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Cathode and anode materials for sodium ion batteries
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Content
CATHODE AND ANODE MATERIALS FOR SODIUM ION BATTERIES
by
Yihang Liu
Dissertation submitted to the
Faculty of the Graduate School
University of Southern California
In Partial Fulfillment of the
Requirements for the Degree
[Doctor of Philosophy]
(Electrical Engineering)
[2018]
Advisory Committee:
Professor Chongwu Zhou, Chair
Professor Wei Wu
Professor Aiichiro Nakano
© Copyright by
[Yihang Liu]
[2018]
ii
Dedication
To my parents:
Thank you for your unconditional love and support. Thank you for your
unwavering faith and absolute belief in my abilities.
iii
Acknowledgements
First and foremost, I would like to express my deep and sincere gratitude to
my supervisor, Dr. Chongwu Zhou. I am really honored to work under his guidance. I
would like to thank Dr. Wei Wu, Dr. Aiichiro Nakano, Dr. Stephen B. Cronin, and
Dr. Han Wang for serving on my committee and taking time attending my Ph. D.
qualifying exam and defense.
I would like to acknowledge all the other members in Dr. Chongwu Zhou’s
group members, with whom I have worked, and because of whom my graduate
experience has been one that I will cherish forever
iv
Table of Contents
Dedication ..................................................................................................................... ii
Acknowledgements ...................................................................................................... iii
Table of Contents ......................................................................................................... iv
List of Figures ............................................................................................................... v
Chapter 1: Introduction of lithium-ion and sodium-ion batteries ................................. 1
1.1 Background of lithium-ion batteries ................................................................... 1
1.2 Working principles of lithium-ion battery system .............................................. 2
1.3 Introduction of sodium-ion batteries ................................................................... 5
Chapter 2: SnO2 Coated Carbon Cloth with Surface Modification as Na-ion Battery
Anode ............................................................................................................................ 7
2.1 Introduction ......................................................................................................... 7
2.2 Experimental ....................................................................................................... 8
2.3 Results and Discussion ..................................................................................... 10
2.4 Summary ........................................................................................................... 18
Chapter 3: Layered P2-Na2/3[Ni1/3Mn2/3]O2 as High-Voltage Cathode for Sodium-Ion
Batteries: The Capacity Decay Mechanism and Al2O3 Surface Modification ........... 20
3.1 Introduction ....................................................................................................... 20
3.2 Experimental ..................................................................................................... 22
3.3 Results and Discussion ..................................................................................... 24
3.4 Summary ........................................................................................................... 34
Chapter 4: Red Phosphorus Nano-Dots on Reduced Graphene Oxide as a Flexible and
Ultra-Fast Anode for Sodium-Ion Batteries ................................................................ 35
4.1 Introduction ....................................................................................................... 35
4.2 Experimental ..................................................................................................... 37
4.3 Results and Discussion ..................................................................................... 39
4.4 Summary ........................................................................................................... 47
Chapter 5: Single-Step Flash-Heat-Treatment Synthesized Flame-Retardant Red
Phosphorus/Graphene Composite as Flexible Anode for Sodium-Ion Batteries ........ 48
5.1 Introduction ....................................................................................................... 48
5.2 Experimental ..................................................................................................... 50
5.3 Results and Discussion ..................................................................................... 53
5.4 Summary ........................................................................................................... 62
Chapter 6: Room-Temperature Pressure Synthesis of Layered Black Phosphorus-
Graphene Composite for Sodium-Ion Battery Anodes ............................................... 63
6.1 Introduction ....................................................................................................... 63
6.2 Experimental ..................................................................................................... 66
6.3 Results and Discussion ..................................................................................... 67
6.4 Summary ........................................................................................................... 76
Bibliography ............................................................................................................... 77
v
List of Figures
Figure 1.1 Schematic structure of a typical lithium-ion battery and its discharge process.
Figure 2.1 The two main steps in the preparation of the virus enabled 3D current collector:
TMV1cys self-assembling and nickel chemical deposition.
Figure 2.2 (a) Optical image of flexible SnO2/CC binder free electrode, (b) XRD profile
of hydrothermal synthesized SnO2, (c) SEM image of SnO2/CC and (d) cross section
SEM image of a single SnO2/CC fiber and its EDX element mapping profile of carbon
(e) and tin (f).
Figure 2.3 (a) TEM image of a large bulk piece of SnO2 layer coated on CC, (b, c)
High resolution TEM image of SnO2 nanocrystal at different magnification, (d) FFT
image of the area marked in (b). TEM images of the carbon layer (e) and Al2O3 layer
(f) coated on CC.
Figure 2.4 (a) Cycling stability and Coulombic efficiency of SnO2/CC, C/SnO2/CC and
Al2O3/SnO2/CC. Potential profiles of (b) C/SnO2/CC and (c) Al2O3/SnO2/CC in cycling
test. (d) Rate performance of C/SnO2/CC and Al2O3/SnO2/CC.
Figure 2.5 The electrochemical impedance spectrum (EIS) profiles of (a) C/SnO2/CC
and (b) Al2O3/SnO2/CC fresh cell and cell after 100 cycles. The SEM images of (c)
C/SnO2/CC and (d) Al2O3/SnO2/CC after 100 cycles.
Figure 3.1 Schematic figure of the exfoliation during the sodiation and de-sodiation
process of the P2-Na2/3[Ni1/3Mn2/3]O2 particle.
Figure 3.2 (a) SEM image of as-prepared Na2/3[Ni1/3Mn2/3]O2. SEM image of a single
Na2/3[Ni1/3Mn2/3]O2 particle (b), and the DEX mapping profile of elements Na (c), Ni
(d) and Mn (e). (f) The TEM image of as-prepared Na2/3[Ni1/3Mn2/3]O2 and (g) the
vi
SAED pattern of the area mark in red circle in (f). The high-resolution TEM image of
(h) as-prepared Na2/3[Ni1/3Mn2/3]O2 particle with Na ion channel marked and (i) Al2O3
coated Na2/3[Ni1/3Mn2/3]O2.
Figure 3.3 (a) Cycling performance of the as-prepared Na2/3[Ni1/3Mn2/3]O2 and Al2O3-
Na2/3[Ni1/3Mn2/3]O2. The charge and discharge profiles of (b) Na2/3[Ni1/3Mn2/3]O2
electrodes and (c) Al2O3-Na2/3[Ni1/3Mn2/3]O2 electrode. (d) Rate performance of the as-
prepared Na2/3[Ni1/3Mn2/3]O2 and Al2O3-Na2/3[Ni1/3Mn2/3]O2. Cyclic voltammetry
profiles of (e) as-prepared Na2/3[Ni1/3Mn2/3]O2 cathode and (f)Al2O3-
Na2/3[Ni1/3Mn2/3]O2 cathode.
Figure 3.4 (a) SEM image of after-cycling Na2/3[Ni1/3Mn2/3]O2 electrode with the
exfoliation and carbon black areas marked. (b) STEM of an after-cycling
Na2/3[Ni1/3Mn2/3]O2 particle. (c) TEM image of an after-cycling Na2/3[Ni1/3Mn2/3]O2
particle. (d) Enlarged TEM image of the area marked in (c). (e) High-resolution TEM
image of the exfoliation opening in (d). (f) The SAED pattern of the area marked in (c).
Figure 3.5 The XRD pattern of the as-prepared Na2/3[Ni1/3Mn2/3]O2, after-cycling
Na2/3[Ni1/3Mn2/3]O2 and after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2. (a) is the enlarged
view of (b) from 15
o
to 17
o
.
Figure 3.6 (a) The SEM image of the after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2
electrode. (b) The high-resolution TEM image of an after-cycling Al2O3-
Na2/3[Ni1/3Mn2/3]O2 particle. (c) The SAED pattern of the after-cycling Al2O3-
Na2/3[Ni1/3Mn2/3]O2 particle.
Figure 4.1 Schematic description of P@RGO synthesis.
vii
Figure 4.2 (a) SEM image of the P@RGO composite. (b) The enlarged SEM image
and (c) the corresponding EDS mapping profile for phosphorus element in the area
marked with red rectangle in (a). (d) TEM and (e) STEM image of the P@RGO
composite. (f) EDS line-scan profile for phosphorus element of three particles marked
in (e). (g) TEM image of a single phosphorus particle on a RGO sheet. (h) The high-
resolution TEM image of the area marked with blue dashed rectangle in (g), the
graphene layer is marked with light blue dashed line. (i) The bending test of the
P@RGO flexible film with the resistance measurement at different bending radius, the
optical images of the flexible film are inserted.
Figure 4.3 (a) Thermogravimetric analysis of the P@RGO composite. (b) Cyclic
voltammetry of the P@RGO anode with a scan rate of 0.1 mV/s between 0 and 3.0 V
vs. Na/Na
+
. (c) Cycling performance of the P@RGO anode at a charge/discharge
current density of 1593.9 mA/g with its potential profiles presented in (d). (e) Rate
performance of the P@RGO anode with its potential profiles presented in (f).
Figure 4.4 Cross-sectional (a) and planar (d) SEM images of the flexible P@RGO
electrode after 300 cycles with the corresponding EDS element mapping profile of
phosphorus in (b) and (e) and carbon in (c) and (f), respectively.
Figure 5.1 Schematic diagrams of the RP/rGO composite synthesis process at (a) pre-
heat-treatment, (b) flash-heat-treatment, and (c) cooling-down steps, respectively. (d)
Schematic diagram of the nano-structure of the resulted RP/rGO composite.
Figure 5.2 (a, b) SEM images of the RP/rGO composite at × 1500 and × 4000
magnifications. (c, d) TEM and STEM images of a single piece of RP/rGO flake. (e)
Enlarged STEM image of the RP/rGO composite with the corresponding EDS mapping
viii
profile of the P element in (f). The RP filled into the void space between the rGO layers
was marked by blue lines in (e).
Figure 5.3 (a) Raman spectra of the RP/rGO composite and commercial RP powder.
(b) XPS spectrum of the synthesized RP/rGO composite and pristine GO powder. (c)
The high-resolution C 1s XPS spectrum of the synthesized RP/rGO composite and
pristine GO powder. (d) Optical images of a flexible RP/rGO film was wrapped on a
glass rod with a diameter of 5 mm, and a punched electrode was bent with a pair of
tweezers.
Figure 5.4 (a) Cycling performance of the fabricated RP/rGO flexible film anode, the
commercial RP mixed with rGO, and the commercial RP mixed with carbon black (CB)
at a charge/discharge current density of 1 A/g with its potential profiles exhibited in
(b). Rate performance of the RP/rGO anode with its potential profiles exhibited in (d).
Figure 5.5 (a, b) TEM and STEM images of the flexible RP/rGO anode after 200 cycles,
with the corresponding EDS spectrum (c) and elemental mapping profiles of
phosphorus in (d), sodium in (e), and carbon in (f), respectively.
Figure 5.6 Combustion tests of an RP/rGO flexible film (a) and an rGO film (b) at t=0,
0.5, 2, 5 and 10 s, respectively. t=0 s is defined as the moment when the flame on the
films can be clearly visualized.
Figure 6.1 Schematic description of the BP/rGO synthesis.
Figure 6.2 (a) Planar and (c) cross-section SEM images of the precursor RP/rGO film;
the same fields of view are shown EDS elemental maps of phosphorus concentration
in (b) and (d). (e) Planar and (f) cross-sectional SEM images of the pressure-
synthesized BP/rGO film. (g) XRD patterns of the GO, RP/rGO precursor, as-prepared
ix
and post-cycling BP/rGO samples. (h) The Raman spectra of RP/rGO, as-prepared and
post-cycling BP/rGO samples.
Figure 6.3 (a) TEM and (b) STEM image of the BP/rGO composite, with the
corresponding maps of X-ray intensity by EDS for (c) phosphorus and (d) carbon. (e)
Enlarged TEM image and (f) high-resolution TEM image of the area marked with light
blue rectangle in (e) with the corresponding FFT images shown in the insets.
Figure 6.4 (a) Cycling performance of the BP/rGO anodes at charge and discharge
current densities of 1 and 40 A/g in black and red color, respectively; specific capacity
is plotted as solid curves against the left-hand axis, whereas Coulombic efficiency is
plotted as open circles against the right-hand axis. (b) Charging and discharging
potential profiles at 1 A/g current density (black curve in (a)), shown for selected cycle
numbers. (c) Rate performance of the BP/rGO anode for a series of tests with five
cycles at each value of current density. (d) Charging and discharging potential profiles
for the rate tests presented in (c), color-coded by current density. (e) Electrochemical
impedance spectra for RP/rGO and BP/rGO anodes.
Figure 6.5 TEM results for the post-cycling BP/rGO anode: (a) TEM image with FFT
inset; (b) STEM image; (c) EDS map of phosphorus X-ray intensity; (d) EDS map of
carbon X-ray intensity.
1
Chapter 1: Introduction of lithium-ion and sodium-ion batteries
1.1 Background of lithium-ion batteries
Due to the increasing global demand for fossil fuels and the direct or indirect
environmental consequences of their use, great attention is being devoted to alternative
technologies for both energy generation and storage. Secondary battery systems have
been considered as one of the energy storage strategy for the energy generated by
renewable such as solar, wind and geothermal energy. Among the different types of
secondary batteries available in the marketplace, lithium-ion batteries became the
predominant battery technology for portable electronics, medical devices and electric
vehicles during the past two decades.[1]
Lithium-ion batteries have around three to four times the energy of standard
lead-acid or nickel metal hydride batteries for the same size and weight. Lithium-ion
batteries can operate over a wide temperature range, typically from -20
o
C to +50
o
C,
and can last for hundreds or even thousands of charge/discharge cycles with near 100%
energy efficiency with tiny memory effect and low self-discharge. The high operation
voltage (~ 3.5 V) of lithium-ion batteries allows a single-cell-pack design for most
portable devices.[1] All the advantages mentioned above make lithium-ion batteries
acquire great popularity for portable devices, implantable devices and electric vehicles.
2
1.2 Working principles of lithium-ion battery system
Figure 1.1 Schematic illustration of a typical lithium-ion battery during its discharge
process.[2]
The three primary functional components of a lithium-ion battery are negative
electrode, positive electrode and electrolyte. Insertion materials have been intensively
used as electrode materials for many rechargeable lithium-ion battery systems.
Electrochemical insertion reaction can be referred to a redox reaction, where the charge
transfer occurs at the electrode/electrolyte interface, during the intercalation of lithium
ions into the solid electrode material. The liquid electrolyte allows the diffusion of ions,
but is electronically insulating. The electrons are transferred through an external circuit
to provide powder to the load. In order to keep the system electrical neutral, the positive
electrode material compensates charge for the removal of lithium ion by oxidizing the
transition metal present in the crystal lattice. Upon reaching the negative electrode, the
lithium ions insert into the electrode material and recombine with the electrons from
3
external circuit, resulting in the electrochemical reduction of the negative electrode
material.[3]
During discharge in a lithium-ion cell, the process is reversed as that illustrated
in Figure 1.1. Lithium ions were extracted from the negative electrode material through
the liquid electrolyte, and back to the positive electrode material; while electrons shuttle
from the negative electrode material to the positive electrode material through the
external circuit. The electrode material compensates charge via reducing the transition
metal ion to its original oxidation state. In a typical LiCoO2/graphite cell, the
electrochemical reactions can be presented as the following:
Positive electrode (charge): LiCoO2 → Li1-xCoO2 + xLi
+
+ xe
-
Negative electrode (charge): xC6 + xLi
+
+ xe
-
→ xLiC6
Positive electrode (discharge): Li1-xCoO2 + xLi
+
+ xe
-
→ LiCoO2
Negative electrode (discharge): xLiC6 → xC6 + xLi
+
+ xe
-
Since lithium ions are transferred back and forth between the positive and
negative electrode materials, as ions flow through the electrolyte, and the electrons flow
through the external circuit, powering the portable device and electric vehicle. Thus,
the electrode materials in a battery system must allow for the conductivity of both
lithium ions and electrons. The rate capability of lithium-ion batteries mainly depends
on the particle size of the host material during insertion and extraction of lithium ions
and the kinetics of the lithium ion diffusion and electron transport in the host material.
Mechanical stresses and strains occur during lithiation and de-lithiation processes,
causing cracks in the material particles, which can eventually lead to the material
peeling off from the current collector and cause a collapse of the battery system, and
4
thus limiting the cycle life of the cell. Nevertheless, many electrode materials were
found to be stable during lithium insertion and removal processes with small strains
and stresses. These low strain materials have volume changes around 10% between the
fully lithiated and de-lithiated states.
Graphite has been used as the anode material for current commercial lithium-
ion batteries since the launch of the first commercial lithium-ion battery. During the
past two decades, a lot of effort has been devoted toward identifying alternative anode
materials that have higher lithium-ion capacity, higher charge/discharge rate and better
electrochemical stability. Insertion alloys (Si, Sn, Ge), redox metal oxides, and carbon
allotropes have been explored as anode materials for the next-generation lithium-ion
batteries.[1]
Unlike the anode, for which high-storage capacity materials are known to exist,
the comparatively low storage capacity of most known cathode materials has been
recognized as a major limiting factor in the overall performance of lithium-ion batteries.
Since the successful introduction of the LiCoO2 cathode in 1991,[4] other positive
electrodes mainly fall into two categories.[5] The first group is layered compounds with
a close packed oxygen negative ions lattice, in which transition metal positive ions
present in layers between the negative ions, and lithium ions are inserted into the
available layers. The materials belong to this group gain the advantages of higher
operating voltage and specific energy density than the second group, due to their highly
oxidizing redox-active couples and compact lattices. These materials are compositional
variations of layered LiCoO2, such as LiNiO2, LiMnO2 and LiNi1−xCoxO2, as well as
spinel structures derived from LiMn2O4.[6] The second group consists of metal oxides
5
such as V2O5 and MnO2, and transition metal phosphates such as the olivine type
LiFePO4. Although their operating voltage is slightly lower than the materials in the
first group, the less cost, enhanced safety and better kinetics of these compounds make
them competitive cathode candidates.[7]
1.3 Introduction of sodium-ion batteries
Since its commercialization two decades ago, Li-ion batteries have essentially
dominated the portable electronics applications in small formats while being poised to
enter more profitable and strategically important markets of automotive and grid
storage. However, the limited abundance of Li in earth-crust, its uneven geographic
distribution and difficulties in recycling Li resources have raised concerns about large
scale application of this chemistry.[8] As an alternative to Li-ion chemistry, Na-ion
batteries have attracted increasing attention, because of the low cost associated with its
high natural occurrences in both earth and ocean, and decent energy densities blessed
by its similar chemical natures to Li. Given this similarity, many mature electrode
materials for Li-ion chemistry have been investigated as drop-in replacement for Na-
ion; however, most of the efforts were rendered ineffective, as evidenced by the low
capacities utilization, inferior rate capability, poor cycling stability or even complete
electrochemical inactivity, for which the larger size of Na-ion relative to Li-ion is
generally believed to be responsible.[8,9] More specifically, although a few cathode
candidates borrowed from Li-ion chemistry do intercalate/de-intercalate Na-ion
reversibly,[8] the availability of such anode materials is much rare. It has been reported
that carbonaceous materials,[10-14]
TiO2,[15]
Sb2O4,[16]
Sb,[17]
SnSb,[18] Sn,[19-21]
6
SnO2,[22,23] and phosphorus-based materials [24-30] can store/release Na-ion with
decent reversibility.
At the same time, improve the energy density of sodium-ion batteries is another
main task. Layered structure materials, such as P2-type and O3-type materials, have
been extensively studied and considered as one of the most promising cathode material
candidates of the next generation Na-ion batteries.[31-36] Similar to Li[Ni0.5Mn1.5]O4,
P2-type Na2/3[Ni1/3Mn2/3]O2 realizes high operating voltage as a sodium insertion host,
because all the Na ions are reversibly extracted based on the Ni
2+
/Ni
4+
redox.[37,38]
This high operating voltage as Na insertion host is highly beneficial to increase the
energy density for the Na-ion batteries. However, when utilizing the long plateau at
high operation voltage, the transition from P2 to O2 phase together with the large
volume change of O2 phase (more than 20%) is unavoidable (since the O2 structure
has a lower formation energy density than P2 structure in high voltage), which severely
hurts the cycling stability within the high voltage window.[39,40] As a result, the Na2/3-
x[Ni1/3Mn2/3]O2 can only obtain stable electrochemical performance in the region of 0
≤ x ≤ 1/3, which greatly limits the energy and power density of Na2/3[Ni1/3Mn2/3]O2. In
order to achieve sufficiently good cyclability, several attempts were reported such as
the substitution of transition metal ions and lithium ions;[41,42] however, the stable
cycling performance is still limited to ~ 100 cycles due to unclear capacity decay
mechanism.
7
Chapter 2: SnO
2
Coated Carbon Cloth with Surface
Modification as Na-ion Battery Anode
2.1 Introduction
Sn is a promising anode material because it alloys with Na at a high specific
capacity of 847 mAh/g when Na15Sn4 is formed.[43] Several studies of Sn film and
nanostructured anodes were reported with capacities up to 405 mAh/g after 150
cycles.[44,45] On the other hand, SnO2 can deliver a high theoretical sodium storage
capacity of 1378 mAh/g, but with the same volume variation problem as Sn metal, and
hence further surface treatment or matrix scaffold is needed to stabilize SnO2 anode
and improve the electronic conductivity.[46,47].
Figure 2.1 Schematic figure of surface-coating/nanocrystal-active-material-
layer/conductive-soft-platform structure.
In this study, we demonstrate a surface-coating/nanocrystal-active-material-
layer/conductive-soft-platform multilayer nanocomposite microfiber electrode by
using atomic layer deposition (ALD), hydrothermal synthesis method and carbon cloth
(CC) as a soft platform to solve the problems mentioned above, as demonstrated in
8
Figure 2.1. More specifically, we have developed a binder-free multilayer
nanocomposite core-shell microfiber electrode consisting of a hydrothermal
synthesized SnO2 nanocrystal layer on conductive carbon cloth (SnO2/CC) with surface
coatings. The carbon fibers in carbon cloth are intrinsically soft and porous based on
the layered structure with strong interlayer interactions and weak van der Waals
interplanar interactions between adjacent graphene sheets. As the core of the core-shell
structure, during the sodiation and de-sodiation process, the soft carbon fibers can
change in shape to prevent the pulverization of the active material, and the pores in the
fiber together with the aligned core-shell array structure can also provide extra space
to accommodate the volume change to avoid the detachment of the active material layer.
An ALD Al2O3 coating and a hydrothermal carbon coating are further applied on
SnO2/CC electrodes to enhance the cycle life and rate capability. The comparison of
impedance and morphology between fresh and cycled electrodes revealed the
mechanism of capacity decay and how Al2O3 surface coating can protect the active
material layer. The hierarchical core-shell SnO2 anode with designed surface-
coating/nanocrystal-active-material-layer/conductive-soft-platform structure described
is ideal for high power and large scale sodium ion storage. The hydrothermal synthesis,
ALD technology and conductive fiber substrates are scalable for large throughput
manufacturing.
Section 2.2 Experimental
SnO2 coating was synthesized through hydrothermal reaction. 0.1 mol/L SnCl2
was dissolved into 5 ml ethanol and 15 ml DI water, then 1ml HCl acid was added.
Oxidized carbon cloth was immersed into the SnO2 precursor solution overnight in
9
Teflon autoclave and followed by heat treatment at 110
o
C for 6 hours. Obtained
SnO2/CC was washed in DI water and ethanol carefully then dried overnight. The
weight of SnO2 was measured by checking the weight of carbon cloth (CC) before and
after SnO2 hydrothermal deposition. The loading mass density of the SnO2 is ~ 0.9
mg/cm
2
which is calculated by using the weight of SnO2 divided by the area of the
carbon cloth.
Carbon coating process was completed using a hydrothermal method, SnO2
coated carbon cloth was immersed into 0.1 mol/L sucrose solutions, and then
transferred into a Teflon autoclave and heated up to 250
o
C for 3 hours. Obtained
C/SnO2/CC was washed in DI water and then annealed at 500
o
C for half an hours. The
Obtained C/SnO2/CC was washed in DI water and dried in air, then annealed at 500
o
C
for half an hour in argon gas.
The ALD Al2O3 coating was performed on the fabricated SnO2/CC electrode
with a homemade ALD system at the temperature of 90 º C and pressure of 6× 10
-1
torr
in a vacuum chamber.
The SnO2 electrodes were cut into 1× 1 cm
2
pieces and tested in half cells
configuration with Na metal as counter electrodes and 1 M NaPF6 in polycarbonate
(PC) electrolyte with 5% FEC by vol. as additive. The batteries were cycled in the
voltage range of 0.1 V ~ 2.5 V at room temperature. Electrochemical impedance spectra
(EIS) were collected with AC voltage at 5 mV amplitude and frequency range of 100
kHz to 10 mHz. The batteries were fully charged and then rested to reach equilibrium
before the impedance test. All electrochemical experiments were conducted at room
temperature, and all capacities were calculated based on the weight of SnO2.
10
2.3 Results and Discussion
Figure 2.2 (a) Optical image of flexible SnO2/CC binder free electrode, (b) XRD
profile of hydrothermal synthesized SnO2, (c) SEM image of SnO2/CC and (d) cross
section SEM image of a single SnO2/CC fiber and its EDX element mapping profile of
carbon (e) and tin (f).
The carbon fibers are initially treated with nitric acid to enrich the hydroxyl (-
OH) groups on its surface, which can provide strong hydrogen bonding between the
fiber and active material. A thin layer of SnO2 was then synthesized onto carbon cloth
using hydrothermal method, the resulted flexible and binder free electrode is shown in
Figure 2.2a and is denoted as SnO2/CC. In Figure 2.2b, the crystal structure of
11
hydrothermal synthesized SnO2 is confirmed by X-ray diffraction (XRD), and no
impurity is detected from the XRD pattern.
The hierarchical structure of SnO2/CC is illustrated in the SEM image (Figure
2.2c). All of the SnO2/CC fibers have a diameter on the order of 5 µ m uniformly. In
order to confirm the hierarchical core-shell structure of the nanocomposite microfiber,
the freshly cut SnO2/CC sample was further characterized in detail. In Figure 2.2d, a
SnO2 shell layer with a thickness of ~ 500 nm is marked on the cross-section SEM
image. Figure 2.2e and 2.2f show the EDX mapping profiles of the carbon and tin
elements in Figure 2.2d, both of which show a good agreement with the contrast profile
of Figure 2.2d. Both SEM images and EDX mapping cooperatively reveal that all
carbon fibers are coated with a SnO2 layer uniformly to form SnO2/CC hierarchical
core-shell nanocomposite microfibers. Similar core-shell structure with conductive
core has been proven effective in resolving issues associated with electron-transfer
kinetics along the high aspect-ratio rod- or wire-like materials [38]. Compared to
conventional rigid metallic substrates, the carbon cloth is extremely soft and porous
with a high capacity for electrolyte absorption which can provide a dual diffusion
channel for Na-ions. In order to study the effect of substrate and surface coating on
electrochemical performance, additional sucrose hydrothermal carbon coating and
ALD Al2O3 coating are also applied on the SnO2/CC samples and are denoted as
C/SnO2/CC and Al2O3/SnO2/CC, respectively.
12
Figure 2.3 (a) TEM image of a large bulk piece of SnO2 layer coated on CC, (b, c)
High resolution TEM image of SnO2 nanocrystal at different magnification, (d) FFT
image of the area marked in (b). TEM images of the carbon layer (e) and Al2O3 layer
(f) coated on CC
The hierarchy, crystalline and configuration of SnO2/CC nanocomposite with
different surface coating were further confirmed by transmission electron microscopy
(TEM). In Figure 2.3a, the TEM image shows a large bulk piece from the SnO2
cylindrical active material shell layer coated on a single carbon fiber. According to the
high resolution TEM image shown in Figure 2.3b and the FFT image (Figure 2.3d) of
the area marked in Figure 2.3b, the SnO2 layer is polycrystalline with uniform grain
size at ~ 5 nm. The high-resolution TEM images of carbon and Al2O3 surface coatings
are shown in Figure 2.3e and 2.3f. Several carbon layers are observed in Figure 2.3e,
the thickness of the carbon coating is 7.8 nm, which is partially graphitized after the
annealing. In Figure 2.3f, the thickness of the uniform ALD Al2O3 coating is ~ 8.3 nm
13
and in an amorphous nature. In order to study the degree of graphitization of the carbon
coating, we designed the following control experiment and characterized the annealed
sample with Raman spectrum equipped with a 532 nm laser source. Sucrose solution
with 0.1 mol/L concentration was dropped onto a quartz substrate and dried in air, and
then the substrate was annealed at 500
o
C for half an hour in argon gas.
Figure 2.4 (a) Cycling stability and Coulombic efficiency of SnO2/CC, C/SnO2/CC
and Al2O3/SnO2/CC. Potential profiles of (b) C/SnO2/CC and (c) Al2O3/SnO2/CC in
cycling test. (d) Rate performance of C/SnO2/CC and Al2O3/SnO2/CC.
The cycling performance of SnO2/CC, C/SnO2/CC and Al2O3/SnO2/CC
electrodes was investigated by galvanostatic charge and discharge of the electrodes
between 0.05 and 2.5 V at the current density of 0.1 C ( 134 mA/g), and the results are
shown in Figure 2.4a. Electrodes assembled in a traditional way by mixing carbon black,
14
PVDF binder and commercial SnO2 powder were also tested for comparison. For the
SnO2/CC sample without any surface coating, its initial specific charge capacity is 529
mAh/g and it maintains 30% of initial capacity after 40 cycles, which is an obvious
improvement over the conventional electrodes made of commercial SnO2 powder. The
degradation was suppressed because the high surface area, porous structure and the
intrinsic soft characteristics of carbon cloth can effectively accommodate the volume
charge during sodiation and de-sodiation process. After 40 cycles, the capacity decay
becomes slower and finally the capacity reaches 74 mAh/g at the 100th cycle. The
cycling stability is further enhanced by a thin layer of carbon and Al2O3 coating. The
C/SnO2/CC sample reaches a 522 mAh/g initial charge capacity which is similar to the
SnO2/CC sample without any surface coating, but it shows a 60% retention which is
much higher than the SnO2/CC sample. This demonstrates that the designed surface-
coating/nanocrystal-active-material-layer/conductive-soft-platform structure can
greatly suppress the volume change and the pulverization of SnO 2 crystal to improve
the cycling stability. Meanwhile, although the Al2O3/SnO2/CC anode has a 470 mAh/g
initial capacity which is slightly lower than C/SnO2/CC, it shows the best retention of
80% and delivers a 377 mAh/g capacity which is even 63 mAh/g higher than
C/SnO2/CC after 100 cycles.
Aside from capacity, cycling potential profiles of C/SnO2/CC and
Al2O3/SnO2/CC anodes for Na-ion batteries at 1, 2, 5, 10, 50 and 100
th
cycle are
depicted in Figure 2.4b and 2.4c. All curves shown in the two figures represent the
typical reaction between SnO2 and Na, and they also reveal that the additional surface
coating processes didn’t reduce SnO2 to SnO or Sn metal. For the C/SnO2/CC anode,
15
potential plateaus in discharge and charge curves gradually disappeared, suggesting
that part of the SnO2 layer disintegrated and detached electronically from the carbon
cloth. Comparing to the C/SnO2/CC anode, the Al2O3/SnO2/CC anode suffers a larger
hysteresis between charge and discharge curves due to the insulating Al2O3 coating.
However it maintained these distinct step plateaus even after 100 deep cycles, which is
a much pronounced difference comparing with the carbon coated samples. For the
Al2O3/SnO2/CC anode, a bump located around 1.3 V is observed on the first discharge
curve, which can be considered as the stress release happened between the Al2O3
surface coating layer and the SnO2 nanocrystal layer during the sodiation process, since
no similar peak is observed from the C/SnO2/CC sample. It is clear that graphite carbon
layer and dense Al2O3 layer are not as soft as the carbon cloth fibers, the stress
introduced by the first sodiation process may lead to the activation of the carbon fibers,
reconstruction of the active material layer and creating more contact between the active
material and electrolyte. The TEM image clearly showed the conformal coating of
Al2O3 and the electrochemical performance before and after Al2O3 coating, which
revealed that the Al2O3 coating prevents the mechanical degradation of SnO2 and also
explains the cycling stability improvement of the Al2O3/SnO2/CC sample.
In Figure 2.4d, the rate capability test results show that the C/SnO2/CC anode
can deliver average capacities of 501, 460, 422, 391 and 342 mAh/g at 0.1, 0.5, 1, 5
and 10 C charge/discharge rate. Even at a 20 C and 30 C charge/discharge rate, a 255
mAh/g and a 144 mAh/g specific charge capacities are obtained. For the
Al2O3/SnO2/CC sample, slightly lower capacities are observed at different current rates
compared to the C/SnO2/CC sample: a 455 mAh/g average capacity at 0.1 C rate,
16
followed by 413, 331, 245 and 95 mAh/g at 0.5, 1, 5 and 10 C charge/discharge rates.
At 20 C and 30 C, the Al2O3/SnO2/CC anode gives very low capacities (< 10 mAh/g)
and there is no typical potential plateau of SnO2 on the charge/discharge curves. When
the rate was restored to 0.1 C, the Al2O3/SnO2/CC anode showed excellent stability
with a 447 mAh/g average capacity which is 27 mAh/g higher than that of the
C/SnO2/CC anode. Obviously, due to the high electrical conductivity of the carbon
cloth platform, both carbon and Al2O3 coated samples gain fast charge/discharge
capability. Especially for the C/SnO2/CC sample, the conductive carbon fiber core and
carbon coating shell provide dual channels for electron transfer and acquire supreme
fast charge-discharge capability.
17
Figure 2.5 The electrochemical impedance spectrum (EIS) profiles of (a) C/SnO2/CC
and (b) Al2O3/SnO2/CC fresh cell and cell after 100 cycles. The SEM images of (c)
C/SnO2/CC and (d) Al2O3/SnO2/CC after 100 cycles.
The electrochemical impedance spectrum (EIS) of C/SnO2/CC and
Al2O3/SnO2/CC cells before and after 100 cycles is characterized and shown in Figure
2.5a and 2.5b. For the fresh C/SnO2/CC cell, the contact resistance is as small as 5
and the charge transfer resistance of was calculated to be 165 based on the semicircle.
After 100 cycles, the contact and charge transfer resistance increased to 20 and 191
. Meanwhile, the fresh Al2O3/SnO2/CC cell have a 372 charge transfer resistance
and a 105 contact resistance, the increased contact and charge transfer resistance can
be attributed to the dense and insulating nature of the Al2O3 coating. After 100 cycles,
the charge transfer resistance calculated from the first semicircle almost remains the
same, however, the contact resistance increased to 175 , and another big semicircle
with a 450 charge transfer resistance appears following the first semicircle, which
implies a strong solid electrolyte interface (SEI) film already formed on the Al2O3
surface. In order to prove that hypothesis and gain insight into the capacity fading
mechanism, SEM images of two samples with different surface coatings are obtained
after washed with PC and DI water carefully to remove the Na salt. In Figure 2.5c,
there is no obvious SEI film on the surface of the C/SnO2/CC sample after cycling.
However, as we expected, SnO2 nanocrystals agglomerated together and became
porous bulks. Parts of the active material layer also detached from the carbon cloth,
which should be considered as the main reason to the capacity decay of C/SnO2/CC
18
sample. In the SEM image of the Al2O3/SnO2/CC sample after cycling, no obvious
agglomerates and detachments were found, and a significant SEI film was observed at
the same time, which gradually started to decompose after exposing to the electron
beam for tens of seconds. Since the washing process for both samples are the same, we
believe that the SEI film formed on Al2O3 surface has different composition or a larger
thickness than that on the carbon coating surface. The different SEI film observed from
Al2O3 coated sample may be formed because of the unique surface property of Al2O3
layer or the physical and electrochemical reaction between the electrolyte and Al2O3
surface coating. It is obvious that although the Al2O3 coating together with strong SEI
film leads to the increasing of the contact and charge transfer resistance of
Al2O3/SnO2/CC electrode, on the other hand, they provide the superior mechanical
support to the SnO2 active layer and suppress the volume change and pulverization of
SnO2 layer during sodiation and de-sodiation process. Compared to C/SnO2/CC sample,
the strong SEI film on Al2O3 surface also consumes more Na-ion during sodiation
process and results in a lower Coulombic efficiency during cycling as shown in Figure
2.4a.
2.4 Summary
In summary, hierarchical core-shell nanocomposite anode consisting of
individual conductive carbon fiber, a SnO2 intermediate layer and carbon or Al2O3
surface coating, was fabricated by hydrothermal and ALD method. The
Al2O3/SnO2/CC anode maintained a 371 mAh/g specific charge capacity at 100
th
cycle
which demonstrated superior electrochemical stability and the C/SnO2/CC anode
delivered a 342 mAh/g and a 144 mAh/g capacity at 10 C and 30 C high
19
charge/discharge current rate. The conductive soft platform as well as the precise
hierarchical control of various sublayers of materials in designed order was believed to
function synergistically to maintain an anode host mechanically, electronically, and
electrochemically active and stable, despite its large volume change upon
sodiation/desodiation cycles and semiconductor material nature. In particular, the
designed surface-coating/nanocrystal-active-material-layer/soft-platform core-shell
electrode system minimized the stress of the volume change and maximized the Na-
ion transport kinetics of SnO2 anode, which have been two rather severe challenges that
this kind of Na-ion host is facing. The excellent rate capability and cycling stability,
the easy processability of carbon fibers and the mature hydrothermal synthesis and
ALD technology make this surface-coating/nanocrystal-active-material-layer/soft-
platform core shell system an excellent candidate for Na-ion storage. Although Na-ion
batteries has a great potential for the low-cost and material-abundant large scale energy
storage application, we believe the fast charge/discharge capability is very meaningful
to keep a high power density to match with several high power cathodes for full Na-
ion batteries. The superior cycling performance, the excellent rate capability combined
with the simplicity of the fabrication process, represents a new strategy for the
development of inexpensive and versatile synthesis techniques for Na-based energy
storage applications.
20
Chapter 3: Layered P2-Na
2/3
[Ni
1/3
Mn
2/3
]O
2
as High-V oltage
Cathode for Sodium-Ion Batteries: The Capacity Decay
Mechanism and Al
2
O
3
Surface Modification
3.1 Introduction
Similar to Li[Ni0.5Mn1.5]O4, P2-type Na2/3[Ni1/3Mn2/3]O2 can serve as a sodium
cathode at high operating voltages based on the Ni
2+
/Ni
4+
redox.[37,38] However, the
P2-O2 crystal phase transition and the large volume change of the O2 phase (more than
20%) is unavoidable since the O2 structure has a lower formation energy density than
the P2 structure at high voltages.[39,40] This severely damages the cycling stability
within the high voltage window. As a result, the Na2/3-x[Ni1/3Mn2/3]O2 can only obtain
stable electrochemical performance in the region of 0 ≤ x ≤ 1/3, which greatly limits
the energy and power density of Na2/3[Ni1/3Mn2/3]O2 for real applications. In order to
achieve sufficiently good cyclability, several attempts were reported such as the
substitution of transition metal ions and lithium ions.[41,42] Despite these substitutions,
the stable cycling performance has still been limited to ~ 100 cycles due to an unclear
capacity decay mechanism.
21
Figure 3.1 Schematic figure of the exfoliation during the sodiation and de-sodiation
process of the P2-Na2/3[Ni1/3Mn2/3]O2 particle.
In this study, we first explored the electrochemical performance and capacity
decay mechanism of Na2/3[Ni1/3Mn2/3]O2 within a high voltage window from 2.5 V to
4.3 V. Based on the discovered degradation mechanism, we applied a surface-coating
solution to stabilize the Na2/3[Ni1/3Mn2/3]O2 electrode using a simple wet chemistry
method. As illustrated schematically in Figure 3.1, the exfoliation phenomena resulting
from an unfavorable crystal structure transition was observed in the after-cycling
Na2/3[Ni1/3Mn2/3]O2 particles. Furthermore, an Al2O3 surface coating was applied on
Na2/3[Ni1/3Mn2/3]O2 particles (Al2O3-Na2/3[Ni1/3Mn2/3]O2) and effectively solved the
issues discussed above and enhanced the cycling life. We employed an Al2O3 coating
because Al2O3 been shown to be an excellent coating material for mitigating volume
expansion and contraction during Li/Na ion insertion and extraction.[48-50] The
comparison between the fresh and cycled electrodes together with the enhanced cycling
performance by robust Al2O3 coating revealed the mechanism of capacity decay, and
further demonstrated that the Al2O3 surface coating can suppress the side reaction and
protect the layered metal oxide particles during long cycling within the high voltage
window. The designed Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathode is ideal for large scale high
power/energy Na-ion storage, not only because of its excellent cycling stability and
improved voltage profile, but also due to the scalability of the solid state reaction of
Na2/3[Ni1/3Mn2/3]O2 and wet chemistry coating.
22
3.2 Experimental
Na2/3[Ni1/3Mn2/3]O2 particles are synthesized through a solid state reaction.
Nickel acetate (Ni(Ac)2· 4H2O) and manganese acetate (Mn(Ac)2· 4H2O) were first
mixed and hand-milled in a mortar with a molar ratio of Ni: Mn=1: 2. The mixture was
heated up to 500
o
C for 5 h with a heating rate of 3
o
C/min. Sodium acetate (NaAc· 2H2O)
was then added to the mixture with a molar ratio of Na: Ni: Mn=2.1: 1: 2 (5% excess
sodium acetate was added in order to make up for the volatilization of Na during
calcination). After that, the mixture was heated to 500
o
C for 5 h again. The mixture
was hand-milled once more and sintered at 950
o
C for 10 h followed by annealing at
700
o
C for 10 h and quenched after the annealing.
The Al2O3 coating was performed on the Na2/3[Ni1/3Mn2/3]O2 sample using a
wet chemistry method targeting at 5 wt.% of the total mass. The Al(NO3)3 salt was
dissolved into deionized water with a concentration of 0.1 mol/L before a suitable
amount of ammonia was added into the solution. The resulting solution was milled with
Na2/3[Ni1/3Mn2/3]O2 powder and dried under stirring overnight. The mixture was
annealed at 200
o
C for 10 h, then sintered at 650
o
C for 10 h with a heating rate of
3
o
C/min and quenched after the annealing.
X-ray diffraction (XRD) Bruker AMX-500 diffractometer with Cu-Kα
radiation source operated at 44 kV. The surface morphology of the samples was
characterized by scanning electron microscopy (SEM, JOEL JSM-7001). A field
emission transmission electron microscopy (TEM, JOEL JEM 2100F) was employed
to obtain the TEM images and scanning transmission electron microscopy (STEM)
images.
23
Electrodes were prepared by casting the slurry containing 85 wt.% active
material, 10 wt.% carbon black and 5 wt.% polyvinylidene fluoride (PVDF) binder
onto an Al foil. The loading mass on the Al foil current collector was ~ 1.5 mg/cm
2
.
Electrochemical tests were conducted in CR2032 coin cells with Na metal as
counter electrodes and 1.2 M NaClO4 in polycarbonate (PC) electrolyte with 5%
fluoroethylene carbonate (FEC) by volume as additive. The stainless steel coin cell
parts were coated with Al2O3 to prevent the side electrochemical reactions that occur
at high voltage. Batteries were cycled in the voltage range of 2.5 V ~ 4.3 V at room
temperature. All capacities were calculated based on the weight of Na2/3[Ni1/3Mn2/3]O2
active material.
Electrochemical impedance spectra (EIS) and cyclic voltammetry (CV) were
collected using a GAMRY Reference 600 test station. In the EIS tests, cells were tested
as prepared and after charging to 4.3 V with a 10-hour rest, the scan range is from 10
6
Hz to 5mHz. In the CV tests, the scan rate is 0.01 mV/s and the scan range is from 2.5
to 4.5 V.
24
3.3 RESULTS AND DISCUSSIONS
Figure 3.2 (a) SEM image of as-prepared Na2/3[Ni1/3Mn2/3]O2. SEM image of a single
Na2/3[Ni1/3Mn2/3]O2 particle (b), and the DEX mapping profile of elements Na (c), Ni
(d) and Mn (e). (f) The TEM image of as-prepared Na2/3[Ni1/3Mn2/3]O2 and (g) the
SAED pattern of the area mark in red circle in (f). The high-resolution TEM image of
(h) as-prepared Na2/3[Ni1/3Mn2/3]O2 particle with Na ion channel marked and (i) Al2O3
coated Na2/3[Ni1/3Mn2/3]O2.
25
In Figure 3.2a, the X-ray diffraction (XRD) pattern of the as-prepared
Na2/3[Ni1/3Mn2/3]O2 powder agrees well with the P2 type Na2/3[Ni1/3Mn2/3]O2 sample
(PDF 54-0894). The XRD pattern was refined through LeBail fitting and the d-spacing
of the (002) reflection is confirmed as 5.59 angstrom accordingly. From the scanning
electron microscopy (SEM) image of the as-prepared Na2/3[Ni1/3Mn2/3]O2 powder
(Figure 3.2b), the dimension of the particles varies from several micrometers to
hundreds of nanometers. The energy-dispersive X-ray (EDX) mapping profiles of a
single particle (Figure 3.2c) are presented in the order of Ni (Figure 3.2d) and Mn
(Figure 3.2e) elements respectively. From the transmission electron microscopy (TEM)
image of Na2/3[Ni1/3Mn2/3]O2 particles (Figure 3.2f), fine hexagonal crystal structures
are observed, and the crystallinity of the particles is confirmed to be single crystal from
the selected area electron diffraction (SAED) pattern presented in Figure 3.2g, where
the SAED area is marked by the red circle in Figure 3.2f. The (002) lattice planes with
5.5 angstrom spacing are observed in the high-resolution TEM image (Figure 3.2h),
which can be confirmed as both the major ion transport channel and storage space for
Na ions. The lattice parameter reaches a good agreement with the d-spacing of the (002)
reflection discussed above in Figure 3.2a. For the Al2O3-coated sample (Al2O3-
Na2/3Ni1/3Mn2/3O2), a ~ 12 nm amorphous Al2O3 layer was clearly observed under TEM
(Figure 3.2i). The coating is conformal both along and perpendicular to the (002) lattice
plane around the particle.
26
Figure 3.3 (a) Cycling performance of the as-prepared Na2/3[Ni1/3Mn2/3]O2 and Al2O3-
Na2/3[Ni1/3Mn2/3]O2. The charge and discharge profiles of (b) Na2/3[Ni1/3Mn2/3]O2
electrodes and (c) Al2O3-Na2/3[Ni1/3Mn2/3]O2 electrode. (d) Rate performance of the as-
prepared Na2/3[Ni1/3Mn2/3]O2 and Al2O3-Na2/3[Ni1/3Mn2/3]O2. Cyclic voltammetry
profiles of (e) as-prepared Na2/3[Ni1/3Mn2/3]O2 cathode and (f)Al2O3-
Na2/3[Ni1/3Mn2/3]O2 cathode.
The cycling performance of Na2/3[Ni1/3Mn2/3]O2 and Al2O3-Na2/3[Ni1/3Mn2/3]O2
cathodes was investigated by galvanostatic charge and discharge of the electrodes
between 2.5 V and 4.3 V at a current density of 0.5 C rate (86.5 mA/g) as shown in
Figure 3.3a. Two electrodes presents high initial specific discharge capacities around
160 mAh/g at the first sodiation process, which is close to the theoretical capacity (173
mAh/g) of P2-Na2/3[Ni1/3Mn2/3]O2. The high initial capacities of the two electrodes
suggest that both samples can be fully sodiated/de-sodiated in the voltage range from
27
2.5 V to 4.3 V. The as-prepared Na2/3[Ni1/3Mn2/3]O2 and Al2O3-Na2/3[Ni1/3Mn2/3]O2
sample exhibit first charge capacities of 193 and 192 mAh/g, respectively. The extra
charge capacity might involve the oxidative decomposition of sodium propyl carbonate
generated at the Na metal negative electrode during the first cycle. The
Na2/3[Ni1/3Mn2/3]O2 cathode without any surface coating suffered from a rapidly fading
capacity, retaining only 53.3% at the 50th cycles and 26.8% at the 300
th
cycle. In
contrast, the Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathode maintains 88.4% of its initial capacity
after 50 cycles, which is a visible improvement over the uncoated Na2/3[Ni1/3Mn2/3]O2
samples. At the 300
th
cycle, its capacity stabilized at ~ 115 mAh/g. Thus, the cycling
stability was significantly extended by a thin layer (12 nm) Al2O3 coating.
For the coulombic efficiency of the two electrodes presented in Figure 3.3a, the
Na2/3[Ni1/3Mn2/3]O2 cathode has an 85.3% initial efficiency that quickly rises to ~ 100%
after the 2
nd
cycle. Meanwhile the Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathode acquires 80.7%
initial efficiency. Although Al2O3 coating is known to be able to suppress electrolyte
oxidation at high voltage, here the Al2O3 coated sample shows a slightly lower initial
coulombic efficiency. This may be due to the Al2O3 coating layer further increases the
kinetic battier for extraction of Na ions out of the metal oxide component. The slightly
lower efficiency of the Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathode in the following cycles may
be related to the low electronic conductivity of the Al2O3 coating.
Aside from capacity, voltage profiles actually reveal additional details about the
electrochemical sodiation/de-sodiation process in these cathode materials. The voltage
profiles of the Na2/3[Ni1/3Mn2/3]O2 and Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathodes are
depicted in Figure 3.3b and 3.3c respectively, with cycle numbers labeled on individual
28
curves. The uncoated Na2/3[Ni1/3Mn2/3]O2 sample shows three potential plateaus at 3.3-
3.4, 3.7 and 4.2 V in its first de-sodiation process corresponding to the Na content of
2/3, 1/2 and 1/3, respectively. On the other hand, the Al2O3-Na2/3[Ni1/3Mn2/3]O2 shows
a slope from 3.3 to 3.7 V and a plateau at 4.2 V. For the charge curves of the 2
nd
, 10
th
,
50
th
and 300
th
cycle, all plateaus of the uncoated Na2/3[Ni1/3Mn2/3]O2 decay rapidly with
increasing cycle number. However, for the Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathode, the
charge capacity decay mainly happens on the high potential plateau at 4.2 V, which
implies the P2-O2 phase transition only occurs in the region of 0 ≤ x ≤ 1/3 for the Na2/3-
x[Ni1/3Mn2/3]O2 cathode. For all discharge curves of both samples, two major potential
plateaus at 4.1 and 3.7 V were observed for both cathodes. However, in contrast to the
uncoated Na2/3[Ni1/3Mn2/3]O2, the discharge plateaus of the Al2O3-Na2/3[Ni1/3Mn2/3]O2
evolve to slopes, which indicates that the Na ion transition in the Al2O3-
Na2/3[Ni1/3Mn2/3]O2 is more inclined to a solid solution diffusion mode. Similar to the
charge curves, the discharge plateaus of the uncoated Na2/3[Ni1/3Mn2/3]O2 were
shortened as the cycle number increased, but the Al2O3-Na2/3[Ni1/3Mn2/3]O2 showed a
much smaller capacity decay owing to the Al2O3 surface coating.
In the rate capability test (Figure 3.3d), the capacities of the two electrodes decay
in the first 60 cycles from 0.1 C to 5 C charge/discharge rate. The uncoated
Na2/3[Ni1/3Mn2/3]O2 sample shows a larger fading rate than the Al2O3-
Na2/3[Ni1/3Mn2/3]O2 sample, which reaches a good agreement with the cycling
performance presented in Figure 3.3a. However, the capacity of the Al2O3-
Na2/3[Ni1/3Mn2/3]O2 was stabilized after 60 cycles, and the cathode can deliver average
capacities of 140.2, 132.4, 116.9, and 65.3 mAh/g at 0.1, 0.5, 1 and 3 C rate. At 5 C
29
charge/discharge rate, the capacity of the Al2O3 coated sample decays to almost zero;
on the other hand, the uncoated Na2/3[Ni1/3Mn2/3]O2 cathode can deliver around 20
mAh/g capacity, which suggests that the Al2O3 coating increases the electronic
resistance and further enlarges the polarization of the material. The capacity of the
Al2O3-Na2/3[Ni1/3Mn2/3]O2 was stabilized at ~ 130 mAh/g at 0.1 C after 110 cycles,
meanwhile the capacity of the uncoated Na2/3[Ni1/3Mn2/3]O2 sample continued to decay.
The sodiation and de-sodiation processes of as-prepared Na2/3[Ni1/3Mn2/3]O2 and
Al2O3-Na2/3[Ni1/3Mn2/3]O2 cathodes were also characterized using cyclic voltammetry
(CV), as shown in Figure 3.4e and 3.4f. For the uncoated Na2/3[Ni1/3Mn2/3]O2 sample,
four peaks at 3.3, 3.6, 4.1 and 4.5 V were observed in the first and subsequent de-
sodiation processes. The peak at 4.5 V in the de-sodiation curves can be mainly
assigned to the side reactions at high voltage such as electrolyte decomposition and
surface corrosion of the material. One sharp peak at 3.9 V with fast decay and another
two peak at 3.3 and 3.5 V were observed in the subsequent sodiation processes. In
Figure 3.4f, the Al2O3-Na2/3[Ni1/3Mn2/3]O2 sample shows a very different curve for the
first de-sodiation process with three peaks located at 3.6 and 4.1 V, indicating a
different SEI formation mechanism on the Al2O3 surface. Compared to the uncoated
Na2/3[Ni1/3Mn2/3]O2 sample, Al2O3-Na2/3[Ni1/3Mn2/3]O2 shows similar peak positions in
the following cycles but with broader peak features except the peaks at 3.3 V, which
becomes to a small shoulder. The peak at 4.4 V was suppressed for the Al2O3-
Na2/3[Ni1/3Mn2/3]O2 sample, indicating that the Al2O3 coating can effectively mitigate
side reactions at high voltage.
30
Figure 3.4 (a) SEM image of after-cycling Na2/3[Ni1/3Mn2/3]O2 electrode with the
exfoliation and carbon black areas marked. (b) STEM of an after-cycling
Na2/3[Ni1/3Mn2/3]O2 particle. (c) TEM image of an after-cycling Na2/3[Ni1/3Mn2/3]O2
particle. (d) Enlarged TEM image of the area marked in (c). (e) High-resolution TEM
image of the exfoliation opening in (d). (f) The SAED pattern of the area marked in (c).
To gain insight into the capacity fading mechanism of Na2/3[Ni1/3Mn2/3]O2, SEM
and TEM images were collected after 300 cycles. In the after-cycling SEM image of
the uncoated Na2/3[Ni1/3Mn2/3]O2 electrode shown in Figure 3.4a, several
semitransparent layers with clear hexagonal shape were observed, which were
identified as the exfoliated layers from the Na2/3[Ni1/3Mn2/3]O2 single particle. This
exfoliation phenomenon can also be confirmed from the after-cycling scanning
transmission electron microscope (STEM) image in Figure 3.4b, which shows one
Na2/3[Ni1/3Mn2/3]O2 single crystal particle exfoliated into 3 pieces. More details about
31
the exfoliation can be found in the after-cycling TEM image of the Na2/3[Ni1/3Mn2/3]O2.
In Figure 3.4c and its enlarged view (Figure 3.4d), a very obvious exfoliation opening
can be clearly observed. Furthermore, in the high-resolution TEM image (Figure 3.4e),
the exfoliation opening can be defined to be along the (002) plane direction, which is
the Na ion storage space and also the major transfer channel described in Figure 3.2h.
Figure 3.5 The XRD pattern of the as-prepared Na2/3[Ni1/3Mn2/3]O2, after-cycling
Na2/3[Ni1/3Mn2/3]O2 and after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2. (a) is the enlarged
view of (b) from 15
o
to 17
o
.
The exfoliation phenomena observed above also agrees well with the changes in
the XRD pattern of the after-cycling Na2/3[Ni1/3Mn2/3]O2 electrode: in Figure 3.5b and
its enlarged view from 15
o
to 17
o
(Figure 3.5a), the (002) lattice peak shows an obvious
shift to the left from 15.8
o
to 15.3
o
, suggesting the (002) lattice space was expanded
during cycling, and further proves that exfoliation should be a consequence of the
lattice space expansion. Compared to the single crystal SAED pattern of the as-
(a) (b)
32
prepared sample, the SAED pattern of the after-cycling Na2/3[Ni1/3Mn2/3]O2 indicates a
phase transition during cycling within the high voltage window, which may be due to
the introduction of the O2 stacks. Part of this crystal structure change can also be
detected from the XRD pattern of the after-cycling Na2/3[Ni1/3Mn2/3]O2 electrode in
Figure 3.5a and 3.5b where almost all peaks show phase separation besides the shifted
(002) peak. This implies that the phase separation occurs during cycling and part of this
transition is irreversible and harmful to the crystal structure stability. Since the P2-O2
transformation requires no bond breaking between oxygen and transition metal, and
also considering the O2 phase has a larger volume change (more than 20%) than the
P2 phase during sodiation/de-sodiation, it is reasonable to believe that the exfoliation
occurred during the Na-ion intercalation in the high voltage region. Besides the
insertion and extraction of Na ions, the exfoliation also possibly involves side reactions
such as surface oxidation at high voltage, since the exfoliation was observed starting
from the surface of the bulk material. Once the exfoliation happens, the P2-O2 phase
transition loop will be broken permanently and part of the Na ion storage space will be
eliminated. The exfoliation is definitely harmful for a cathode material that aims for
stable electrochemical performance.
33
Figure 3.6 (a) The SEM image of the after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2
electrode. (b) The high-resolution TEM image of an after-cycling Al2O3-
Na2/3[Ni1/3Mn2/3]O2 particle. (c) The SAED pattern of the after-cycling Al2O3-
Na2/3[Ni1/3Mn2/3]O2 particle.
In hope of investigating how the Al2O3 coating can improve the cycling stability
of the cathode material, SEM and TEM images with the SAED pattern of the after-
cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2 electrode are also collected. In the SEM image of
after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2 electrode (Figure 3.6a), several active material
particles are embedded in the electrode without any sign of exfoliation. In the high-
resolution TEM image of the after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2 (Figure 3.6b),
part of the Al2O3 coating is peeled off from the particles, which may be due to the large
volume change of the O2 phase during Na ion storage and release. Whether the surface
coating can survive in longer cycles need more investigation. However, no exfoliation
opening was found in the TEM image, which proved that the Al2O3 coating efficiently
protected the Na2/3[Ni1/3Mn2/3]O2 crystal structure by suppressing the exfoliation. From
the SAED pattern in Figure 3.6c, only a slight phase transition is detected for the after-
cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2 compared to the uncoated sample. The morphology
difference between the after-cycling samples with and without surface coating can also
be observed from their XRD patterns. In Figure 3.5a, the shift of the (002) peak is only
~ 0.1
o
for the after-cycling Al2O3-Na2/3[Ni1/3Mn2/3]O2 compared to the as-prepared
Na2/3[Ni1/3Mn2/3]O2, which is much smaller than the shift of the after-cycling
Na2/3[Ni1/3Mn2/3]O2 without surface coating (~ 0.5
o
). In Figure 3.5b, no obvious peak
separation was detected from the XRD pattern of the after-cycling Al2O3-
34
Na2/3[Ni1/3Mn2/3]O2 sample. The difference of the XRD data of those two after-cycling
samples suggests that the Al2O3 surface coating can efficiently suppress the phase
separation and avoid the crystal structure damage introduced by the phase transition
during long-term cycling.
Our results demonstrate that the exfoliation of the Na2/3[Ni1/3Mn2/3]O2 involves
both a crystal phase transition and volume expansion of the O2 phase. Thus, we believe
that external mechanical support is necessary in order to stabilize the crystal structure
and enhance the electrochemical performance. The above experimental data revealed
that the Al2O3 coating prevents the mechanical degradation of Na2/3[Ni1/3Mn2/3]O2 and
also explains the electrochemical cycling stability improvement of the Al2O3-
Na2/3[Ni1/3Mn2/3]O2 sample. It is clear that the Al2O3 surface coating cannot only reduce
the side reaction at high voltage, but can also give mechanical support to help the bulk
material to maintain its layered structure, which can increase the reversibility of the P2-
O2-P2 phase transition loop during charge/discharge process. This Al2O3 surface
coating unlocks the stable cycling performance of P2-Na2/3[Ni1/3Mn2/3]O2 within high
voltage window and releases the high energy density of this layered structure cathode
material for Na-ion batteries. Other surface coating methods or scaffold matrices may
also improve the performance of the layered structure cathode for Na-ion batteries
based on the exfoliation induced decay mechanism discovered in this paper and will be
studied in the future.
3.4 CONCLUSIONS
In summary, we first investigated the capacity decay mechanism of the layered
structure P2-Na2/3[Ni1/3Mn2/3]O2, and the exfoliation phenomena associated with the
35
phase transition was analyzed. Furthermore, the cycling stability of the
Na2/3[Ni1/3Mn2/3]O2 cathode in the high voltage range was significantly enhanced and
the exfoliation was efficiently suppressed by a thin layer of Al2O3 surface coating,
which demonstrated that the Al2O3 coating can provide superior protection to this P2
type layered structure material. The stable high voltage Al2O3-Na2/3[Ni1/3Mn2/3]O2
represents a new strategy for the development of inexpensive and high power/energy
Na-based energy storage applications.
Chapter 4: Red Phosphorus Nano-Dots on Reduced Graphene
Oxide as a Flexible and Ultra-Fast Anode for Sodium-Ion
Batteries
4.1 Introduction
Phosphorus offers much promise because it alloys with sodium at a high
specific capacity of 2596 mAh/g when Na3P is formed.[51] Due to the unstable
property and toxicity of white phosphorus, black phosphorus and red phosphorus are
the two major forms explored for sodium-ion storage. Thanks to the high conductance
and the two-dimensional material nature, several promising studies of black
phosphorus were reported with stable high capacity with impressive rate capability up
to 10 C charge/discharge rate were reported.[52-55] However, the intrinsic challenges
stemming from the high cost and sophisticated synthesis of black phosphorus promote
red phosphorus to a better position in sodium-ion chemistry. Recently, several red
36
phosphorus anodes developed with ball-milling technique and carbonaceous scaffold
were presented to solve the poor conductivity, slow sodium ion transport kinetics and
large volume variation during cycling.[56-60]
Figure 4.1 Schematic description of P@RGO synthesis.
In this study, we demonstrate an effective solution to solve the above issues of
red phosphorus, which allows the growth of red phosphorus particles with diameters
vary from serval hundreds to tens of nanometers on reduce graphene oxide (RGO)
densely and uniformly by physical vapor deposition (PVD) method, resulting in the
P@RGO composite. The preparation steps of P@RGO are presented in Figure 4.1. Red
phosphorus precursor and RGO powder was placed at two ends of a quartz tube, which
was sealed into an ampoule under vacuum with the materials loaded. During a 15-
minute heat-treatment at 600
o
C, red phosphorus precursor was vaporized and filled into
the gaps between RGO layers due to the vacuum environment. After that, the
phosphorus nano-dots would grow on the RGO surface during the condensation of the
phosphorus vapor. In this design, the RGO sheets serve as an intimate electronic
37
pathway between the red phosphorus active material and external circuit, and the
precisely controlled deposited red phosphorus nano-dots effectively facilitates the
sodium ion transfer, thus accelerating the electrochemistry reaction rate of the whole
battery system. The free space between the RGO sheets can accommodate the volume
variation of the red phosphorus to stabilize the P@RGO anode further. Moreover, the
architectural P@RGO is proved to be a convenient and reliable solution to sodium-ion
batteries which is not only mechanically stable but also amenable toward red
phosphorus PVD processes, hence allowing the architecture of flexible power sources
for wearable electronics.
4.2 Methods and Experimental
Graphene oxide water suspension (Graphene Laboratories Inc.) was further
chemically exfoliated by a modified method described in the supporting information
with reduction details. Red phosphorus fine powder with 99 % purity (Spectrum
Chemical Mfg. Corp.) was used as precursor. The red phosphorus precursor was dried
at 90
o
C to remove the moisture and meshed with a 30 µ m mesh after drying. Then the
precursor and RGO powder were placed at two ends of a quartz tube with 10 cm in
length with phosphorus to RGO mass ratio of 3:1, and then the quartz tube was sealed
into an ampoule under vacuum. The ampoule was annealed in a tube furnace at 600
o
C
for 15 minutes and maintained at 280
o
C for 10 hours to covert white phosphorus to red
phosphorus. After the heat-treatment process, the ampoule was transferred into glove
box and opened to take the composite out. The final product was washed in ethanol and
then dried at 90
o
C in glove box.
38
The surface morphology and energy-dispersive X-ray spectrum of the samples
was characterized by a JEOL: JSM-7001 microscope operating at 15 kV. A field
emission transmission electron microscopy (JEOL JEM 2100F) was employed to
obtain the TEM images and scanning transmission electron microscopy (STEM)
images with EDS profiles. Samples were first dispersed in ethanol and then collected
using carbon-film-covered copper grids for analysis. Thermogravimetric analysis was
carried out using a Netzsch STA at a heating rate of 1 ° C min
-1
under N2 atmosphere.
Electrochemical tests were conducted in CR2032 coin cells with Na metal as
counter electrodes and 1 M NaClO4 in dimethyl carbonate (DMC) electrolyte with 10%
fluoroethylene carbonate (FEC) by volume as additive in order to form a strong and
stable solid electrolyte interface (SEI) film. The P@RGO film with a thickness of ~
110 μm was cut into electrodes with a diameter of 16 mm for the electrochemical tests,
and the mass of each electrode is ~ 2.6 mg. All batteries are assembled inside an argon-
filled glovebox with both water and oxygen <0.1 ppm. For the commercial red
phosphorus control sample, phosphorus powder is manually mixed with polyvinylidene
fluoride (PVDF) in 1-methyl-2-pyrrolidinone (NMP) and carbon black in a weight ratio
of 60:10:30. In both cycling stability and rate capability tests, batteries were cycled in
the voltage range of 0.01 to 1.75 V vs. Na/Na
+
at room temperature. All capacities were
calculated based on the total mass of P@RGO composite. Cyclic voltammetry were
collected using a GAMRY Reference 600 test station with a scan rate of 0.1 mV/s and
a scan range from 0 to 3.0 V vs Na/Na
+
.
39
4.3 Results and Discussion
Figure 4.2 (a) SEM image of the P@RGO composite. (b) The enlarged SEM image
and (c) the corresponding EDS mapping profile for phosphorus element in the area
marked with red rectangle in (a). (d) TEM and (e) STEM image of the P@RGO
composite. (f) EDS line-scan profile for phosphorus element of three particles marked
in (e). (g) TEM image of a single phosphorus particle on a RGO sheet. (h) The high-
resolution TEM image of the area marked with blue dashed rectangle in (g), the
graphene layer is marked with light blue dashed line. (i) The bending test of the
40
P@RGO flexible film with the resistance measurement at different bending radius, the
optical images of the flexible film are inserted.
In the scanning electron microscopy (SEM) image of the as-prepared P@RGO
as shown in Figure 4.2a, most RGO sheets have planar surface area larger than 100
μm
2
, and phosphorus particles were deposited onto RGO sheets densely and uniformly.
In Figure 4.2b, the enlarged SEM image of the area marked by red rectangle in Figure
4.2a with the corresponding energy-dispersive X-ray spectroscopy (EDS) element
mapping profile of phosphorus elements shown in Figure 4.2c, the particles on RGO
sheets can be confirmed as phosphorus nano-dots, and the diameter of the nano-dots
varies from hundreds to tens of nanometers. In the transmission electron microscopy
(TEM) image (Figure 4.2d) and the corresponding scanning transmission electron
microscopy (STEM) dark-field image (Figure 4.2e), phosphorus nano-dots can be
easily identified from the brightness contrast. In Figure 4.2f, the EDS line-scan profiles
of three typical particles with different diameters (numbered in Figure 4.2e), diameters
of different phosphorus nano-dots were measured as 79, 161 and 223 nm, which
reached a good agreement with the particle sizes observed under SEM. In the enlarged
TEM image of a single phosphorus particle, as shown in Figure 4.2f, the amorphous
nature of the particle can be observed. In the high-resolution TEM image of the edge
of the particle, as shown in Figure 4.2h, a few graphene layers (marked with light blue
dashed line) can be clearly observed wrinkling around the red phosphorus nano-dot
with good contact.
The synthesized P@RGO composite can be easily fabricated into a highly
flexible and free-standing film by filtration method, as shown in the optical image
41
inserted in figure 4.2i. In order to investigate the mechanical properties of the P@RGO
flexible film, the relative resistance (R/R0) of the film was measured at different
bending radius from flat to 12.3, 8.4. 5.8, 3.4 and 2 mm, and the experimental details
were specified in the Supporting Information. No obvious resistance change (< 10%)
was observed at the bending radius larger than 3.4 mm; a 17% resistance increase was
observed at the bending radius as small as 2 mm. The film was free of breaking during
the entire bending tests, indicating that the RGO film can provide great mechanical
support to the red phosphorus nano-dots against the stresses from the sodiation and de-
sodiation processes.
42
Figure 4.3 (a) Thermogravimetric analysis of the P@RGO composite. (b) Cyclic
voltammetry of the P@RGO anode with a scan rate of 0.1 mV/s between 0 and 3.0 V
vs. Na/Na
+
. (c) Cycling performance of the P@RGO anode at a charge/discharge
current density of 1593.9 mA/g with its potential profiles presented in (d). (e) Rate
performance of the P@RGO anode with its potential profiles presented in (f).
43
Figure 4.3a exhibits the thermogravimetric analysis (TGA) data of the RGO
film and P@RGO composite in nitrogen atmosphere. The P@RGO sample has a sharp
weight loss between 400
o
C and 430
o
C due to the red phosphorus vaporization, and the
weight percentage of phosphorus in P@RGO composite can be calculated as 61.4%.
The electrochemical performance of the P@RGO anode was first tested in cyclic
voltammetry (CV) for the initial three cycles with a voltage window from 0 to 3.0 V vs
Na/Na
+
, as shown in Figure 4.3b. Two peaks at 0.8 and 0 V were observed from the
first sodiation process. The peak located at 0.8 V disappeared in the following cycles
indicating a stable solid electrolyte interphase (SEI) film formation caused by the
decomposition of electrolyte, which contributes to the irreversible capacity at the first
cycle. The major peak at 0 V can be assigned to the sodium intercalation and observed
in the subsequent sodiation process consistently. Only one peak at 0.7 V was observed
during the de-sodiation process for three cycles consistently. A minor decay was
observed at the 0.7 V peak from the first to the second de-sodiation process without
peak position shifting, and the unchanged peak current intensity in the following cycles
implies excellent reversibility of the P@RGO anode.
The cycling performance of our P@RGO anode samples were investigated by
galvanostatic charge/discharge of the electrodes between 0.01 and 1.75 V at a current
density of 1593.9 mA/g, as shown in Figure 4.3c. The P@RGO electrode presents a
75.2% initial efficiency that quickly rises to ~ 99% after 5 cycles, and a high initial
specific discharge capacity of 1611 mAh/g at the first sodiation process indicating a
fully sodiated status. After that, the specific charge capacity decayed from 1074.5 to
930.3 mAh/g from the 2
nd
to the 100
th
cycle with an 86.6% retention, and then was
44
stabilized ~ 940 mAh/g after 100 cycles with a 914 mAh/g capacity at the 300
th
cycle
finally. The corresponding volumetric capacity of the P@RGO film electrode can be
calculated to be around 111 mAh/cm
3
over 300 cycles, based on the fact that each
electrode has a mass of ~ 2.6 mg, 16 mm in diameter and ~ 110 μm in thickness. The
superior cycling performance of P@RGO is comparable to and exhibits an advantage
over most silicon anodes in lithium-ion batteries with a 15-20% weight ratio of
conductive carbon and binder as additives, and also with copper foil as current collector.
Thus, the cycling stability of P@RGO anode was significantly extended by the
precisely controlled nano-sized red phosphorus particles and the architectural RGO
enabled conductive supporting network with superior mechanical properties.
Aside from capacity, voltage profiles actually reveal additional details about the
electrochemical sodiation/de-sodiation process in this anode. The voltage profiles of
the cycling performance test of the P@RGO anode are depicted in Figure 4.3d with
cycle numbers labeled on individual curves. In the first sodiation process, one small
potential plateaus at 0.8 V and the major sodiation potential plateaus at 0.5-0.1 V are
observed. The small plateau is due to the SEI formation in dimethyl carbonate (DMC)
electrolyte with fluoroethylene carbonate (FEC) additive, which reached a good
agreement with the CV curve in Figure 4.3b. The discharge curves present a major
plateau at 0.2-0.1 V at the 2
nd
, 50
th
, 100
th
and 300
th
cycle consistently. On the other
hand, all charge curves of the 1
st
, 2
nd
, 10
th
, 50
th
and 300
th
cycle exhibit a major de-
sodiation plateau at 0.7-0.8 V.
In Figure 4.3e, the rate capability test with charge/discharge current densities
ranging from 159.4 to 47818.3 mA/g, the capacities of the P@RGO anode decays
45
slowly at the first 20 cycles, which reaches a good agreement with the cycling
performance presented in Figure 4.3c. The anode can deliver average capacities of
1165.4, 1105.4, 1081.8, 1039.2, 973.7, 892.9, 755.9, 510.6 and 135.3 mAh/g at the
charge/discharge current density of 159.4, 318.8, 797, 1593.9, 3187.9, 7969.7, 15939.4,
31878.9 and 47818.3 mA/g, and then the capacity was stabilized at ~ 1100 mAh/g after
switching back to 159.4 mA/g after 90 cycles. The representative charge/discharge
curves at different current densities are presented in Figure 4.3f, while the hysteresis
was enlarged attributing to the increasing of current densities. At lower current density
below 10 A/g, the major sodiation and de-sodiation plateaus at 0.2-0.1 V and 0.6-0.7
V can be located. At high current density ~ 47 A/g, the plateaus almost disappeared
and the anode was showing a capacitor-like potential profiles. However, the typical
phosphorus potential plateaus are visible at 31878.9 mA/g current density with an
average specific charge capacity of 510.6 mAh/g, which is a visible improvement over
the red phosphorus anodes reported recently. The superior rate performance presented
above proved that the sodium ion transport kinetics was significantly improved by the
nano-sized phosphorus particles and the rate capability of the whole electrode was
greatly boosted by the excellent electronic conductance of the architectural RGO
network which serves as electron pathway.
46
Figure 4.4 Cross-sectional (a) and planar (d) SEM images of the flexible P@RGO
electrode after 300 cycles with the corresponding EDS element mapping profile of
phosphorus in (b) and (e) and carbon in (c) and (f), respectively.
To gain insight into the accompanying structural change of P@RGO anode
during long cycling, both cross-section and planar SEM images at different
magnifications with corresponding EDS element mapping profiles were collected from
the electrode after 300 deep cycles, as shown in Figure 4.4. Since NaClO4 was
employed as the sodium ion conducting salt in this report, the EDS phosphorus signal
from the electrolyte salt such as NaPF6 can be excluded. In the after-cycling cross-
section SEM as shown in Figure 4.4a, the multi-layer electrode structure was well
maintained. And in Figure 4.4b and 4.4c, both phosphorus and carbon EDS signals can
be detected from the cross-sectional area, indicating that most of the phosphorus was
confined in the designed structure with a negligible mass loss from peeling off from
the RGO sheets during cycling because of the minimized sodiation/de-sodiation
stresses from the amorphous phosphorus nano-dots and excellent mechanical support
47
from the architectural RGO structure. In the after-cycling planar SEM image of the
electrode as shown in Figure 4.4d, the P@RGO network can be clearly observed with
the EDS profiles specifying the distribution of phosphorus and carbon elements.
Although we observed several phosphorus particles in sub-micrometer size, which
were probably formed by the agglomeration between the phosphorus nano-dots during
cycling, the EDS profiles of the two elements reached a good coincidence, which
implies that most of the phosphorus nano-dots survived during the 300 deep cycles and
the architectural P@RGO electrode is favorable for long cycling usage.
4.4 Conclusions
In summary, red phosphorus nano-dots were deposited on highly conductive
RGO sheets densely and uniformly, the phosphorus particle size was controlled
precisely in the range from hundreds to tens of nanometers. In particular, the
phosphorus nano-dots not only minimized the stresses during sodiation/de-sodiation
for long cycle life, but also decreased the sodium ion diffusion length for fast
charging/discharging; the electrochemical performance of the phosphorus anode was
improved by taking the advantage from the RGO network which served as electron
pathway and provided excellent mechanical support against the volume variation of
phosphorus particles. The fabricated highly flexible P@RGO electrode can provide a
1211 mAh/g charge capacity toward sodium-ion at the initial cycles and retain 914
mAh/g after 300 cycles with a 1593.9 mA/g charge/discharge current density; the anode
also achieved a 510.6 mAh/g capacity at ~ 31.8 A/g current density with typical
phosphorus potential plateaus and a 135.3 mAh/g capacity at ~ 47.8 A/g current density,
demonstrating the best rate performance of flexible red phosphorus anodes for Na-ion
48
batteries reported in the literature to date. The superior cycling and rate performance,
combined with the excellent mechanical properties of the P@RGO electrode,
represents a suitable strategy for the development of inexpensive and versatile
techniques for flexible and wearable sodium-based energy storage applications.
Chapter 5: Single-Step Flash-Heat-Treatment Synthesized
Flame-Retardant Red Phosphorus/Graphene Composite as
Flexible Anode for Sodium-Ion Batteries
5.1 Introduction
Red phosphorus (RP) is the most cost-effective allotrope among all kinds of P with
reasonably good chemical stability for battery studies, and the adoption of RP can
potentially enhance the flame retardancy of the electrodes since it has been widely used
as an effective flame retardant additive for decades [61]. However, the micron-sized
RP particles have been proven to be unsuitable for stable Na-ion storage due to the poor
electronic conductivity and huge volume expansion during the sodiation process [58].
Although encouraging improvement was achieved, multifaceted issues such as the
nonuniform particle size and poor contact between the RP active material and the
conductive platforms ascribed to simple mechanical mixing still hamper the
electrochemical performance optimization of the RP anodes. Traditional vaporization-
condensation methods have been adopted for RP anodes recently, which led to
impressive performance enhancement over the micron-sized RP particles by taking
advantages of the uniform RP particle size and good contact between the active
materials and the conductive scaffold. However, the sophisticated preparation process
49
of such synthesis methods involving quartz ampoule sealing makes it encounter
challenges originated from the productivity and scalability.
Figure 5.1 Schematic diagrams of the RP/rGO composite synthesis process at (a) pre-
heat-treatment, (b) flash-heat-treatment, and (c) cooling-down steps, respectively. (d)
Schematic diagram of the nano-structure of the resulted RP/rGO composite.
Herein, we report a facile flash-heat-treatment synthesis method to grow nano-
sized RP onto the surface of reduced graphene oxide (rGO) sheets and into the void
50
space between rGO layers, and the RP deposition and the GO reduction were completed
in a single-step heat-treatment simultaneously. The synthesis method of the RP/rGO
composite is schematically described in Figure 5.1. In Figure 5.1a, the RP and GO
precursors were placed in a ceramic boat following a RP/GO/RP three-layer structure
with a ceramic cover in a tube furnace with Ar/H2 atmosphere. The boat was placed on
the side of the heating zone initially. After heating the furnace to 500
o
C, the boat was
moved into the heating zone, as shown in Figure 5.1b. Once the P condensation was
visualized on the quartz tube at the downstream of gas flow, the boat was moved back
to the original position immediately, as described in Figure 5.1c. Then the boat was
maintain at 300
o
C for 6 hours to convert the white P to red P with a cooling-down
process. The structure of the resulted RP/rGO composite is schematically described in
Figure 5.1d, which combines multiple advantages: 1. Nano-sized RP can shorten the
ion diffusion length and thus boost both ionic and electronic kinetics of the anode; 2.
The RP filled into the gaps between the rGO layers can be protected by the rGO sheets
conformally against the volume variation during the Na ion intercalation and extraction;
3. The rGO network can serve as the electron pathway and thus increases the electronic
conductivity of the RP anode. Furthermore, the flash-heat-treatment synthesis process
reported is more simple and cost-effective than the traditional ball-milling and
vaporization-condensation methods, and thus enabled large throughput manufacturing
of the RP anodes.
5.2 Experimental
Synthesis of the RP/rGO composite: RP precursor was dried at 90
o
C to remove the
moisture and meshed with a 30 µ m mesh after drying. The RP and GO powder
51
precursors were placed in a ceramic boat with a RP/GO/RP three-layer structure with
a ceramic cover on the boat. The boat with chemicals was loaded into a tube furnace
under argon flow with a mixture of 5% hydrogen (Ar/H2). The boat was first placed on
the side of the heating zone. After heating the furnace to 500
o
C, the boat was moved
into the heating zone, and once the P condensation is visualized on the inner surface of
the quartz tube at the downstream of the gas flow (~ 1 minutes), the boat was moved
back to the original position immediately. Then the temperature of the boat was
maintain at 300
o
C for 6 hours to convert the white P to red P. After a cooling-down
process, the resulted RP/rGO composite was transferred into an Ar-filled glovebox, and
washed with methanol and dried. The preparation of the rGO and P-treated rGO control
samples and the flexible film electrode fabrication are specified in the Supporting
Information.
The surface morphology and energy-dispersive X-ray spectrum (EDS) of the
samples was characterized by a JEOL: JSM-7001 microscope operating at 15 kV, the
samples were assembled onto the specimen by carbon tapes. A field emission
transmission electron microscopy (JEOL JEM 2100F, 200 kV) was employed to obtain
the TEM images and scanning transmission electron microscopy (STEM) images with
EDS profiles. Samples were first dispersed in ethanol through ultra-sonication and then
collected using carbon-film-covered copper grids for analysis. Thermogravimetric
analysis (TGA) was carried out using a Netzsch STA at a heating rate of 1 ° C min
-1
under N2 atmosphere with the temperature range from room temperature to 800
o
C.
Raman spectra was taken using a Renishaw InVia spectrometer with a 532 nm layer
(10 μW) focused through a 100× objective lens. Raman spectra were collected at room
52
temperature under ambient conditions, glass slides were used as substrates to carry
powder samples.
Electrochemical tests were conducted in CR2032 coin cells with Na metal as
counter electrodes and 1 M NaClO4 in dimethyl carbonate (DMC) electrolyte with 10%
fluoroethylene carbonate (FEC) by volume as additive in order to form a strong and
stable solid electrolyte interface (SEI) film in the first cycle. The RP/rGO film with a
thickness of ~ 110 μm was cut into electrodes with a diameter of 16 mm for the
electrochemical tests, and the mass of each electrode is ~ 1.2 mg. All batteries are
assembled inside an argon-filled glovebox with both water and oxygen <0.1 ppm. For
the commercial RP control anode samples mixed with rGO and carbon black (CB), RP
powder is mixed with polyvinylidene fluoride (PVDF) in 1-methyl-2-pyrrolidinone
(NMP) and rGO or carbon black (CB) in a weight ratio of 50:10:40 by high energy ball
milling for 1 hour with argon filled in the ball mill jar. The resulted slurry was casted
onto an Al foil and dried at 90
o
C in air overnight, and then punched into electrodes with
the size of CR2032 type cell. The active material mass loading on the Al foil current
collector was ~ 1.5 mg/cm
2
. In the cycling stability and rate capability tests, batteries
were cycled in the voltage range of 0.01 to 1.75 V vs. Na/Na
+
at room temperature. All
capacities and current densities were calculated based on the mass of RP only unless
specified differently.
53
5.3 Results and Discussion
Figure 5.2 (a, b) SEM images of the RP/rGO composite at × 1500 and × 4000
magnifications. (c, d) TEM and STEM images of a single piece of RP/rGO flake. (e)
Enlarged STEM image of the RP/rGO composite with the corresponding EDS mapping
profile of the P element in (f). The RP filled into the void space between the rGO layers
was marked by blue lines in (e).
The as-produced RP/rGO composite has both RP nano-dots deposited onto rGO
surface and chunks of nano-sized RP deposited into the void space between rGO layers,
as described below according to the scanning electron microscope (SEM), transmission
electron microscope and scanning transmission electron microscope (STEM) images.
In Figure 5.2a and 5.2b, both SEM images of the RP/rGO composite at different
magnifications show nano-sized P particles were grown on the rGO sheets in a dense
and uniform way. Figure 5.2c and 5.2d show the TEM and STEM images of one piece
54
of RP/rGO flake, which confirms the presence of P nano-particles with morphology
consistent with the SEM images. The dense growth of RP particles can be easily
identified from the dark-field STEM image shown in Figure 5.2d from the sharp
brightness contrast difference between P and rGO. More details of the architectural
structure of the RP/rGO composite can be found in the enlarged STEM image shown
in Figure 5.2e, with the corresponding energy-dispersive X-ray spectrum (EDS)
mapping profile of the P element exhibited in Figure 5.2f: besides the nano-sized P
particles deposited on the surface, considerable amount of P was filled into the void
space between rGO layers through gaps and wrinkle tunnels (marked by blue dash lines
in Figure 5.2e), indicating conformal protection from the robust rGO layers. The
synthesized RP/rGO composite has a sharp weight loss between 400
o
C and 450
o
C due
to the RP vaporization, and the weight percentage of RP in the composite was
calculated as 57.9%.
55
Figure 5.3 (a) Raman spectra of the RP/rGO composite and commercial RP powder.
(b) XPS spectrum of the synthesized RP/rGO composite and pristine GO powder. (c)
The high-resolution C 1s XPS spectrum of the synthesized RP/rGO composite and
pristine GO powder. (d) Optical images of a flexible RP/rGO film was wrapped on a
glass rod with a diameter of 5 mm, and a punched electrode was bent with a pair of
tweezers.
The Raman spectrum of the RP/rGO composite and commercial RP powder is
exhibited in Figure 5.3a. Besides the D and G band peaks at 1331 and 1577 cm
-1
, the
spectrum of the RP/rGO composite contains a small broad peak of 2D signal, indicating
a multi-layered nature of the rGO sheets. Compared to the commercial RP sample, the
RP signal was also detected from the Raman spectrum of the synthesized RP/rGO
56
composite in the range from 300 to 500 cm
-1
, which indicates that the P domain
observed in Figure 5.2 is in the form of RP. The GO precursor shows a major (002)
peak at ~ 10
o
, which disappeared after the flash-heat-treatment process, indicating the
successful reduction of GO. Figure 5.3b presents the XPS spectrum of the RP/rGO
composite and the pristine GO which has been used as the precursor of the RP/rGO
synthesis. Both 2p and 2s peaks of the P element can be detected from the spectrum of
the RP/rGO composite, and the intensity of the O 1s peak was greatly weakened
compared to the GO sample, indicating most of the functional groups containing the O
element were eliminated in the heat-treatment process. More details about the reduction
of GO are exhibited in Figure 5.3c, which shows the C 1s high-resolution XPS
spectrum of the two samples. The spectrum pocket of the pristine GO sample can be
fitted into three Gaussian-Lorentzian peaks: C=C/C-C bond at 284.6 eV, C-O bond at
286.3 eV, and O-C=O bond at 287.2 eV; and for the RP/rGO composite, C=C/C-C and
C-O bond at 284.5 and 285.6 eV can be identified. In contrast to the GO sample, the
C-O bond peak intensity of the RP/rGO composite was greatly reduced, and the O-
C=O peak was almost vanished, illustrating the GO was reduced thermally during the
heat-treatment synthesis process of the RP/rGO composite in Ar/H2 atmosphere. No
obvious C-P bond signal was detected in the Raman and XPS spectrum of the RP/rGO
composite, implying that the heating temperature might be too low and/or heating time
might be too short to achieve a relatively high P-doping concentration in graphene
during the synthesis process. The synthesized RP/rGO composite can be easily
fabricated into highly flexible free-standing films through a filtration method as shown
in the optical images in Figure 5.3d with details specified in the Supporting Information,
57
and the film was free of visible fracture during the bending process on a glass rod with
a diameter of 5 mm. The film was punched into electrodes with a diameter of 17 mm
for the electrochemical performance tests described below.
Figure 5.4 (a) Cycling performance of the fabricated RP/rGO flexible film anode, the
commercial RP mixed with rGO, and the commercial RP mixed with carbon black (CB)
at a charge/discharge current density of 1 A/g with its potential profiles exhibited in
(b). Rate performance of the RP/rGO anode with its potential profiles exhibited in (d).
The electrochemical performance of the RP/rGO composite and the commercial
RP control sample was first investigated by galvanostatic charge/discharge cycling
experiments between 0.01 and 1.75 V at a current density of 1 A/g with Na metal as
counter electrodes, as exhibited in Figure 5.4a. Since no obvious P heteroatom-doping
was detected in rGO and the rGO network works as current collector in this RP/rGO
58
free-standing electrode, all current densities and capacities were calculated based on
the mass of RP only unless specified differently. Specific charge capacities of 2197 and
1883 mAh/g at the first and the second cycle are observed, respectively. Afterward, the
capacity increased slightly from 1705 to 1823 mAh/g from the 3rd to the 50th cycle,
and then was stabilized ~ 1550 mAh/g from the 50th to the 200th cycle. In contrast, the
commercial RP mixed with rGO and carbon black (CB) obtained similar initial
capacities but suffered from a rapid capacity fading during the first 30 and 120 cycles.
Thus, the cycling stability of RP/rGO anode was significantly extended by the nano-
sized RP particles deposited onto the surface of rGO sheets and into the void space
between rGO layers, and the architectural rGO enabled conductive supporting network
with superior mechanical properties.
The RP/rGO film electrode presented a 73.2% efficiency at the first cycle and that
quickly raised to ∼98% after 5 cycles. Similar to silicon as anode for lithium-ion
batteries, the relatively low initial efficiency may come from the Na consumption
during the solid electrolyte interphase (SEI) formation on the surface of those RP
particles deposited on graphene sheets without conformal coating protection from
graphene. The low initial coulombic efficiency can be an obstacle for the practical
applications of the RP anodes. However, methods such as pre-sodiation and artificial
SEI film coating can be potentially adopted to solve the above issue.
Aside from capacity, voltage profiles actually reveal additional details about the
electrochemical sodiation/de-sodiation process in this anode. The voltage profiles of
the cycling performance test of the RP/rGO anode are depicted in Figure 5.4b with
cycle numbers labeled on the side. In the first sodiation process, a small potential
59
plateau at 0.7 V and the major sodiation plateau between 0.5 and 0.1 V are observed.
The minor plateau at 0.7 V is due to the SEI film formation caused by the
decomposition of electrolyte, which contributes to the irreversible capacity at the first
cycle. The disappearance of this plateau in the following cycles indicates a stable SEI
film formation. The discharge curves present a major plateau at 0.4-0.1 V at the 2nd,
50th and 200th cycle consistently. All charge curves exhibit a major de-sodiation
plateau at 0.2-0.7 V.
The rate capability test results are exhibited in Figure 5.4c with charge/discharge
current densities ranging from 1 to 8 A/g. The anode can deliver average capacities of
1786, 1597, 1324 and 679 mAh/g at the current densities of 1, 2, 4 and 6 A/g,
respectively; and a capacity of ~ 10 mAh/g is observed at the current density of 8 A/g.
Then the capacity was maintained at ~ 1640 mAh/g after switching the current density
back to 1 A/g after 90 cycles. The representative charge/discharge curves at different
current densities are presented in Figure 5.4d. The hysteresis between the charge and
discharge curves was enlarged attributing to the increasing of current densities, and the
major sodiation and de-sodiation plateaus at 0.4-0.1 V and 0.2-0.7 V can be located for
all curves. The superior rate performance presented above implies that the sodium ion
transport kinetics was significantly enhanced by the nano-sized RP, and the rate
capability of the whole electrode was greatly boosted by the excellent electronic
conductance of the architectural rGO network which serves as electron pathways.
60
Figure 5.5 (a, b) TEM and STEM images of the flexible RP/rGO anode after 200 cycles,
with the corresponding EDS spectrum (c) and elemental mapping profiles of
phosphorus in (d), sodium in (e), and carbon in (f), respectively.
To gain insight into the accompanying structural change of the RP/rGO anode
during long cycling, the TEM image with corresponding EDS element mapping
profiles were collected from the electrode after 200 cycles at a fully sodiated status, as
shown in Figure 5.5. Before analysis, the post-cycling electrode was immersed in
electrolyte solvent overnight to remove the Na salt. Since NaClO4 was employed as the
sodium ion conducting salt in this report, the signal of the P element from the electrolyte
salt such as NaPF6 can be excluded. In the post-cycling TEM and STEM images shown
in Figure 5.5a and 5.5b, a single piece of the RP/rGO composite was observed. The
corresponding EDS elemental spectrum is presented in Figure 5.5c, and the elemental
mapping profiles are presented in Figure 5.5d, 5.5e and 5.5f for carbon, sodium, and
61
phosphorus, respectively. A RP flake in micrometer size is attached on the graphene
sheet according to the mapping profile of the P element, indicating the agglomeration
between RP nano-sized particles on the surface of rGO sheets during cycling. The
mapping profiles of three elements reached a good coincidence, which implies that
most of the RP active material survived during the 200 deep cycles and the architectural
RP/rGO structure is favorable for long cycling usage.
Figure 5.6 Combustion tests of an RP/rGO flexible film (a) and an rGO film (b) at t=0,
0.5, 2, 5 and 10 s, respectively. t=0 s is defined as the moment when the flame on the
films can be clearly visualized.
Although the batteries have achieved great successes in both academia and
industry, safety obstacles associated with the highly flammable liquid organic
electrolytes remain a perennial issue [66-68]. A lot of efforts have been devoted to
62
solving this safety problem, such as developing non-flammable electrolyte [62-64],
adding flame-retardant additives into electrolyte [65-67], and separator engineering
[68-70]. The flame retardancy of the electrodes is one of the key factors to improve the
battery safety by further preventing the flame spreading after the ignition of the
electrolyte. In order to explore the flame retardancy of the RP/rGO composite, the
RP/rGO film was investigated in the combustion test with an rGO film as a control
sample. Both films were immersed in the electrolyte solvent for 3 hours before the tests.
The optical image of the tests at t=0, 0.5, 2, 5 and 10 s are presented in Figure 5.6a and
5.6b for the RP/rGO and rGO films, respectively. t=0 s is defined as the moment when
the flame on the films can be clearly visualized. The flame was fully developed at 0.5
s with the red-hot edges formed on both films due to the combustion of electrolyte.
Afterward, the flame attenuated due to the exhaustion of the electrolyte on the surface
of and absorbed by the films. From 5 to 10 s, most of the rGO film was burned out with
the red-hot edge moving inward. On the other hand, the red-hot edge of the RP/rGO
film vanished at ~ 5 s, and only a small part of the film was burned out from 5 to 10 s.
The videos of the combustion tests discussed above can be found in the Supporting
Information. The RP/rGO film gained excellent flame retardancy as a result of the
phosphoric acid derivative generation which can isolate the burning material from
oxygen and catalyze the char layer formation on the material surface, and thus further
prevents the flame formation [71,72].
5.4 Conclusions
In summary, we have developed a single-step flash-heat-treatment synthesis method to
deposit nano-sized RP onto the rGO sheets surface and into the void space between
63
rGO layers, and the RP growth and the GO reduction were completed within a single-
step flash-heat-treatment process. In particular, the nano-sized RP not only minimized
the stresses during sodiation/de-sodiation for long cycle life, but also decreased the
sodium ion diffusion length for enhanced kinetics; the electrochemical performance of
the RP anode was boosted by taking advantages of the rGO network which served as
electron pathways and provided excellent mechanical support to the RP filled into the
void space between rGO layers against the volume variation during cycling. The
resulted RP/rGO flexible anode achieved 1786, 1597, 1324 and 679 mAh/g specific
charge capacities at 1, 2, 4 and 6 A/g charge/discharge current densities in the rate
performance test, and an average capacity of 1625 mAh/g during 200 deep cycles at 1
A/g current density, which would convert to 941 mAh/g if calculated based on the total
mass of the free-standing RP/rGO film electrode. Moreover, the RP/rGO film also
obtained excellent flame retardancy by taking advantage of the RP ingredient. The
superior electrochemical performance of the RP/rGO flexible film electrode combined
with the improved flame retardancy, represents a suitable strategy for the inexpensive,
safe, and wearable Na-based energy storage applications.
Chapter 6: Room-Temperature Pressure Synthesis of Layered
Black Phosphorus-Graphene Composite for Sodium-Ion Battery
Anodes
6.1 Introduction
The common white phosphorus form, however, is toxic and unstable,
suggesting that the other two allotropes, black phosphorus (BP) and red phosphorus
64
(RP), should be tested and developed for sodium ion storage. Similar to silicon anodes
in lithium ion batteries, due to the intrinsic insulating nature and the huge volume
expansion during the sodiation process, RP were usually bundled with conductive
matrix such as carbon nanotubes, graphene and other carbonaceous materials.[56-60]
Although RP offers lower material cost, the higher electrical conductivity (~300 S/m)
of BP promises better rate performance in sodium-ion electrochemical applications,
and less conductive additives would be needed for BP anodes.[52-54] Most BP-based
energy storage work to date has used BP exfoliated from bulk crystals, an expensive,
complex and poorly-scalable production method whose intrinsic challenges have so far
offset the nominal advantages of BP. Recently, synthesis of BP through application of
pressure at room temperature has been reported, including complete RP to BP
conversion leading to a BP thin film on a flexible substrate for electronic and optical
device applications.[73] However, reports of two-dimensional anisotropic volumetric
expansion of BP during the sodiation reaction imply that a successful energy storage
anode will require complementing BP with other conductive two-dimensional
materials, such as reduced graphene oxide (rGO).
Figure 6.1 Schematic description of the BP/rGO synthesis.
65
In this study, we demonstrate a low-cost and scalable synthesis of BP/rGO
layered structure electrodes under a pressure of 8 GPa at room temperature, as
illustrated in Figure 6.1. The RP/rGO precursor was first synthesized using a flash-
heat-treatment method. Briefly, the commercial RP and GO powder were arranged in
a three-layer structure inside a ceramic boat with cover in a tube furnace under Ar/H 2
flow. The 2-minute flash-heat-treatment was control by moving the boat into and out
of the hot zone. The RP/rGO precursor was assembled into a film through filtration to
form the layered structure. Samples of RP/rGO film were transferred onto alumina foil
current collectors and 10 to 20 discs of RP/rGO/Al were stacked together for the RP to
BP conversion in a multi-anvil cell, as illustrated in Figure 6.1d. The assembly was
held at 8 GPa for 4 hours and then slowly decompressed over 10 hours to ensure
complete RP to BP conversion and minimal cracking during expansion. The
synthesized BP/rGO/Al discs were directly employed as electrodes in the
electrochemical tests, without any carbon black or polymer binder additives. In the
BP/rGO layered structure, the high conductivity of both BP and graphene facilitates
sodium ion transport and thereby accelerates the electrochemical reaction rate of the
whole battery system. The excellent mechanical properties of the graphene phase
accommodate the volume differences between BP and Na3P, stabilizing the
nanostructure. The simple two-step synthesis of these layered BP/rGO electrodes,
without carbon black or polymer binder, is convenient and reliable method of preparing
phosphorus anodes, offering excellent possibility for scalable production of composites
for sodium-ion batteries.
66
6.2 Methods and Experimental
The RP/rGO precursor was assembled to a film through filtration, and then
transferred onto 16 mm diameter alumina foil discs that serve as current collectors. The
loading mass of the electrode material is ~ 2 mg. Then, 10 to 20 of these RP/rGO/Al
discs were stacked together for RP to BP conversion. The discs were placed at the
center of a 16 mm diameter hole drilled face-to-face through a 25 mm edge-length
chromium-doped magnesium oxide octahedron, with the ends of the hole filled cell by
two pryrophyllite rods. The assembly was placed in a 6-8 Kawai-type multi-anvil
apparatus with pyrophyllite gaskets and 18 mm truncation edge length on the anvils
and loaded in a 1000-ton hydraulic press. Using a room-temperature pressure
calibration for this assembly based on the electrical resistivity transitions of Bi, the
pressure was increased to 8 GPa over 10 hours, maintained for 6 hours, and then slowly
released to ambient pressure over another 10 hours. The synthesized BP/rGO/Al discs
were recovered and then directly employed as electrodes in the electrochemical tests.
The surface morphology and energy-dispersive X-ray maps of the samples were
obtained with a JEOL JSM-7001 scanning electron microscope operating at 15 kV.
Power X-ray diffraction (XRD) of materials at various stages of preparation and testing
were obtained with Rigaku Ultima IV powder/thin-film diffractometer with Cu Kα
radiation. Raman spectra were obtained with a Renishaw Raman spectrometer and 532
nm excitation laser and a laser spot size of ~ 1 μm. A field emission transmission
electron microscopy (JEOL JEM 2100F) was employed to obtain the TEM images and
scanning transmission electron microscopy (STEM) images with EDS profiles.
Samples were first dispersed in ethanol and then collected using carbon-film-covered
67
copper grids for analysis. Thermogravimetric analysis (TGA) was carried out using a
Netzsch STA at a heating rate of 1 ° C min
-1
under N2 atmosphere.
Electrochemical tests were conducted in CR2032 coin cells with Na metal as
counter electrodes and 1 M NaClO4 in dimethyl carbonate (DMC) electrolyte with 10%
fluoroethylene carbonate (FEC) by volume as additive in order to form a strong and
stable solid electrolyte interface (SEI) film. The electrodes were immersed in
electrolyte for 24 hours before the battery cell assembling. All battery cells are
assembled inside an argon-filled glovebox with both water and oxygen <0.1 ppm. In
both cycling stability and rate capability tests, batteries were cycled in the voltage range
of 0.01 to 1.5 V vs. Na/Na
+
at room temperature. All capacities were calculated based
on the total mass of BP/rGO electrode material. In the cycling stability test of the
BP/rGO anode at 40 A/g current density, the first sodiation process was performed with
1 A/g current density.
6.3 Results and Discussion
68
Figure 6.2 (a) Planar and (c) cross-section SEM images of the precursor RP/rGO film;
the same fields of view are shown EDS elemental maps of phosphorus concentration
in (b) and (d). (e) Planar and (f) cross-sectional SEM images of the pressure-
synthesized BP/rGO film. (g) XRD patterns of the GO, RP/rGO precursor, as-prepared
and post-cycling BP/rGO samples. (h) The Raman spectra of RP/rGO, as-prepared and
post-cycling BP/rGO samples.
Initial characterization: After the pressure synthesis, the BP/rGO layered
composite and its precursor were characterized by scanning electron microscopy (SEM)
with energy-dispersive X-ray spectroscopy (EDS), X-ray diffraction (XRD) and
Raman spectroscopy. The planar and cross-sectional SEM images of the precursor
RP/rGO film and as-prepared BP/rGO film are shown alongside EDS maps of
phosphorus K X-ray intensity in Figure 6.2a to 6.2f. In Figure 6.2a and 6.2b, red
phosphorus particles with dimensions varying from hundreds of nm to several m can
be clearly visualized on the graphene sheet. The cross-sectional SEM image and
phosphorus X-ray map of the RP/rGO film in Figure 6.2c and 6.2d clearly show the
layered structure of the red phosphorus between graphene sheets. In the planar and
cross-section SEM images (Figure 6.2e and 6.2f) of the pressure-synthesized BP/rGO
anode, the surface is transformed to a uniform flat plane, no discrete phosphorus
particles are visible, and the layered structure is notably denser than in the starting
material shown in Figure 6.2c. The overall thickness of the film decreased upon
pressing and recovery from ~ 35 μm to ~ 18 μm. The XRD patterns (Cu Kα radiation)
of the GO, the RP/rGO precursor, a sample of the as-prepared BP/rGO, and a sample
of the tested BP/rGO after a cyclic performance test are all given in Figure 6.2g. The
69
GO control sample exhibits a peak at 2 = 12° that disappears and is replaced by a
broad peak at ~ 24° in the RP/rGO precursor, indicating that GO is reduced to graphene
during the flash-heat-treatment synthesis process. On the other hand, the as-prepared
BP/rGO sample shows four characteristic peaks of BP, which can be assigned to the
lattice plane reflections (020), (040), (060) and (080). The Raman spectra (532 nm laser
radiation) of the RP/rGO precursor and of the BP/rGO samples both as-prepared and
after cyclic testing are shown in Figure 6.2h. After pressure synthesis, the broad RP
Raman band from 300 to 500 cm
-1
disappeared entirely, whereas the as-prepared
BP/rGO material clearly shows the characteristic peaks of BP at 364, 440 and 469 cm
-
1
, corresponding to A
1
g
, B
2
g
and A
2
g
lattice vibrational modes. The disappearance of the
broad RP Raman band and the appearance of the BP characteristic peaks suggest most
of the RP precursor has been converted to BP.
70
Figure 6.3 (a) TEM and (b) STEM image of the BP/rGO composite, with the
corresponding maps of X-ray intensity by EDS for (c) phosphorus and (d) carbon. (e)
Enlarged TEM image and (f) high-resolution TEM image of the area marked with light
blue rectangle in (e) with the corresponding FFT images shown in the insets.
The transmission electron microscopy (TEM) image and the corresponding
scanning transmission electron microscopy (STEM) dark-field image are shown
alongside maps of the EDS intensity of the phosphorus and carbon X-rays are given in
Figure 6.3a to 6.3d. The images show that the graphene layers wrap around the BP,
indicating good contact between the active material and supporting network in the
designed BP/graphene layered structure. In the enlarged and high-resolution TEM
images (Figure 6.3e and 6.3f), the BP crystal lattice spacing can be clearly observed
forming coherent crystallites with dimensions up to ~ 100 nm, but the poly-crystalline
nature of the material at the scale of the whole images is confirmed by reflections from
multiple crystallites, approaching Bragg rings, in the fast Fourier transform (FFT)
images shown as insets in Figure 6.3e and 6.3f.
71
Figure 6.4 (a) Cycling performance of the BP/rGO anodes at charge and
discharge current densities of 1 and 40 A/g in black and red color, respectively; specific
capacity is plotted as solid curves against the left-hand axis, whereas Coulombic
efficiency is plotted as open circles against the right-hand axis. (b) Charging and
discharging potential profiles at 1 A/g current density (black curve in (a)), shown for
selected cycle numbers. (c) Rate performance of the BP/rGO anode for a series of tests
with five cycles at each value of current density. (d) Charging and discharging potential
72
profiles for the rate tests presented in (c), color-coded by current density. (e)
Electrochemical impedance spectra for RP/rGO and BP/rGO anodes.
Electrochemical tests: The electrochemical performance of layered BP/rGO
anodes were tested in coin cells with sodium foils as counter electrodes. Because the
synthesized BP/rGO anodes are free of carbon black and polymer binder, all capacities
and current densities are calculated based on the total mass of the BP/rGO electrode
material in this work, unless specified otherwise. First, the cycling performance of our
BP/rGO anodes was investigated by galvanostatic charge and discharge between 0.01
and 1.5 V at current densities of 1 and 40 A/g, as shown in Figure 6.4a. At 1 A/g current
density, the BP/rGO electrode presented a high initial coulombic efficiency of 89.5%
and an initial specific discharge capacity of 1680.3 mAh/g for the first reaction cycle
with sodium. After that, a charge capacity of 1503.9 mAh/g was observed for the first
cycle and capacity decayed from 1474.8 to 1364.3 mAh/g over the first 100 cycles
(92.5% retention). Charge capacity stabilized at ~ 1250 mAh/g after 500 cycles. On the
other hand, at 40 A/g current density, the anode presented a slightly lower initial
efficiency of 86.6%. Charge capacities of 851.9 and 791.6 mAh/g were observed in the
first and second cycles and capacity stabilized at ~ 640 mAh/g after 500 cycles.
However, compare to current density of 1 A/g, cycling performance at 40 A/g shows a
larger fluctuation. We believe higher current rate would introduce larger over potential,
which results in larger efficiency fluctuation in sequence. In particular, at higher current
rate, due to the kinetic limitation and hysteresis, the Na ions can get trapped in the host
materials, and the sodiation/de-sodiation reactions may become partially irreversible
with new SEI film formation at other locations.
73
Capacity is a key performance indicator, but voltage profiles reveal additional
details about the electrochemical sodiation and de-sodiation reactions in this anode.
The voltage profiles of the cycling performance test are depicted in Figure 6.4b with
cycle numbers labeled on individual curves. The first sodiation process displays a small
potential plateau at 0.7-0.9 V and a major sodiation potential plateau at 0.5-0.2 V. The
small plateau is due to SEI film formation in the electrolyte, dimethyl carbonate (DMC)
with fluoroethylene carbonate (FEC) additive, which is responsible for the irreversible
capacity. The discharge curves consistently present major plateau at 0.4-0.2 V for the
2
nd
, 100
th
and 500
th
cycle. On the other hand, all the charge curves (1
st
, 2
nd
, 100
th
, and
500
th
cycle) exhibit major de-sodiation plateau at 0.4-0.6 V.
A second anode was tested in a rate capability protocol with five
charge/discharge cycles at each current density, ranging from 0.1 to 60 A/g. The data
in Figure 6.4c show that the anode can deliver average capacities of 1460.1, 1401.2,
1377.6, 1339.7, 1277.8, 1123.78, 720.8 and 17.3 mAh/g at current densities of 0.1, 0.5,
1, 5, 10, 20, 40 and 60 A/g, respectively. Charge capacity stabilized after 40 cycles at
~ 1400 mAh/g after switching back to 0.1 A/g. Representative voltage profiles for
charge and discharge at various current densities are given in Figure 6.3e. The major
sodiation and de-sodiation plateaus at 0.2-0.1 V and 0.6-0.7 V are evident for current
density from 0.1 to 20 A/g. Although hysteresis between charge and discharge curves
becomes significant at 40 A/g current density, the typical phosphorus potential plateaus
persist and the anode can still deliver an adequate average charge capacity of 720.8
mAh/g, indicating the ultra-fast electronic and ionic transport of our BP/rGO layered
composite anode. However, the highest current density tested, 60 A/g, exceeds the
74
transport rates achievable and the anode can only deliver an average capacity of 17.3
mAh/g with a capacitor-like potential profile without any plateau.
In order to investigate the origin of the observed rate capability, we measured
electrochemical impedance spectrum (EIS) curves for both RP/rGO and BP/rGO
anodes, as seen in Figure 6.4f. The Nyquist plots of both anodes yield a depressed
semicircle in the high-to-medium frequency range (corresponding to the charge transfer
impedance at the electrolyte/electrode interfaces) followed by a straight line at low
frequency (corresponding to the bulk diffusion impedance in the composites). Both
anodes present very small contact resistance at high frequency; the charge-transfer
impedance value is 38 for the BP/rGO anode and 62 for the RP/rGO anode. The
much smaller charge transfer impedance of BP/rGO arises from the higher electrical
conductivity of BP. For the pressure-synthesized BP/rGO anode, the slope of the low
frequency straight line is much higher than that of the RP/rGO anode, indicating higher
sodium ion diffusivity in BP.
.
75
Figure 6.5 TEM results for the post-cycling BP/rGO anode: (a) TEM image with FFT
inset; (b) STEM image; (c) EDS map of phosphorus X-ray intensity; (d) EDS map of
carbon X-ray intensity.
Post-cycling characterization: To gain insight into phase and structural changes
in the BP/rGO anode after repeated charge and discharge cycling, a fully de-sodiated
anode that experienced 500 cycles of charge and discharge at 1 A/g current density was
prepared for a second round of TEM and X-ray map analysis, as shown in Figure 6.5.
Since NaClO4 was employed as the ion conducting salt in this work, none of the
phosphorus signal in the map is derived from any electrolyte salt such as NaPF6. The
graphene sheets are apparent in Figure 6.5a, 6.5b and 6.5d, whereas the active
phosphorus active is evident in the EDS mapping profile in Figure 6.5c. It is clear that
the graphene structure was maintained and remained wrapped around the active
material after 500 deep cycles. This suggests that, as anticipated in the composite
design, most of the phosphorus remained confined in the graphene network with
negligible mass loss due to volume variation or peeling off. The BP/graphene layered
structure appears sufficiently robust and favorable for usage over long cycles.
Interestingly, most active material in the anode developed porosity, as shown in the
TEM and STEM images (Figure 6.5a and 6.5b). The FFT image in the inset to Figure
6.5a displays a prominent broad ring due to amorphous scattering, indicating that most
of the polycrystalline BP was converted to amorphous phosphorus during the
sodiation/de-sodiation cycling. This result agrees well with the XRD and Raman
patterns of the post-cycling BP/rGO anode (Figure 6.2f and 6.2g), from which all the
characteristic peaks of BP disappeared after cycling. The amorphization of the BP may
76
result from the high concentration of sodium when the sample if fully sodiated, which
results in a transition from intercalation to an alloying reaction and the breaking of P-
P bonds in BP.
6.4 Conclusions
In summary, BP/rGO layered composite were synthesized by the application of
pressure at room temperature and the resulting anodes present excellent cycling
stability and rate capability. In particular, the high electronic conductivity of the active
BP material and the graphene network facilitates ion transfer kinetics of BP/rGO
anodes for fast charging/discharging and the graphene network provides robust
mechanical support despite volume changes in the phosphorus, leading to stable
electrochemical performance. After 500 deep cycles, the synthesized BP/rGO
electrodes continue to provide ~ 1250 mAh/g charge capacity at 1 A/g charge/discharge
current density and ~ 640 mAh/g capacity at 40 A/g current density. The anode
delivered average capacities of 1460.1, 1401.2, 1377.6, 1339.7, 1277.8, 1123.78 and
720.8 mAh/g at current densities of 0.1, 0.5, 1, 5, 10, 20 and 40 A/g, demonstrating the
best high-rate phosphorus anode performance reported in the sodium-ion literature to
date. Only at 60 A/g did charging rate exceed the kinetic capability of the anode. The
superior cycling and rate performance and straightforward pressure synthesis of this
carbon-black-free and binder-free electrode material represents a suitable strategy for
practical application of phosphorus-based anodes in sodium-ion batteries.
77
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Abstract (if available)
Abstract
Due to the increasing global demand for fossil fuels and the direct or indirect environmental consequences of their use, great attention is being devoted to alternative technologies for both energy generation and storage. Secondary battery systems have been considered as one of the energy storage strategy for the energy generated by renewable such as solar, wind and geothermal energy.
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Liu, Yihang
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Cathode and anode materials for sodium ion batteries
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Viterbi School of Engineering
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Doctor of Philosophy
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Electrical Engineering
Publication Date
10/16/2020
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