller (Eurotherm, Thermcraft, Inc.). Prior to
heating, the atmosphere was purged with Ar for 10 min. All aerogels
were heated under a constant flow of argon gas using the following
program: ramp (25−100 °C, 2 °C min−1), hold (100 °C for 60 min),
ramp (100−200 °C, 2 °C min−1), hold (200 °C for 60 min), ramp
(200−250 °C, 2 °C min−1), hold (250 °C for 7 h), cooling to RT
(under an Ar atmosphere).
Characterization. Scanning Electron Microscopy (SEM). Micrographs in Figure 2B,F were obtained from a S4700 field emission
scanning electron microscope (Hitachi). Samples were prepared by
breaking the aerogel and placing it onto double-sided carbon adhesive
tape loaded onto an aluminum stub. The samples were sputter-coated
with 4 nm of platinum/palladium (80/20). Samples were measured at
a working distance of 5 mm with an accelerating voltage of 5 kV and a
filament current of 10 μA. Micrographs in Figure 2A,D,E,H were
obtained from a S-3400N scanning electron microscope (Hitachi)
following a similar sample preparation method. The samples were
sputter-coated with 10 nm of gold. Samples were measured at an
acceleration voltage of 15 kV, a working distance of 7−9 cm, and a
filament current of 9.4−10.9 μA.
Focused Ion Beam (FIB) SEM. Micrographs in Figure 2C,G were
obtained from a Crossbeam 550 (Carl Zeiss) FIB-SEM. No sputtercoating was done due to the risk of sample modification. Instead, a
cascade of FIB induced carbon depositions was done starting with low
current (10 pA, 50 pA, 100 pA, and finally 300 pA). The pad size was
20 × 10 μm. The probe voltage and current were 30 kV and 300 pA,
respectively, with a milling dose of 400 mC. For serial sectioning and
imaging, the FIB slice thickness was 25 nm. The nanoparticle size was
analyzed using ImageJ. To ensure an accurate nanoparticle size was
estimated, we measured the size of nanoparticles from two separate
Article
images taken at different depths during milling. In each image, 300
nanoparticles were measured as shown in Figure S2.
Powder X-ray Diffraction (PXRD). PXRD patterns were recorded
on a D8 Advance (Bruker) diffractometer equipped with a Cu Kα
sealed tube X-ray source and a NaI scintillation detector. Samples
were mounted onto a zero-background silica plate and secured using
double-sided tape.
Polarized Optical Microscopy (POM). Images were collected on an
Olympus BX41 (Olympus Corp.) optical microscope by removing the
alcogels from an ethanol solution and placing them on a glass slide.
Ethanol was periodically added to the specimen during imaging to
prevent dehydration.
Thermogravimetric Analysis (TGA). TGA was performed on a TG
209 F1 Libra (Netzsch) thermogravimetric analyzer. Samples were
heated from 25 to 800 °C at 10 °C min−1 under air (oxidative
atmosphere).
Nitrogen Sorption (N2 sorption). N2 sorption isotherms were
collected at 77 K on an accelerated surface area and porosimetry
(ASAP) 2020 analyzer (Micromeritics). Prior to analysis, aerogel
samples (50 mg) were degassed at 120 °C for 8 h under the following
conditions: evacuation phase (10 °C min−1, 120 °C, 4 h), heating
phase (10 °C min−1, 120 °C, 4 h). Nitrogen sorption isotherms were
collected and evaluated using Brunauer−Emmett−Teller (BET) and
Barrett−Joyner−Halenda (BJH) methods for surface area and pore
size analysis, respectively.
Mechanical Tests. Properties of the aerogels were evaluated in
compression mode on an AGS-X universal testing machine
(Shimadzu) equipped with a 500 N load cell in displacement control
mode at a rate of 0.5 mm min−1. Cylinder-shaped 3.5 ± 0.5 mm thick
aerogels with ⌀ = 15 ± 1 mm were used. The tests were performed on
either dry aerogels or wet aerogels soaked in an aqueous solution of
Na2SO4 (1 M) for 24 h. The compressive modulus was calculated
from the linear region in the compressive stress−strain curves.
Ultraviolet−Visible (UV−Vis) Spectroscopy. Measurements were
performed with a Cary 60 UV−vis double beam spectrophotometer
(Agilent). Absorbance values within the 200−800 nm range were
obtained using a sampling interval of 1 nm at 200 nm min−1.
Approximately 100 mg of CNC-based aerogels or ca. 45 mg of
carbon-based aerogels were soaked in 5 mL of 1 M Na2SO4 for 24 h
and then removed before measuring the supernatant.
Raman Spectroscopy. Measurements were performed on a
Senterra II Raman microscope (Bruker). Data were collected using
a 532 nm laser at 12.5 mW, with a 5 s integration time. A 20×
objective aperture was used to focus the laser onto the sample. The
spectral baseline was fit to an exponential decay, and the curves were
deconvoluted using a Lorentz fit to extract the peak intensity of the D
and G bands of the carbonized aerogels, as done in the literature.47
Electrochemical Measurements. The supercapacitor behavior of
the aerogels was investigated using cyclic voltammetry (CV)
experiments in a Swagelok two-electrode cell with stainless-steel
current collectors with a VMP3 potentiostat (Bio-Logic Science
Instruments). Symmetric supercapacitor cells were fabricated using
two aerogels of the same formulation (⌀ = 8 mm) separated by a
mesoporous CNC membrane (⌀ = 10 mm) soaked in an aqueous
Na2SO4 solution (1 M) for 24 h. The mesoporous CNC membrane
was fabricated through evaporation-induced self-assembly (EISA)
according to Lizundia et al.48 CV data were collected at scan rates
between 10 and 500 mV s−1 in the 0−0.9 V potential range. The
specific capacitance (Cp, F g−1) was calculated from eq 1
Cp =
∫ I dV
f ΔVm
(1)
where I represents the current intensity (A), f accounts for the scan
rate (V s−1), ΔV is the potential window (V), and m is the aerogel
mass (g).
For solid-state symmetric supercapacitors, a cellulosic-gel was used
as both the separator and the electrolyte. Carboxymethyl cellulose (1
g) was dissolved into Na2SO4 (50 mL, 1 M) with continuous stirring
until the solution became clear (6 h). Citric acid (0.2 g) was added to
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Scheme 1. Formation of the CNC/GeO2 Aerogel Using the Combined Method of CNC Self-Assembly and Sol−Gel
Chemistrya
a
(A) An aqueous suspension of 4 wt % CNCs is added to a vial and allowed to sit until (B) the CNCs phase separate into anisotropic (lower) and
isotropic (upper) layers. The isotropic phase is discarded, and (C) the appropriate ratio of H2O:DMF is obtained through rotary evaporation after
which (D) TPOG is added to the vial and mixed. Over the course of 3 days, TPOG undergoes hydrolysis and condensation to form GeO2
nanoparticles. Next, the solvent is replaced with ethanol, and the resultant alcogel is sc-CO2 dried to yield the resultant (E) CNC/GeO2 aerogel,
which shows a homogenous distribution of GeO2 throughout the chiral nematic cellulose structure.
Figure 1. (A) POM images of prepared CNC-Control, CNC/GeO2-100, and CNC/GeO2-300 alcogels demonstrating characteristic chiral nematic
fingerprint regions. Scale bar is 5 μm. Photographs showing the macroscopic appeareance of (B) CNC-Control (I), CNC/GeO2-50 (II), CNC/
GeO2-100 (III), CNC/GeO2-200 (IV), and CNC/GeO2-300 (V) alcogels. Photographs were taken of the alcogel in a Petri dish over top of the
UBC text (white) on a black background to illustrate the difference in opacity of the samples. (C) CNC/GeO2-300 aerogel and (D) C/GeO2-300
aerogel. The scale bar for (B), (C), and (D) is 2.5 cm.
the solution, and the aerogels were soaked in the Na2SO4−CMC gel
electrolyte for 60 min, picked out, and finally treated at 80 °C for
different times to allow gel formation through double esterification
cross-linking.49 Finally, aerogels were assembled facing each other in
an all-solid-state flexible supercapacitor using a Swagelok twoelectrode cell.
the anisotropic layer could be captured through a sol−gel
method to yield neat cellulose or cellulose/silica aerogels.
Importantly, the CNCs needed time undisturbed to reorganize
into a chiral structure prior to gelation; otherwise, the resultant
aerogel would lack structural ordering.45
In water, tetraisopropoxygermane (TPOG) undergoes the
following hydrolysis (eq 2) and condensation (eqs 3 and 4)
processes:
■
RESULTS AND DISCUSSION
Synthesis. Chiral nematic CNC/GeO2 aerogels were
prepared using a sol−gel method that gave a homogeneous
distribution of germania nanoparticles (NPs) in an ordered
chiral nematic CNC network; the preparation route is shown
in Scheme 1. When left to stand, aqueous suspensions of
CNCs will phase-separate into isotropic (upper) and
anisotropic (lower) phases over the course of a few days,
where the lower phase of the CNCs has a chiral nematic
organization.10,45 Xu et al. demonstrated that the chirality of
Hydrolysis:
Ge(Oi Pr)4 + H 2O F Ge(OH)(Oi Pr)3 + i PrOH
(2)
Condensation:
Ge(OH)(Oi Pr)3 + GeOH(Oi Pr)3
F (Oi Pr)3 GeOGe(Oi Pr)3 + H 2O
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Table 2. Brunauer−Emmett−Teller (BET) Surface Area, Barrett−Joyner−Halenda (BJH) Pore Size and Volume, and
Calculated Density and Porosity for the Prepared Aerogel Materials
sample
CNC-Control
Carbon-Control
CNC/GeO2-50
C/GeO2-50
CNC/GeO2-100
C/GeO2-100
CNC/GeO2-200
C/GeO2-200
CNC/GeO2-300
C/GeO2-300
surface area [m2 g−1] pore size [nm] pore volume [cm3 g−1]
225
518
217
599
211
656
260
705
236
695
14.8
16.5
11.7
15.1
13.7
10.8
11.3
9.6
8.8
10.9
0.80
1.59
0.63
1.61
0.72
1.23
0.72
1.07
0.52
1.30
porosity
[%]
shrinkage after carbonization [%]
0.07
0.09
0.07
0.11
0.06
0.10
0.06
0.13
0.50
0.09
95.2
93.7
95.8
92.9
94.0
94.6
96.8
95.1
97.6
97.0
87
N/A
86
N/A
80
N/A
76
N/A
77
N/A
The alcogel was converted into an aerogel by supercritical
CO2 drying to give a dried product that maintained the pore
characteristics of the original alcogel (Figure 1C). Supercritical
drying was used instead of freeze drying to prevent damaging
the porous macrostructure, which would have resulted in pore
collapse.51 In this work, the aerogels were cylindrical because
the alcogels were prepared in a cylindrical vial. To afford a
material suitable for use in energy storage applications,22,41,42
the aerogels were carbonized (Figure 1D) under a flow of
argon while secured between two glass slides to prevent them
from curling during heating. When temperatures exceeded 300
°C, the samples shrunk more than 75% and became too brittle
to handle. For this reason, carbonization was performed at 250
°C.
PXRD patterns of the aerogels pre- and postcarbonization
(CNC/GeO2 and C/GeO2, respectively) are shown in Figure
S4. In general, there is an increase in intensity of the germania
peaks (PDF: 00-036-1463)52 with TPOG loading. We
estimated the size of the crystallites using the Scherrer
equation53 after deconvoluting the diffraction peak corresponding to the (101) plane (Figure S5) and found them to be
ca. 15 nm prior to carbonization and ca. 25 nm after
carbonization. There is likely annealing and fusion of
crystalline domains at elevated temperature.54 PXRD analysis
of the CNC/GeO2 aerogels also showed diagnostic peaks for
cellulose at ∼18° and ∼22° 2θ, which disappeared following
carbonization. The carbonized samples instead showed a broad
peak between ∼17 and 27° 2θ that is diagnostic of partially
graphitized carbon. This is in contrast to other works where
crystalline graphitic materials were obtained when CNCs were
carbonized at temperatures exceeding 400 °C.55,56 Further
characterization by Raman spectroscopy (Figure S6) shows the
presence of broad D (ca. 1364 cm−1) and G (ca. 1581 cm−1)
bands with a D/G ratio of 0.78, that, when combined with the
lack of a strong diffraction peak at (002), suggests the material
is largely amorphous and lacks graphitic structure.
The surface area and porosity of the composite aerogels
were determined with nitrogen adsorption/desorption isotherms. N2 gas sorption isotherms for all aerogels indicated a
type IV isotherm with an H2 hysteresis loop consistent with
mesoporous materials with cylindrical and spherical pores
(Figure S7).57 The surface areas, pore sizes, pore volumes, and
porosity of the prepared aerogels are reported in Table 2. In
general, the cellulosic aerogels have surface areas between 220
and 260 m2 g−1 and pore sizes and volumes of 9−15 nm and
0.5−0.8 cm3 g−1, respectively. The total pore volume (VT)
(Table S1, Equation S1, and Equation S2) calculated based on
Ge(OH)(Oi Pr)3 + Ge(Oi Pr)4
F (Oi Pr)3 GeOGe(Oi Pr)3 + i PrOH
apparent density [g cm3]
(4)
This process occurred very rapidly when TPOG was added
directly to aqueous CNC suspensions, forming germania
aggregates that precipitated. Our challenge for the synthesis of
these aerogels was to find conditions that enabled control over
the hydrolysis and condensation while simultaneously
maintaining the colloidal stability of the system. Diluting the
aqueous CNC suspension with DMFa well-known drying
control chemical additive for sol−gel processes50could lower
the reactivity of TPOG; however, there is a trade off in CNC
stability. Therefore, ratios of H2O:DMF were judiciously
chosen relative to the concentration of the germanium
precursor to delay the onset of gelation until at least 72 h
after mixing to achieve a chiral structure.
Table 1 lists the gelation time observed after TPOG addition
to suspensions of CNCs in different H2O:DMF ratios. To
illustrate the importance of these ratios, PXRD patterns of
CNC/GeO2-300 prepared from different ratios of H2O:DMF
are shown in Figure S3. At low ratios of water, the germania in
the resultant aerogel was amorphous, but the sol gelled so
quickly that the aerogel had no structural ordering. At high
concentrations of water, the germania in the aerogel is highly
crystalline, but the resultant sol and aerogel contained visible
aggregates. In this study, we focused on preparing aerogels that
maintained chiral nematic order while hosting a homogeneous
distribution of the germania NPs. Suitable H2O:DMF ratios for
TPOG exceeding 1 mM were not found, and all H2O:DMF
ratios led to gelation and precipitation within minutes of
adding the TPOG. As tetraethoxygermane (TEOG) condensed far too quickly, our study only pertains to aerogels
prepared from TPOG.
The prepared CNC/TPOG solutions were kept in a
refrigerator while the CNCs reassembled and the TPOG
hydrolyzed and condensed, at which point the water was
solvent-exchanged with ethanol to afford the corresponding
alcogel. Polarized optical microscopy (POM) of these alcogels
revealed strong birefringence and fingerprint textures characteristic of chiral nematic organization (Figure 1A). The
absence of large aggregates suggested that the germania was
well-dispersed. Photographs of the alcogels (Figure 1B) show
that the alcogels are translucent, but their opacity increases
with the concentration of TPOG. This is likely due to
increased concentration of germania NPs, and not aggregation
as supported by POM.
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Figure 2. SEM images of the (A−D) control and (E−H) high germania loading aerogels. (A) low magnification image of CNC-Control, which is a
neat cellulose aerogel; (B, C) high magnification images of CNC-Control; (D) low magnification image of C-Control; (E) low magnification image
of CNC/GeO2-300; (F, G) high magnification images of CNC/GeO2-300 (GeO2 NPs highlighted with blue circles); and (H) low magnification
image of C/GeO2-300.
Figure 3. Compressive mechanical properties of CNC/GeO2 and C/GeO2 aerogels. (A) compressive stress−strain curves; (B) compressive
modulus; and (C) compressive strength at 85% strain.
composites showed similar degradation onset profiles as CNC
aerogels with decomposition beginning around 220 °C.59 For
CNC/GeO2 composites, the measured loading of germania
was lower than the theoretical loading (Table S2). This loss is
attributed to the solvent exchange steps, where a milky white
supernatant was occasionally observed. After carbonization,
there is an entirely different degradation profile with a main
mass loss step occurring between 350 and 400 °C which
suggests complete conversion of the cellulose structure into
carbon. Additionally, as the concentration of germania
increases so too does the thermal stability of the materials
with degradation occurring near 450 °C for C/GeO2-300. A
slight mass increase at low temperatures was observed in the
C/GeO 2 aerogels (and to a lesser extent CNC/GeO 2
aerogels), which we attribute to possible side reactions and
oxidation occurring in the air atmosphere. The germania
content was notably higher for the carbonized derivatives, with
C/GeO2-200 and C/GeO2-300 showing germania loadings of
ca. 40 and 60 wt %, respectively. Given that a high metal
loading is necessary for energy storage applications, further
discussion will pertain to analogues of those two samples.
Morphology. The chirality of the prepared aerogels was
verified by SEM. Figure 2 shows typical cross sections of the
pure CNC and CNC/GeO2-300 aerogels and their carbonized
derivatives. All cellulosic aerogels have a twisted layer structure
characteristic of left-handed chiral nematic ordering with
the apparent density is much larger than the pore volume (VP)
calculated from the Barrett−Joyner−Halenda, BJH, method.
Since nitrogen sorption can only be used to determine the pore
volumes of pores below 300 nm, the difference between VT
and VP may be attributed to the existence of macropores.
Additionally, a decrease in apparent density is observed with
increasing germania content that may correspond to reduced
shrinkage of the material during sc-CO2 drying of the alcogel
(Table 2). Consequently, this translates into a larger porosity
with increasing germania. After carbonization, there is a large
increase in the surface area of the materials up to 705 m2 g−1.
Those surface areas remain above the results reported for most
of the cellulosic aerogels, which usually fail to obtain values
exceeding 600 m2 g−1.58 Likely, there is increased accessibility
of the pores after CNC degradation. In addition, there was a
decrease in VT along with an increase in micropore and
mesopore volume. This suggests that although the macrostructure shrinks during the carbonization process, the carbon/
germania backbone prevents collapse of the porous structure,
resulting in a substantially higher surface area for the C/GeO2300 aerogels when compared to the Carbon-Control aerogel
(695 m2 g−1 versus 518 m2 g−1). Samples were analyzed by
TGA under air to determine the amount of germanium oxide
present in the materials (Figure S8, Table S2). Although the
germanium oxide is stable over 800 °C, its presence does not
significantly affect the thermal stability of the CNCs, and the
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pitches of ca. 3 μm for all aerogels (Figure 2A,E). Importantly,
this chirality is maintained even after carbonization for all
structures (Figure 2D,H) despite shrinkage of the bulk aerogel
by up to ca. 87% (the pitch is ca. 2 μm for all carbonized
samples). The aerogels also have a vast porous network
(Figure 2C), which is important for electrolyte uptake.
Germania NPs are well-dispersed throughout the chiral
network of CNCs and appear to be 86 ± 22 nm in size
(Figure 2G, Figure S2, Video 1). Although some larger
microstructures were observed, the majority of germania is
indeed nanosized. This is an important finding that hints at the
utility of these composite materials for use as electrode
materials in supercapacitors as the electrolyte may easily access
the pseudocapacitive surface of germania nanoparticles.
Mechanical Properties. Determining the compressive
properties of the prepared aerogels provides insight into their
capacity to withstand applied external forces when used in
supercapacitor applications. Accordingly, representative uniaxial compressive stress−strain curves of CNC/GeO2 and C/
GeO2 aerogels are shown in Figure 3A, while the compressive
modulus and strength at 85% compression are summarized in
Figure 3B,C, respectively. The compressive stress−strain
curves are characterized by an elastic behavior at strains
below 10% (the compressive stress increases linearly upon
strain application) followed by a plastic deformation region
and a final stiffening region at larger strains (>25% for C/GeO2
aerogels and >55% for CNC/GeO2 aerogels).60 Upon
compression, all aerogels showed a good deformability,
accommodating up to 95% strain without breaking or cracking.
Compressive modulus values of 180−220 kPa for cellulosic
aerogels and 350−400 kPa for carbonaceous aerogels were
achieved. The larger modulus observed for carbonized aerogels
indicates the formation of a mechanically robust threedimensional structure upon carbonization. This may also
explain why an increased surface area and porosity were
observed following carbonization despite shrinkage of the bulk
material. Similarly, C/GeO2 aerogels show a substantially
improved compressive strength at 85% strain when compared
to CNC/GeO2 aerogels (ca. 3200 kPa versus ca. 1300 kPa)
(see Figure S9). Remarkably, compressive modulus and
strength values of CNC/GeO2 aerogels remain well above
the results reported for other CNC-based aerogels reported in
the literature, which range from 7 to 41 kPa at 80% strain.61
There are two factors that contribute to this improved
resistance to deformation. First, as opposed to freeze-drying
the aerogels, which is a commonly used approach for aerogel
preparation, the aerogels prepared here were obtained through
sc-CO2 drying, thereby avoiding pore collapse and keeping the
macrostructure of the original wet gel intact.51 Second, our
approach yields a highly organized chiral structure with longrange order, which provides enhanced specific strength and
toughness. This has been explored by Tripathi et al., who
prepared both isotropic cellulose aerogels and chiral nematic
cellulose aerogels (following the same method we used) and
compared the mechanical properties.62 There was a very clear
benefit to the hierarchical ordering, which gave an enhanced
specific strength and specific toughness when compared to the
isotropic aerogels. Consequently, these enhanced mechanical
properties are important for maintaining the porosity and
macrostructure of the aerogel during both carbonization and
electrochemical testing.
To further understand the deformation process of the chiral
carbonaceous aerogels, which are especially interesting for
Article
energy storage applications, we investigated the compression in
situ by SEM. We built a custom sample holder comprising two
parallel plates where the aerogel thickness was controlled using
four screws (Figure S10A,B). The same region of a given
aerogel was observed by SEM before and after 30%
compression relative to its starting thickness. The micrographs
of C/GeO2-300 (Figure S10C,D) show that the original chiral
morphology was maintained, with small changes in the helical
pitch spacing. As the disordered regions are preferentially
deformed, the chiral areas are able to withstand applied
external forces with no significant deformation. These results
suggest that, similar to many biological systems featuring a
twisted plywood structure,63 the hierarchical nanoscale
structuring of the aerogels is responsible for the high
compressive modulus observed.
The highly porous character of aerogels suggests they may
be useful for electrochemical energy storage devices.36 Ideally,
electrochemical energy storage devices should have a strong
affinity toward the liquid electrolyte to provide a good
compatibility with conducting media. We immersed the
aerogels into a 1 M Na2SO4 electrolyte, observing that the
liquid was rapidly absorbed and retained by the aerogel. We
defined the electrolyte uptake (EU) after immersing aerogels
in 1 M Na2SO4 for 24 h as
EU =
100
× (mwet − mdry)
mdry
(5)
where mwet and mdry are the mass of the wet and dry aerogels,
respectively. As depicted in Figure S11A, EU values up to 956
and 585% were obtained for CNC-Control and CarbonControl, respectively. Interestingly, these high EU values were
obtained despite the predominantly hydrophobic character of
the carbonized CNCs. Although the presence of germania
slightly lowers the EU due to a reduction of the liquid-holding
phase, the obtained results are well above the EU values
reported so far for most other aerogel systems including
cellulose aerogel membranes (615%),64 bacterial cellulose/
Li0.33La0.557TiO3 aerogel (586%),65 and cellulose/polyacrylamide aerogels (548%).66 The enhanced EU is explained by the
fact that the hierarchical, interconnected pore structure of
aerogels provides a reduced surface tension and enhanced
capillary flow of the electrolyte within the solid67 and that scCO2 drying also helps mitigate pore collapse. Additionally, the
oxygen-containing functional groups of CNCs offer a good
affinity with 1 M Na2SO4.68 UV−vis measurements of the
liquid electrolyte after aerogel soaking were conducted to
check the stability of the prepared aerogels in the liquid
electrolyte. UV−vis spectra of the resultant solutions are
shown in Figure S11B. Although the CNC/GeO2 aerogels
released a small amount of CNCs into the electrolyte, as
denoted by the absorption maxima in the 250−310 nm range,
the aerogels did not redisperse in the solvent. This is in
contrast to early reports of nanocellulose aerogels, which could
be redispersed unless they underwent additional chemical
cross-linking.69
Additionally, the UV−vis spectra of the C/GeO2 aerogels
remain barely modified, suggesting that the carbonization step
provides a long-lasting physical structure where aerogel
components (both carbonized CNCs and GeO2) remain
tightly bound. A broad, weakly absorbing species at 275 nm
can be seen in the aerogel samples containing germania. It is
possible this could be from unbound GeO2 leaching from the
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Figure 4. Compressive mechanical properties of electrolyte-soaked CNC/GeO2 and C/GeO2 aerogels. (A) Compressive stress−strain curves at
60% compression for the 1st and 10th cycle for (i) CNC/GeO2-300, (ii) CNC/GeO2-200 (ii), (iii) CNC-Control, (iv) C/GeO2-300, (v) C/GeO2200, and (vi) C-Control aerogels. Mechanical properties from the obtained stress−strain curves: (B) compressive modulus and (C) compressive
strength at 60% compression for the 1st and 10th cycles.
aerogels suffer from a marked compressive strength decrease
upon cycling (up to 45%), while carbon-based aerogels kept
their compressive strength strain almost unchanged (<10%).
Energy loss coefficient (η) values (further discussion in Figure
S14) ranged from 0.7 to 0.9, indicating a stable threedimensional structure provided by the highly porous carbonaceous skeleton, yielding aerogels able to accommodate external
stresses without pore collapse (aerogels did not break apart but
recovered their original shape upon release, Figure S15.51 A
priori, the good mechanical adaptability of these aerogels,
which shows no collapse during compression and shaperecovery after release, may result in an improved adhesion and
interfacial compatibility when assembled into a supercapacitor,
which is a prime requirement for energy storage devices having
optimal charge transfer between negative and positive
electrodes. This is facilitated by the hierarchical ordering of
the cellulose which offers enhanced mechanical properties
when compared to its isotropic counterpart, and when coupled
with the porous morphology, large specific surface areas, and
large electrolyte uptake capacity, the developed aerogels are
good candidates to sustain local stresses during supercapacitor
applications.
Electrochemical Performance. In contrast to brittle silica
or early reports of CNC aerogels,69,72 the fabricated aerogels
are mechanically strong and keep their structure after
compression when soaked in a liquid electrolyte. Such
properties are promising for use of these materials in
aerogels; however, we have not been able to accurately identify
this species.
The compressive response of electrolyte-soaked aerogels was
studied to assess their mechanical flexibility and ability to
withstand multiple loading−unloading cycles without damage
(Figure S12). Aerogels soaked in 1 M Na2SO4 were
compressed to 60% of their original thickness and then
immediately released for 10 successive cycles. Figure 4A shows
the representative compression stress−strain curves for the 1st
and 10th cycles of the synthesized aerogels (see Figure S13 for
all 10 cycles). Generally, all materials showed a continuous
compressive stress increase upon loading, but the curves
followed different decompression pathways likely due to some
structural collapse. To our surprise, C/GeO2 aerogels showed
comparable and even improved flexibility and shape recovery
performance, in spite of the generally brittle character of the
carbonized fibers.70 While CNC/GeO2 aerogels presented a
marked hysteresis and have a residual strain of 50% (strain at
which 0 kPa is reached during unloading), C/GeO2 aerogels
show a lower residual strain. This behavior indicates a better
shape recovery of the carbonaceous aerogels with no plastic
deformation.71 The compressive modulus varied from ca. 7
kPa for CNC/GeO2 aerogels up to ca. 40 kPa for the
carbonaceous analogues (Figure 4B). Similarly, the maximum
strength during the first compressive cycle remained in the
10−60 kPa and 330−650 kPa ranges for CNC/GeO2 and C/
GeO2 aerogels, respectively. As shown in Figure 4C, CNC5204
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Figure 5. (A) Cyclic voltammograms of aerogels in 1 M Na2SO4 measured at a scan rate of 100 mV s−1; (B) cyclic voltammograms of the C/GeO2300 aerogel in 1 M Na2SO4 for scan rates varying between 10 and 500 mV s−1; (C) gravimetric rate performance of aerogels in 1 M Na2SO4; and
(D) cyclic voltammograms of the C/GeO2-200 aerogel after the 1st, 5th, 10th, 20th, 50th, 100th, and 300th cycles.
200, suggesting the importance of having a high germania
loading. The CV response of C/GeO2-300 at scan rates
between 10 and 500 mV s−1 is summarized in Figure 5B (see
Figure S17 for the rate performance of all prepared aerogels).
In general, there is an increase in the current with increasing
scan rate, which demonstrates good capacitive behavior in the
studied voltage window. Further, the area within the CV curves
increases with the applied scan rate, and the curves keep their
shapes, indicating low ionic transfer resistance at high current
rates.77 Figure 5C summarizes the gravimetric rate performance (Cp specific capacitance variation as a function of applied
current rate) for all the studied compositions (reported Cp is
normalized to the total aerogel mass). It was observed that
aerogels containing large amounts of GeO2 present an
increased specific capacity for a given current rate, while
carbonization increases the charge storage capacity of CNCbased aerogels. These findings confirm the initial hypothesis
that the combination of a highly porous hierarchical structure
and the presence of GeO2 NPs hosted within an electrically
conducting carbonaceous skeleton result in improved energy
storage performance. A maximum Cp of 113 F g−1 was
achieved for the C/GeO2-300 aerogel at a scan rate of 10 mV
s−1. In contrast, neat carbon aerogels had a Cp of 44 F g−1,
while the neat CNC aerogel had a Cp of 36 F g−1 (similar to
the 34 F g−1 shown by cellulose nanofiber aerogels).28
Although capacities up to 328 F g−1 have been reported for
porous carbon aerogels from cellulose upon a carbonizing−
activating process,78 the results obtained here are well above
the 43.7 F g−1 measured for symmetric CNC/polypyrrole
supercapacitors at 2 mV s−179 or the 90 F g−1 obtained for
poly(3,4-ethylenedioxythiophene)/nanocellulose electrodes at
1 mA cm−2.80 The good rate stability of our materials arises
from the large specific surface area of the aerogels (up to 700
symmetric supercapacitor applications. By using a mesoporous
cellulose membrane (see Figure S16), which has already shown
good performance in LIBs, in conjunction with the cellulosebased aerogels, we report for the first time a supercapacitor
device originating entirely from cellulose.73
The electrochemical properties of aerogels are shown in
Figure 5. Cyclic voltammetry (CV) scans of the symmetric
supercapacitors containing two identical aerogels separated by
the mesoporous CNC membrane soaked in 1 M Na2SO4 are
shown in Figure 5A (no binders or conducting additives were
incorporated). The curve shape varies depending on the
sample, but in general, the aerogels display symmetrical
rectangular-shaped curves indicative of electrochemical double-layer capacitors (EDLCs, based on ion adsorption/
desorption).74 C/GeO2-300 presents a distorted symmetrical
quasi-rectangular shape related to a pseudocapacitance
behavior provided by the redox reactions occurring at the
GeO2−carbon interfaces.75 It can be observed that this feature
is absent for CNC/GeO2-300. We postulate that the different
pore structures achieved during synthesis may be the
underlying reason. Indeed, the BET surface area of C/GeO2300 is 695 cm3 g−1 as opposed to 236 cm3 g−1 for the CNC/
GeO2-300 aerogel. In order to achieve the pseudocapacitive
contribution, it is necessary to guarantee the exposure of GeO2
to the liquid electrolyte ions.76 Unfortunately, the relatively
poor surface area of the CNC/GeO2-300 aerogel denotes a
pore clogging effect, reducing the effective GeO2 area exposed
to accumulate charges and thus lowering the overall Cp.
Thanks to the combination of a pseudocapacitive character
with EDLC behavior, the electrochemical performance of C/
GeO2 is effectively enhanced,74 as indicated by the increased
area within the CV curve. It is also important to note that this
quasi-rectangular shape is significantly attenuated for C/GeO25205
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m2 g−1), allowing for enhanced electrolyte accessibility of that
active material, which in turn enables rapid ion diffusion across
the electrodes and increases the amount of material capable of
undergoing redox reactions.36 Moreover, the mesopores
function as electrolyte nanoreservoirs and act as electrolyte
ion diffusion channels, further enhancing the supercapacitor
performance.81
Cell cycling durability is one of the most relevant parameters
determining the practical implementation of supercapacitor
electrodes. Extensive cycling often results in undesired large
volume changes of the carbon-based electrodes, which in turn
reduces the overall Cp of the material over time.82 Therefore,
mitigating the mechanical degradation and the associated
capacity fading in supercapacitors is critical to obtaining
electrochemical energy storage systems with long-life cycles. In
this framework, the incorporation of inorganic nanoparticles
into 3D hosts has been proven as a suitable way to protect the
electrode structure from collapse by buffering the volume
changes of the skeleton.83 Accordingly, we assessed the cycling
stability of the germania-containing aerogel further at a scan
rate of 50 mV s−1 (Figure 5D, see Figure S18 for all aerogels).
No shape change was observed, and the aerogel kept adequate
charge-storage characteristics for the whole 300 cycles. This is
in contrast to the neat carbon aerogel (Carbon-Control, Figure
S18F), which showed continuous capacity decay upon cycling.
In this context, Figure S19 summarizes the Cp decay after 300
cycles for all of the synthesized aerogels. Generally, the
aerogels show a decrease in Cp, with the exception of CNC/
GeO2-300 which shows a 5% Cp increase after 20 cycles, prior
to decaying. Such behavior is explained by the relatively poor
mechanical properties of CNC aerogels that allow for
deformation upon cycling to provide additional active GeO2
surfaces. Overall, carbonization leads to aerogels capable of
holding their initial specific capacity to a larger extent; i.e., the
neat CNC aerogel keeps ca. 77% of its initial Cp after 300
successive charge/discharge cycles, while the neat carbonized
aerogel keeps ca. 83% of its Cp. Remarkably, both in the CNCand carbon-based aerogels, the presence of GeO2 extends the
lifetime of the supercapacitors, resulting in enhanced capacity
retention values of 91% and 96% for the CNC/GeO2-200 and
C/GeO2-200 aerogels, respectively (as opposed to the 77%
and 83% achieved for the monolithic CNC and carbon
aerogels). Those results are in contrast with the poor cycling
stability of pseudocapacitive polymer electrodes, which suffer
from structural breakdown during repeated charge/discharge
cycles.84 These results prove that the formation of GeO2
nanoparticles within the chiral nematic structure through a
sol−gel approach protects both the CNC- and the carbonbased aerogels from undesired structural deterioration during
cycling, reducing the capacity fade observed in the control
(neat CNC and carbon) samples. This finding has also been
observed during mechanical testing in both dry and electrolytesoaked states (Figure 3 and Figure 4), where aerogels
containing GeO2 present an enhanced resistance to deformation (seen as a larger compressive modulus and strength). In
this way, GeO2 prevents aerogel stacking and specific surface
area loss during charge/discharge cycles, promoting efficient
ion transport and keeping decent capacities after many cycles.
Based on these promising findings, we constructed a solidstate supercapacitor (SSC) using a gel electrolyte that
combines ionically conducting media with liquid-leakage-free
characteristics.85 As opposed to conventional supercapacitors
relying on aqueous or organic-based liquid electrolytes, the
Article
SSC avoids electrolyte leakage issues and can achieve larger
energy densities (i.e., lighter devices for a given delivered
energy). Nowadays, poly(vinyl alcohol) is a popular choice for
SSCs because of its good gelling ability and chemical/
mechanical stability.86 We replaced poly(vinyl alcohol) with
a cellulosic gel electrolyte to construct a solid-state symmetric
supercapacitor. Although the performance of solid-state energy
storage devices remains below that of traditional systems due
to limited ion transfer at electrode−electrolyte interfaces and
reduced ionic conductivities, their improved safety makes them
attractive alternatives to liquid systems.87 To the best of our
knowledge, this is the first example showing the development
of a solid-state symmetric supercapacitor fully based on
renewable resources (modern supercapacitors require petroleum-derived materials soaked in a liquid electrolyte36 or
poly(vinyl alcohol)/Na2SO4 gels).74 The electrochemical
characterization of the solid-state symmetric supercapacitor
comprised of C/GeO2 aerogels is shown in Figure S20. We
found that the gel electrolyte synthesis plays a pivotal role in
ensuring an efficient ion transfer between aerogel electrodes.
Long thermal treatments result in extensive gel cross-linking,
which in turn limits the ion migration process within the gel
electrolyte.88 As a result, CV curves with extremely low current
densities were obtained (Figure S20A). Limiting the double
esterification cross-linking to 5 min ensures an adequate
balance between cross-linking (supercapacitor safety, electrically insulating both electrodes) and ionic conductivity. As a
consequence, CV curves in Figure S20B are characterized by a
combination of an EDLC behavior of carbon (ion adsorption/
desorption) with the pseudocapacitance of germania (reversible redox reactions), providing a Cp of 29 F g−1 at 10 mV s−1
(14 F g−1 at 500 mV s−1). That said, there is still room for
improvement as obtained Cp values can be further increased by
optimizing the synthesis of the cellulose gel. Similar to
supercapacitors having a liquid electrolyte, the prepared solidstate symmetric supercapacitors show long-term stability,
holding 90% of their initial capacity after 250 cycles at 50
mV s−1 (Figure S20C,D). Those results reveal the suitability of
binder-free high-surface-area carbon/germania aerogels for
energy storage applications, including symmetric solid-state
supercapacitors based on renewable resources, of which scarce
examples exist in the literature.
■
CONCLUSIONS
We report the development of high surface area cellulose and
carbon germania aerogel composite materials with high
porosity and chiral nematic order for use in supercapacitors.
In preparing these materials, we demonstrated a simple
method for incorporating highly reactive precursors into
carbonaceous supports while retaining long-range order. Indepth compressive mechanical measurements demonstrated
the utility of having a chiral nematic structure in the aerogel,
which prevents collapse of the macrostructure during carbonization as well as during electrochemical testing. This is
important for maintaining a high porosity and surface of the
material which ultimately results in high electrolyte uptake
(>500%). The ordered chiral nematic structure also serves as a
pathway for efficient electron transport within the whole 3D
skeleton, which is vital for electrochemical applications.
Preparation of symmetric supercapacitors based entirely on
CNCs showed good specific capacitance (ca. 114 F g−1) and
retention of up to 95% after 300 cycles at a scan rate of 50 mV
s−1. The presence of GeO2 not only increased the specific
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HORSE), Bioimaging Facility (BIF), and Shared Instrument
Facility (SIF) for assistance with XRD, SEM, and TGA,
respectively. The authors would also like to acknowledge Dr.
Andreas Schertel and Carl Zeiss Microscopy GmbH for FIBSEM measurements and Dr. Andy Tran for designing the
graphics in Scheme 1. The authors are grateful for the Open
Access funding provided by the University of Basque Country
(UPV/EHU).
capacitance of the carbon aerogels by combining the inherent
EDLC behavior of carbon (ion adsorption/desorption) with
the pseudocapacitance of germania (reversible redox reactions)
but also resulted in symmetric supercapacitors with an
enhanced life span. These results reveal the suitability of
binder-free high-surface-area germania/carbon aerogels for
energy storage applications, including symmetric solid-state
supercapacitors based on renewable resources.
■
■
ASSOCIATED CONTENT
sı Supporting Information
*
REFERENCES
(1) Lizundia, E.; Kundu, D. Advances in Natural Biopolymer-Based
Electrolytes and Separators for Battery Applications. Adv. Funct.
Mater. 2021, 31, 2005646.
(2) Berglund, L. A.; Burgert, I. Bioinspired Wood Nanotechnology
for Functional Materials. Adv. Mater. 2018, 30, 1704285.
(3) Li, T.; Chen, C.; Brozena, A. H.; Zhu, J. Y.; Xu, L.; Driemeier,
C.; Dai, J.; Rojas, O. J.; Isogai, A.; Wågberg, L.; Hu, L. Developing
Fibrillated Cellulose as a Sustainable Technological Material. Nature
2021, 590, 47−56.
(4) Altaner, C. M.; Thomas, L. H.; Fernandes, A. N.; Jarvis, M. C.
How Cellulose Stretches: Synergism between Covalent and Hydrogen
Bonding. Biomacromolecules 2014, 15, 791−798.
(5) Wei, Z.; Cai, C.; Huang, Y.; Wang, P.; Song, J.; Deng, L.; Fu, Y.
Strong Biodegradable Cellulose Materials with Improved Crystallinity
via Hydrogen Bonding Tailoring Strategy for UV Blocking and
Antioxidant Activity. Int. J. Biol. Macromol. 2020, 164, 27−36.
(6) Habibi, Y.; Lucia, L. A.; Rojas, O. J. Cellulose Nanocrystals:
Chemistry, Self-Assembly, and Applications. Chem. Rev. 2010, 110,
3479−3500.
(7) Rånby, B. G.; Banderet, A.; Sillén, L. G. Aqueous Colloidal
Solutions of Cellulose Micelles. Acta Chem. Scand. 1949, 3, 649−650.
(8) Revol, J.-F.; Godbout, L.; Dong, X.-M.; Gray, D. G.; Chanzy, H.;
Maret, G. Chiral Nematic Suspensions of Cellulose Crystallites; Phase
Separation and Magnetic Field Orientation. Liq. Cryst. 1994, 16,
127−134.
(9) Marchessault, R. H.; Morehead, F. F.; Walter, N. M. Liquid
Crystal Systems from Fibrillar Polysaccharides. Nature 1959, 184,
632−633.
(10) Revol, J.-F.; Bradford, H.; Giasson, J.; Marchessault, R. H.;
Gray, D. G. Helicoidal Self-Ordering of Cellulose Microfibrils in
Aqueous Suspension. Int. J. Biol. Macromol. 1992, 14, 170−172.
(11) Guidetti, G.; Atifi, S.; Vignolini, S.; Hamad, W. Y. Flexible
Photonic Cellulose Nanocrystal Films. Adv. Mater. 2016, 28, 10042−
10047.
(12) Walters, C. M.; Boott, C. E.; Nguyen, T. D.; Hamad, W. Y.;
MacLachlan, M. J. Iridescent Cellulose Nanocrystal Films Modified
with Hydroxypropyl Cellulose. Biomacromolecules 2020, 21, 1295−
1302.
(13) Shopsowitz, K. E.; Qi, H.; Hamad, W. Y.; MacLachlan, M. J.
Free-Standing Mesoporous Silica Films with Tunable Chiral Nematic
Structures. Nature 2010, 468, 422−425.
(14) He, Y.-D.; Zhang, Z.-L.; Xue, J.; Wang, X.-H.; Song, F.; Wang,
X.-L.; Zhu, L.-L.; Wang, Y.-Z. Biomimetic Optical Cellulose
Nanocrystal Films with Controllable Iridescent Color and Environmental Stimuli-Responsive Chromism. ACS Appl. Mater. Interfaces
2018, 10, 5805−5811.
(15) Tang, J.; Berry, R. M.; Tam, K. C. Stimuli-Responsive Cellulose
Nanocrystals for Surfactant-Free Oil Harvesting. Biomacromolecules
2016, 17, 1748−1756.
(16) Zhang, J.-H.; Xie, S.-M.; Zhang, M.; Zi, M.; He, P.-G.; Yuan, L.M. Novel Inorganic Mesoporous Material with Chiral Nematic
Structure Derived from Nanocrystalline Cellulose for High-Resolution
Gas Chromatographic Separations. Anal. Chem. 2014, 86, 9595−
9602.
(17) Walters, C. M.; Adair, K. R.; Hamad, W. Y.; MacLachlan, M. J.
Synthesis of Chiral Nematic Mesoporous Metal and Metal Oxide
Nanocomposites and their Use as Heterogeneous Catalysts. Eur. J.
Inorg. Chem. 2020, 2020, 3937−3943.
The Supporting Information is available free of charge at
https://pubs.acs.org/doi/10.1021/acs.chemmater.1c01272.
PXRD, Raman, TGA, nitrogen sorption isotherms, SEM,
mechanical testing, electrolyte uptake, and CV data of
the studied aerogels (PDF)
FIB-SEM milling video (MP4)
■
Article
AUTHOR INFORMATION
Corresponding Authors
Mark J. MacLachlan − Department of Chemistry, University
of British Columbia, Vancouver, British Columbia V6T 1Z1,
Canada; Stewart Blusson Quantum Matter Institute,
University of British Columbia, Vancouver, British Columbia
V6T 1Z4, Canada; WPI Nano Life Science Institute,
Kanazawa University, Kanazawa, Ishikawa 920-1192,
Japan; orcid.org/0000-0002-3546-7132;
Email: mmaclach@chem.ubc.ca
Erlantz Lizundia − Life Cycle Thinking Group, Department of
Graphic Design and Engineering Projects, Faculty of
Engineering in Bilbao, University of the Basque Country
(UPV/EHU), Bilbao 48013, Spain; BCMaterials, Basque
Center for Materials, Applications and Nanostructures,
48940 Leioa, Spain; Email: erlantz.liizundia@ehu.eus
Authors
Christopher M. Walters − Department of Chemistry,
University of British Columbia, Vancouver, British Columbia
V6T 1Z1, Canada
Gunwant K. Matharu − Department of Chemistry, University
of British Columbia, Vancouver, British Columbia V6T 1Z1,
Canada
Wadood Y. Hamad − Transformation and Interfaces Group,
Bioproducts Innovation Centre of Excellence, FPInnovations,
Vancouver, British Columbia V6T 1Z4, Canada;
orcid.org/0000-0003-1376-5865
Complete contact information is available at:
https://pubs.acs.org/10.1021/acs.chemmater.1c01272
Author Contributions
The manuscript was written through contributions of all
authors. All authors have given approval to the final version of
the manuscript.
Funding
M.J.M. thanks NSERC for funding (Discovery Grant,
CREATE Nanomat Grant).
Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS
The authors would like to thank The University of British
Columbia’s Centre for Higher Order Structure Elucidation (C5207
https://doi.org/10.1021/acs.chemmater.1c01272
Chem. Mater. 2021, 33, 5197−5209
Chemistry of Materials
pubs.acs.org/cm
(18) Boott, C. E.; Tran, A.; Hamad, W. Y.; MacLachlan, M. J.
Cellulose Nanocrystal Elastomers with Reversible Visible Color.
Angew. Chem., Int. Ed. 2020, 59, 226−231.
(19) Zhuo, H.; Hu, Y.; Chen, Z.; Peng, X.; Lai, H.; Liu, L.; Liu, Q.;
Liu, C.; Zhong, L. Linking Renewable Cellulose Nanocrystal into
Lightweight and Highly Elastic Carbon Aerogel. ACS Sustainable
Chem. Eng. 2020, 8, 11921−11929.
(20) Ke, F.-S.; Mishra, K.; Jamison, L.; Peng, X.-X.; Ma, S.-G.;
Huang, L.; Sun, S.-G.; Zhou, X.-D. Tailoring Nanostructures in
Micrometer Size Germanium Particles to Improve their Performance
as an Anode for Lithium Ion Batteries. Chem. Commun. 2014, 50,
3713−3715.
(21) Hayner, C. M.; Zhao, X.; Kung, H. H. Materials for
Rechargeable Lithium-Ion Batteries. Annu. Rev. Chem. Biomol. Eng.
2012, 3, 445−471.
(22) Hwang, J.; Jo, C.; Kim, M. G.; Chun, J.; Lim, E.; Kim, S.; Jeong,
S.; Kim, Y.; Lee, J. Mesoporous Ge/GeO2/Carbon Lithium-Ion
Battery Anodes with High Capacity and High Reversibility. ACS Nano
2015, 9, 5299−5309.
(23) Wang, X.; Liu, B.; Wang, Q.; Song, W.; Hou, X.; Chen, D.;
Cheng, Y.-B.; Shen, G. Three-Dimensional Hierarchical GeSe2
Nanostructures for High Performance Flexible All-Solid-State Supercapacitors. Adv. Mater. 2013, 25, 1479−1486.
(24) Yang, B.; Nie, A.; Chang, Y.; Cheng, Y.; Wen, F.; Xiang, J.; Li,
L.; Liu, Z.; Tian, Y. Metallic Layered Germanium Phosphide GeP5 for
High Rate Flexible All-Solid-State Supercapacitors. J. Mater. Chem. A
2018, 6, 19409−19416.
(25) Velmurugan, V.; Srinivasarao, U.; Ramachandran, R.; Saranya,
M.; Grace, A. N. Synthesis of Tin Oxide/Graphene (SnO2/G)
Nanocomposite and its Electrochemical Properties for Supercapacitor
Applications. Mater. Res. Bull. 2016, 84, 145−151.
(26) Lim, S. P.; Huang, N. M.; Lim, H. N. Solvothermal Synthesis of
SnO2/Graphene Nanocomposites for Supercapacitor Application.
Ceram. Int. 2013, 39, 6647−6655.
(27) Choudhury, A.; Dalal, A.; Dhar Dwivedi, S. M. M.; Ghosh, A.;
Halder, N.; Das, S.; Mondal, A. Vapour Transport Grown
Photosensitive GeO2 Thin Film. Mater. Res. Bull. 2021, 142, 111397.
(28) Liu, P.; Ru, Q.; Zheng, P.; Shi, Z.; Liu, Y.; Su, C.; Hou, X.; Su,
S.; Chi-Chung Ling, F. One-Step Synthesis of Zn2GeO4/CNT-O
Hybrid with Superior Cycle Stability for Supercapacitor Electrodes.
Chem. Eng. J. 2019, 374, 29−38.
(29) Augustyn, V.; Simon, P.; Dunn, B. Pseudocapacitive Oxide
Materials for High-Rate Electrochemical Energy Storage. Energy
Environ. Sci. 2014, 7, 1597−1614.
(30) Saha, D.; Kruse, P. Conductive Forms of MoS2 and Their
Applications in Energy Storage and Conversion. J. Electrochem. Soc.
2020, 167, 126517.
(31) Wang, R.; Yao, M.; Niu, Z. Smart Supercapacitors from
Materials to Devices. InfoMat 2020, 2, 113−125.
(32) Zhu, M.; Huang, Y.; Huang, Y.; Pei, Z.; Xue, Q.; Li, H.; Geng,
H.; Zhi, C. Capacitance Enhancement in a Semiconductor
Nanostructure-Based Supercapacitor by Solar Light and a SelfPowered Supercapacitor−Photodetector System. Adv. Funct. Mater.
2016, 26, 4481−4490.
(33) Bushick, K.; Mengle, K. A.; Chae, S.; Kioupakis, E. Electron and
Hole Mobility of Rutile GeO2 from First Principles: An UltrawideBandgap Semiconductor for Power Electronics. Appl. Phys. Lett. 2020,
117, 182104.
(34) Peng, H.; Yan, B.; Jiang, M.; Liu, B.; Gu, Y.; Yao, G.; Zhang, Y.;
Ye, L.; Bai, X.; Chen, S. A Coral-Like Polyaniline/Barium Titanate
Nanocomposite Electrode with Double Electric Polarization for
Electrochromic Energy Storage Applications. J. Mater. Chem. A 2021,
9, 1669−1677.
(35) Liu, Y. X., Liang; Wang, M. Barium Titanate-Doped Super
Capacitor Electrode Material Having High Dielectric Constant and
Preparation Method Thereof. Chinese Patent CN106024406A, 2016.
(36) Yang, X.; Shi, K.; Zhitomirsky, I.; Cranston, E. D. Cellulose
Nanocrystal Aerogels as Universal 3D Lightweight Substrates for
Supercapacitor Materials. Adv. Mater. 2015, 27, 6104−6109.
Article
(37) Chen, S.; Xing, W.; Duan, J.; Hu, X.; Qiao, S. Z.
Nanostructured Morphology Control for Efficient Supercapacitor
Electrodes. J. Mater. Chem. A 2013, 1, 2941−2954.
(38) Du, A.; Zhou, B.; Zhang, Z.; Shen, J. A Special Material or a
New State of Matter: A Review and Reconsideration of the Aerogel.
Materials 2013, 6, 941−968.
(39) Barrios, E.; Fox, D.; Li Sip, Y. Y.; Catarata, R.; Calderon, J. E.;
Azim, N.; Afrin, S.; Zhang, Z.; Zhai, L. Nanomaterials in Advanced,
High-Performance Aerogel Composites: A Review. Polymers 2019, 11,
726.
(40) Chen, G.; Chen, B.; Liu, T.; Mei, Y.; Ren, H.; Bi, Y.; Luo, X.;
Zhang, L. The Synthesis and Characterization of Germanium Oxide
Aerogel. J. Non-Cryst. Solids 2012, 358, 3322−3326.
(41) Fang, S.; Shen, L.; Zheng, H.; Zhang, X. Ge−Graphene−
Carbon Nanotube Composite Anode for High Performance LithiumIon Batteries. J. Mater. Chem. A 2015, 3, 1498−1503.
(42) Li, D.; Wang, H.; Liu, H. K.; Guo, Z. A New Strategy for
Achieving a High Performance Anode for Lithium Ion Batteries
Encapsulating Germanium Nanoparticles in Carbon Nanoboxes. Adv.
Energy Mater. 2016, 6, 1501666.
(43) Meng, X.; Al-Salman, R.; Zhao, J.; Borissenko, N.; Li, Y.;
Endres, F. Electrodeposition of 3D Ordered Macroporous Germanium from Ionic Liquids: A Feasible Method to Make Photonic
Crystals with a High Dielectric Constant. Angew. Chem., Int. Ed. 2009,
48, 2703−2707.
(44) Armatas, G. S.; Kanatzidis, M. G. Mesostructured Germanium
with Cubic Pore Symmetry. Nature 2006, 441, 1122−1125.
(45) Xu, Y.-T.; Dai, Y.; Nguyen, T.-D.; Hamad, W. Y.; MacLachlan,
M. J. Aerogel Materials with Periodic Structures Imprinted with
Cellulose Nanocrystals. Nanoscale 2018, 10, 3805−3812.
(46) Nguyen, T.-D.; Lizundia, E.; Niederberger, M.; Hamad, W. Y.;
MacLachlan, M. J. Self-Assembly Route to TiO2 and TiC with a
Liquid Crystalline Order. Chem. Mater. 2019, 31, 2174−2181.
(47) Puech, P.; Kandara, M.; Paredes, G.; Moulin, L.; Weiss-Hortala,
E.; Kundu, A.; Ratel-Ramond, N.; Plewa, J.-M.; Pellenq, R.;
Monthioux, M. Analyzing the Raman Spectra of Graphenic Analyzing
the Raman Spectra of Graphenic Carbon Materials from Kerogens to
Nanotubes: What Type of Information Can Be Extracted from Defect
Bands? C 2019, 5, 69.
(48) Lizundia, E.; Nguyen, T.-D.; Vilas, J. L.; Hamad, W. Y.;
MacLachlan, M. J. Chiroptical, Morphological and Conducting
Properties of Chiral Nematic Mesoporous Cellulose/Polypyrrole
Composite Films. J. Mater. Chem. A 2017, 5, 19184−19194.
(49) Casas, X.; Niederberger, M.; Lizundia, E. A Sodium-Ion Battery
Separator with Reversible Voltage Response Based on Water-Soluble
Cellulose Derivatives. ACS Appl. Mater. Interfaces 2020, 12, 29264−
29274.
(50) Adachi, T.; Sakka, S. The Role of N,N-dimethylformamide, a
DCCA, in the Formation of Silica Gel Monoliths by Sol-Gel Method.
J. Non-Cryst. Solids 1988, 99, 118−128.
(51) Zhou, L.; Zhai, Y.-M.; Yang, M.-B.; Yang, W. Flexible and
Tough Cellulose Nanocrystal/Polycaprolactone Hybrid Aerogel
Based on the Strategy of Macromolecule Cross-Linking via Click
Chemistry. ACS Sustainable Chem. Eng. 2019, 7, 15617−15627.
(52) McMurdie, H. F.; Morris, M. C.; Evans, E. H.; Paretzkin, B.;
Wong-Ng, W.; Ettlinger, L.; Hubbard, C. R. Standard X-Ray
Diffraction Powder Patterns from the JCPDS Research Associateship.
Powder Diffr. 1986, 1, 64−77.
(53) Patterson, A. L. The Scherrer Formula for X-Ray Particle Size
Determination. Phys. Rev. 1939, 56, 978−982.
(54) Kim, H.; Viswanathamurthi, P.; Bhattarai, N.; Lee, D.
Preparation and Morphology of Germanium Oxide Nanofibers. Rev.
Adv. Mater. Sci. 2003, 5, 220−223.
(55) Eom, Y.; Son, S. M.; Kim, Y. E.; Lee, J.-E.; Hwang, S.-H.; Chae,
H. G. Structure Evolution Mechanism of Highly Ordered Graphite
During Carbonization of Cellulose Nanocrystals. Carbon 2019, 150,
142−152.
(56) Zhu, H.; Shen, F.; Luo, W.; Zhu, S.; Zhao, M.; Natarajan, B.;
Dai, J.; Zhou, L.; Ji, X.; Yassar, R. S.; Li, T.; Hu, L. Low Temperature
5208
https://doi.org/10.1021/acs.chemmater.1c01272
Chem. Mater. 2021, 33, 5197−5209
Chemistry of Materials
pubs.acs.org/cm
Carbonization of Cellulose Nanocrystals for High Performance
Carbon Anode of Sodium-Ion Batteries. Nano Energy 2017, 33,
37−44.
(57) Sing, K. S. W.; Everett, D. H.; Haul, R. A. W.; Moscou, L.;
Pierotti, R. A.; Rouquérol, J.; Simieniewska, T. Reporting
Physisorption Data for Gas/Solid Systems with Special Reference to
the Determination of Surface Area and Porosity. Pure Appl. Chem.
1985, 57, 603−619.
(58) Zaman, A.; Huang, F.; Jiang, M.; Wei, W.; Zhou, Z.
Preparation, Properties, and Applications of Natural Cellulosic
Aerogels: A Review. Energy Built Environ. 2020, 1, 60−76.
(59) Roman, M.; Winter, W. T. Effect of Sulfate Groups from
Sulfuric Acid Hydrolysis on the Thermal Degradation Behavior of
Bacterial Cellulose. Biomacromolecules 2004, 5, 1671−1677.
(60) Wicklein, B.; Kocjan, A.; Salazar-Alvarez, G.; Carosio, F.;
Camino, G.; Antonietti, M.; Bergström, L. Thermally Insulating and
Fire-Retardant Lightweight Anisotropic Foams Based on Nanocellulose and Graphene Oxide. Nat. Nanotechnol. 2015, 10, 277−283.
(61) Abraham, E.; Weber, D. E.; Sharon, S.; Lapidot, S.; Shoseyov,
O. Multifunctional Cellulosic Scaffolds from Modified Cellulose
Nanocrystals. ACS Appl. Mater. Interfaces 2017, 9, 2010−2015.
(62) Tripathi, A.; Tardy, B. L.; Khan, S. A.; Liebner, F.; Rojas, O. J.
Expanding the Upper Limits of Robustness of Cellulose Nanocrystal
Aerogels: Outstanding Mechanical Performance and Associated Pore
Compression Response of Chiral-Nematic Architectures. J. Mater.
Chem. A 2019, 7, 15309−15319.
(63) Fabritius, H.-O.; Sachs, C.; Triguero, P. R.; Raabe, D. Influence
of Structural Principles on the Mechanics of a Biological Fiber-Based
Composite Material with Hierarchical Organization: The Exoskeleton
of the Lobster Homarus americanus. Adv. Mater. 2009, 21, 391−400.
(64) Wan, J.; Zhang, J.; Yu, J.; Zhang, J. Cellulose Aerogel
Membranes with a Tunable Nanoporous Network as a Matrix of Gel
Polymer Electrolytes for Safer Lithium-Ion Batteries. ACS Appl.
Mater. Interfaces 2017, 9, 24591−24599.
(65) Ding, C.; Fu, X.; Li, H.; Yang, J.; Lan, J.-L.; Yu, Y.; Zhong, W.H.; Yang, X. An Ultrarobust Composite Gel Electrolyte Stabilizing Ion
Deposition for Long-Life Lithium Metal Batteries. Adv. Funct. Mater.
2019, 29, 1904547.
(66) Li, L.; Lu, F.; Wang, C.; Zhang, F.; Liang, W.; Kuga, S.; Dong,
Z.; Zhao, Y.; Huang, Y.; Wu, M. Flexible Double-Cross-Linked
Cellulose-Based Hydrogel and Aerogel Membrane for Supercapacitor
Separator. J. Mater. Chem. A 2018, 6, 24468−24478.
(67) Haller, P. D.; Bradley, L. C.; Gupta, M. Effect of Surface
Tension, Viscosity, and Process Conditions on Polymer Morphology
Deposited at the Liquid−Vapor Interface. Langmuir 2013, 29,
11640−11645.
(68) Zhang, J.; Yue, L.; Kong, Q.; Liu, Z.; Zhou, X.; Zhang, C.; Xu,
Q.; Zhang, B.; Ding, G.; Qin, B.; Duan, Y.; Wang, Q.; Yao, J.; Cui, G.;
Chen, L. Sustainable, Heat-Resistant and Flame-Retardant CelluloseBased Composite Separator for High-Performance Lithium Ion
Battery. Sci. Rep. 2015, 4, 3935.
(69) Heath, L.; Thielemans, W. Cellulose Nanowhisker Aerogels.
Green Chem. 2010, 12, 1448−1453.
(70) Cho, M.; Karaaslan, M. A.; Renneckar, S.; Ko, F. Enhancement
of the Mechanical Properties of Electrospun Lignin-Based Nanofibers
by Heat Treatment. J. Mater. Sci. 2017, 52, 9602−9614.
(71) Qin, Y.; Peng, Q.; Ding, Y.; Lin, Z.; Wang, C.; Li, Y.; Xu, F.; Li,
J.; Yuan, Y.; He, X.; Li, Y. Lightweight, Superelastic, and Mechanically
Flexible Graphene/Polyimide Nanocomposite Foam for Strain Sensor
Application. ACS Nano 2015, 9, 8933−8941.
(72) Parmenter, K. E.; Milstein, F. Mechanical Properties of Silica
Aerogels. J. Non-Cryst. Solids 1998, 223, 179−189.
(73) Hänsel, C.; Lizundia, E.; Kundu, D. A Single Li-Ion Conductor
Based on Cellulose. ACS Appl. Energy Mater. 2019, 2, 5686−5691.
(74) Zhang, Y.; Shang, Z.; Shen, M.; Chowdhury, S. P.; Ignaszak, A.;
Sun, S.; Ni, Y. Cellulose Nanofibers/Reduced Graphene Oxide/
Polypyrrole Aerogel Electrodes for High-Capacitance Flexible AllSolid-State Supercapacitors. ACS Sustainable Chem. Eng. 2019, 7,
11175−11185.
Article
(75) Ryu, J.; Hong, D.; Shin, S.; Choi, W.; Kim, A.; Park, S.
Hybridizing Germanium Anodes with Polysaccharide-Derived Nitrogen-Doped Carbon for High Volumetric Capacity of Li-Ion Batteries.
J. Mater. Chem. A 2017, 5, 15828−15837.
(76) Yu, Z.; Tetard, L.; Zhai, L.; Thomas, J. Supercapacitor
Electrode Materials: Nanostructures From 0 to 3 Dimensions. Energy
Environ. Sci. 2015, 8, 702−730.
(77) Huang, Y.; Zhong, M.; Huang, Y.; Zhu, M.; Pei, Z.; Wang, Z.;
Xue, Q.; Xie, X.; Zhi, C. A Self-Healable and Highly Stretchable
Supercapacitor Based on a Dual Crosslinked Polyelectrolyte. Nat.
Commun. 2015, 6, 10310.
(78) Zhuo, H.; Hu, Y.; Tong, X.; Zhong, L.; Peng, X.; Sun, R.
Sustainable Hierarchical Porous Carbon Aerogel from Cellulose for
High-Performance Supercapacitor and CO2 Capture. Ind. Crops Prod.
2016, 87, 229−235.
(79) Shi, K.; Yang, X.; Cranston, E. D.; Zhitomirsky, I. Efficient
Lightweight Supercapacitor with Compression Stability. Adv. Funct.
Mater. 2016, 26, 6437−6445.
(80) Wang, Z.; Tammela, P.; Huo, J.; Zhang, P.; Strømme, M.;
Nyholm, L. Solution-Processed Poly(3,4-ethylenedioxythiophene)
Nanocomposite Paper Electrodes for High-Capacitance Flexible
Supercapacitors. J. Mater. Chem. A 2016, 4, 1714−1722.
(81) Zheng, W.; Lv, R.; Na, B.; Liu, H.; Jin, T.; Yuan, D.
Nanocellulose-Mediated Hybrid Polyaniline Electrodes for High
Performance Flexible Supercapacitors. J. Mater. Chem. A 2017, 5,
12969−12976.
(82) Liu, Z.; Yuan, X.; Zhang, S.; Wang, J.; Huang, Q.; Yu, N.; Zhu,
Y.; Fu, L.; Wang, F.; Chen, Y.; Wu, Y. Three-Dimensional Ordered
Porous Electrode Materials for Electrochemical Energy Storage. NPG
Asia Mater. 2019, 11, 12.
(83) Zeng, Y.-F.; Xin, G.-X.; Bulin, C.-K.; Zhang, B.-W. One-Step
Preparation and Electrochemical Performance of 3D Reduced
Graphene Oxide/NiO as Supercapacitor Electrodes Materials. Wuji
Cailiao Xuebao 2018, 33, 1070−1076.
(84) Liu, T.; Finn, L.; Yu, M.; Wang, H.; Zhai, T.; Lu, X.; Tong, Y.;
Li, Y. Polyaniline and Polypyrrole Pseudocapacitor Electrodes with
Excellent Cycling Stability. Nano Lett. 2014, 14, 2522−2527.
(85) Ye, T.; Zou, Y.; Xu, W.; Zhan, T.; Sun, J.; Xia, Y.; Zhang, X.;
Yang, D. Poorly-Crystallized Poly(Vinyl Alcohol)/Carrageenan
Matrix: Highly Ionic Conductive and Flame-Retardant Gel Polymer
Electrolytes for Safe and Flexible Solid-State Supercapacitors. J. Power
Sources 2020, 475, 228688.
(86) Yan, T.; Zou, Y.; Zhang, X.; Li, D.; Guo, X.; Yang, D. Hydrogen
Bond Interpenetrated Agarose/PVA Network: A Highly Ionic
Conductive and Flame-Retardant Gel Polymer Electrolyte. ACS
Appl. Mater. Interfaces 2021, 13, 9856−9864.
(87) Yang, Y. A Mini-Review: Emerging All-Solid-State Energy
Storage Electrode Materials for Flexible Devices. Nanoscale 2020, 12,
3560−3573.
(88) Zhou, Y.; Wan, C.; Yang, Y.; Yang, H.; Wang, S.; Dai, Z.; Ji, K.;
Jiang, H.; Chen, X.; Long, Y. Highly Stretchable, Elastic, and Ionic
Conductive Hydrogel for Artificial Soft Electronics. Adv. Funct. Mater.
2019, 29, 1806220.
5209
https://doi.org/10.1021/acs.chemmater.1c01272
Chem. Mater. 2021, 33, 5197−5209